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MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY
The journal MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY
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is a scientific journal, devoted to original
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and technology. Topics of particular interest include metallic materials, inorganic materials, polymers, vacuum technique and
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VSEBINA – CONTENTS
PREGLEDNI ^LANEK – REVIEW ARTICLE
Additive manufacturing: the future of manufacturing
Dodajalna (3D) tehnologija: prihodnost proizvajanja
S. A. Adekanye, R. M. Mahamood, E. T. Akinlabi, M. G. Owolabi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 709
Pomembna obletnica revije Materiali in tehnologije: petdeset let izhajanja znanstvene periodi~ne publikacije
An important anniversary of the Materials and Technology journal: fifty years of publication
E. Nared . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 717
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Investigation of grain boundaries in Alloy 263 after special heat treatment
Preiskava mej zrn v zlitini 263 po posebni toplotni obdelavi
I. Slatkovský, M. Dománková, M. Sahul . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 721
Fracture toughness of ledeburitic Vanadis 6 steel after sub-zero treatment for 17 h and double tempering
Lomna `ilavost ledeburitnega jekla Vanadis 6 po toplotni obdelavi s 17-urnim podhlajevanjem in dvojnim popu{~anjem
J. Pta~inová, P. Jur~i, I. Dlouhý . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 729
Electronic and optical properties of the spinel oxides MgxZn1-xAl2O4 by first-principles calculations
Elektronske in opti~ne lastnosti spinelnih oksidov MgxZn1-xAl2O4, izpeljane iz teoreti~nih osnov
C. Xiang, J. X. Zhang, Y. Lu, D. Tian, C. Peng . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 735
Surface characteristics of Invar alloy according to micro-pulse electrochemical machining
Karakteristike povr{ine Invar zlitine glede na mikropulzno elektrokemi~no obdelavo
S.-H. Kim, S.-G. Choi, W.-K. Choi, E.-S. Lee . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 743
Durability of materials based on a polymer-silicate matrix and a lightweight aggregate exposed to aggressive influences
combined with high temperatures
Vzdr`ljivost materialov na osnovi iz polimer-silikatnih matric in lahkega dodatka, izpostavljenih agresivnim vplivom v kombinaciji z visokimi
temperaturami
T. Melichar, J. Byd`ovský, Á. Dufka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 751
Influence of thermomechanical treatment on the grain-growth behaviour of new Fe-Al based alloys with fine Al2O3
precipitates
Vpliv termomehanske obdelave FeAl zlitin s finimi Al2O3 izlo~ki na rast zrn
B. Ma{ek, O. Khalaj, H. Jirková, J. Svoboda, D. Bublíková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 759
Analysis of precipitates in aluminium alloys with the use of high-resolution electron microscopy and computer simulation
Raziskave oborin v aluminijevih zlitinah z visokoresolucijsko elektronsko mikroskopijo in ra~unalni{ko simulacijo
K. Matus, A. Tomiczek, K. Go³ombek, M. Pawlyta . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 769
Microstructural evaluation of Ni-SDC cermet from a representative 2D image and/or a 3D reconstruction based on a stack of
images
Vrednotenje mikrostruktur Ni-SDC kermeta z 2D in/ali 3D metodo
G. Kapun, M. Marin{ek, F. Merzel, S. [turm, M. Gaber{~ek, T. Skalar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 775
A facile method to prepare super-hydrophobic surfaces on silicone rubbers
Preprosta metoda za pripravo superhidrofobnih povr{in pri silikonskih gumah
H. Y. Jin, Y. F. Li, S. C. Nie, P. Z., N. K. Gao, W. Li. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 783
Investigation of the static icing property for super-hydrophobic coatings on aluminium
Preiskava lastnosti stati~ne zaledenitve pri superhidrofobnih prevlekah na aluminiju
H. Y. Jin, S. C. Nie, Y. F. Li, T. F. Xu, P. Zhang, W. Li . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 789
Effect of ball milling on the properties of the porous Ti–26Nb alloy for biomedical applications
Vpliv krogli~nega mletja na lastnosti porozne zlitine Ti–26Nb za biomedicinske aplikacije
G. Dercz, I. Matu³a . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 795
Effects of an addition of coir-pith particles on the mechanical properties and erosive-wear behavior of a
wood-dust-particle-reinforced phenol formaldehyde composite
Vplivi dodatka kokosovih vlaken fenol-formaldehidnemu kompozitu, oja~anem z lesnim prahom, na njegove mehanske lastnosti in erozijsko
obrabo
A. S. Jose, A. Athijayamani, K. Ramanathan, S. Sidhardhan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 805
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 51(5)707–879(2017)
MATER.
TEHNOL.
LETNIK
VOLUME 51
[TEV.
NO. 5
STR.
P. 707–879
LJUBLJANA
SLOVENIJA
SEP.–OKT.
SEP.–OCT. 2017
Optimum bushing length in thermal drilling of galvanized steel using artificial neural network coupled with genetic algorithm
Optimalna dol`ina podpore ({ablone, vodila) pri termi~nem vrtanju galvaniziranega jekla z uporabo umetne nevronske mre`e in
genetskega algoritma
N. Rajesh J. Hynes, R. Kumar, J. A. J. Sujana . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 813
Gelling polysaccharide as the electrolyte matrix in a dye-sensitized solar cell
@elirni polisaharid kot elektrolitna osnova v solarnih celicah, ob~utljivih na barvila
J. P. Bantang, D. Camacho . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 823
Development of a heat treatment for increasing the mechanical properties and stress corrosion resistance of 7000 Al alloys
Razvoj toplotne obdelave za izbolj{anje mehanskih lastnosti in napetostno korozijsko odpornost 7000 Al zlitin
M. Shakouri, M. Esmailian, S. Shabestari . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 831
Corrosion resistance of as-plated and heat-treated electroless dublex Ni-P/Ni-B-W coatings
Korozijska odpornost platiranih in neelektri~no topolotno obdelanih dupleks Ni-P/Ni-B-W prevlek
B. Yüksel, G. Erdogan, F. E. Bastan, R. A. Yýldý . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 837
Short-term creep of P91 heat-resistant steels at low stresses and an instantaneous-stress-change testing
Kratkotrajno lezenje toplotno odpornega jekla P91 pri nizkih napetostih in nenadni menjavi napetosti obremenjevanja
J. Zhe, S. Junjie, Z. Pengshuo . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 843
Effect of severe plastic and heavy cold deformation on the structural and mechanical properties of commercially pure
titanium
U~inek plasti~nosti in deformacije pri podhlajevanju na strukturne in mehanske lastnosti ~istega komercialnega titana
J. Palán, P. [utta, T. Kubina, M. Dománková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 849
Effect of yttrium and zirconium microalloying on the structure and properties of weld joints of a two-phase titanium alloy
U~inek mikrolegiranja itrija in cirkonija na strukturo in lastnosti na spoje zavrov dvofazne zlitine titana
A. Illarionov, A. Popov, S. Illarionova, D. Gadeev . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 855
Microstructure evolution and statistical analysis of Al/Cu friction-stir spot welds
Razvoj mikrostrukture in statisti~na analiza vrtilno-tornih to~kastih zvarov Al/Cu
M. P. Mubiayi, E. T. Akinlabi, M. E. Makhatha . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 861
Synthesis of PMMA/ZnO nanoparticles composite used for resin teeth
Sinteza PMMA/ZnO nanodelcev kompozitov za izdelavo zob iz umetnih smol
D. Popovi}, R. Bobovnik, S. Bolka, M. Vukadinovi}, V. Lazi}, R. Rudolf . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 871
ERRATUM . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 879
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
709–715
ADDITIVE MANUFACTURING: THE FUTURE OF
MANUFACTURING
DODAJALNA (3D) TEHNOLOGIJA: PRIHODNOST PROIZVAJANJA
Sheriff Adefemi Adekanye1, Rasheedat Modupe Mahamood1,2,
Esther Titilayo Akinlabi2, Moses Gbadebo Owolabi3
1University of Ilorin, Department of Mechanical Engineering, Nigeria
2University of Johannesburg, Department of Mechanical Engineering Science, South Africa
3Howard University, Department of Mechanical Engineering, Washington, USA
adekanyefm@gmail.com
Prejem rokopisa – received: 2016-08-26; sprejem za objavo – accepted for publication: 2017-03-16
doi:10.17222/mit.2016.261
Additive manufacturing is an advanced manufacturing method used to fabricate prototypes, tooling as well as functional
products. Additive manufacturing helps us produce complex parts as single-unit objects, which was not possible with the
traditional manufacturing methods. There are different types of additive-manufacturing technologies, including selective laser
melting, laser-metal-deposition process, fused deposition modeling and electron-beam melting. All these additive-manu-
facturing technologies produce three-dimensional (3D) objects by adding materials layer after layer. A 3D object is built directly
from a 3D computer-aided-design (CAD) model of the object. Additive manufacturing is a very promising method for the
aerospace industry, in particular because of its ability to reduce the buy-to-fly ratio. This technology is the technology of the
future because it is going to change the way products are designed and manufactured. In this research, various additive-manu-
facturing technologies are described in detail and some of the research works in this field are also presented. The future research
directions are also highlighted.
Keywords: additive manufacturing, fused deposition modelling, laser metal deposition, selective laser melting, selective laser
sintering
Dodajalni proizvodni proces je napredna tehnologija, ki se uporablja za izdelavo prototipov, orodij in funkcionalnih izdelkov. Z
naprednim proizvodnim procesom lahko izdelamo izdelke ali dele kompliciranih oblik, ki jih ni mo`no izdelati s tradicionalnimi
metodami proizvodnje. Obstajajo razli~ne vrste naprednih dodajalnih tehnologij, kot so na primer: selektivno lasersko
nataljevanje, proces laserskega nana{anja (depozicije) kovin, modeliranje z nana{anjem staljenih kapljic in taljenje z
elektronskim curkom. Vse te napredne dodajalne tehnologije omogo~ajo izdelavo tridimenzionalnih (3D) izdelkov z dodajanjem
materialov plast za plastjo. Tridimenzionalni produkt je izgrajen neposredno s pomo~jo 3D-ra~unalni{ko podprtega modeliranja
(angl. CAD). Napredna dodajalna proizvodnja je zelo obetavna proizvodna metoda, predvsem v letalski in vesoljski industriji,
{e posebej zaradi sposobnosti zmanj{evanja stro{kov izdelave (zmanj{anje razmerja med maso materiala, potrebnega za izdelavo
dolo~enega izdelka in njegovo dejansko maso).
Klju~ne besede: dodajalna proizvodnja, modeliranje z nana{anjem staljenih materialov, laserska depozicija kovin, selektivno
lasersko nataljevanje, selektivno lasersko sintranje
1 INTRODUCTION
The additive-manufacturing process is an advanced
manufacturing process that produces three-dimensional
(3D) objects directly from the 3D computer-aided-design
(CAD) digital information of the parts by adding ma-
terials layer by layer.1–3 This is different from the tradi-
tional manufacturing that involves material removal or an
energy-intensive process such as machining, casting and
forging. With additive manufacturing, a complex-shaped
product can be produced as a single object, which was
not possible in the past.
For the traditional manufacturing process, complex
parts have to be designed based on the ease of manufact-
uring the parts. Complex parts are usually broken down
into smaller parts because of the ease of manufacturing
those parts and the pieces are assembled at the later stage
through various joining processes. This practice does not
only involve labor-intensive processes but also results in
wastage of materials, in addition to the materials re-
moved during shaping and cutting operations; the excess
materials added for joining also contribute to material
wastage. The net weight of a product is also very large
because of the excess materials used during the
assembly. The additive-manufacturing technology, on the
other hand, can make a product directly from the 3D
CAD model of the product by just adding materials layer
by layer. Irrespective of the complexity, any part that can
be drawn using any computer-aided-design software can
be made and as a single-unit object. Additive-manu-
facturing technologies are classified into seven main
groups, namely: powder-bed fusion, directed energy
deposition, sheet lamination, material extrusion, binder
jetting, material jetting and vat photopolymerization.1 A
designer does not need to worry about the manufactur-
ability of a product; he/she is only concerned with the
functionality of the part being designed.
At the inception, the additive manufacturing process
was used to make prototypes because of the limitation in
Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715 709
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.183.183.4:621.039.337:316.422.44 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 51(5)709(2017)
Optimum bushing length in thermal drilling of galvanized steel using artificial neural network coupled with genetic algorithm
Optimalna dol`ina podpore ({ablone, vodila) pri termi~nem vrtanju galvaniziranega jekla z uporabo umetne nevronske mre`e in
genetskega algoritma
N. Rajesh J. Hynes, R. Kumar, J. A. J. Sujana . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 813
Gelling polysaccharide as the electrolyte matrix in a dye-sensitized solar cell
@elirni polisaharid kot elektrolitna osnova v solarnih celicah, ob~utljivih na barvila
J. P. Bantang, D. Camacho . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 823
Development of a heat treatment for increasing the mechanical properties and stress corrosion resistance of 7000 Al alloys
Razvoj toplotne obdelave za izbolj{anje mehanskih lastnosti in napetostno korozijsko odpornost 7000 Al zlitin
M. Shakouri, M. Esmailian, S. Shabestari . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 831
Corrosion resistance of as-plated and heat-treated electroless dublex Ni-P/Ni-B-W coatings
Korozijska odpornost platiranih in neelektri~no topolotno obdelanih dupleks Ni-P/Ni-B-W prevlek
B. Yüksel, G. Erdogan, F. E. Bastan, R. A. Yýldý . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 837
Short-term creep of P91 heat-resistant steels at low stresses and an instantaneous-stress-change testing
Kratkotrajno lezenje toplotno odpornega jekla P91 pri nizkih napetostih in nenadni menjavi napetosti obremenjevanja
J. Zhe, S. Junjie, Z. Pengshuo . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 843
Effect of severe plastic and heavy cold deformation on the structural and mechanical properties of commercially pure
titanium
U~inek plasti~nosti in deformacije pri podhlajevanju na strukturne in mehanske lastnosti ~istega komercialnega titana
J. Palán, P. [utta, T. Kubina, M. Dománková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 849
Effect of yttrium and zirconium microalloying on the structure and properties of weld joints of a two-phase titanium alloy
U~inek mikrolegiranja itrija in cirkonija na strukturo in lastnosti na spoje zavrov dvofazne zlitine titana
A. Illarionov, A. Popov, S. Illarionova, D. Gadeev . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 855
Microstructure evolution and statistical analysis of Al/Cu friction-stir spot welds
Razvoj mikrostrukture in statisti~na analiza vrtilno-tornih to~kastih zvarov Al/Cu
M. P. Mubiayi, E. T. Akinlabi, M. E. Makhatha . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 861
Synthesis of PMMA/ZnO nanoparticles composite used for resin teeth
Sinteza PMMA/ZnO nanodelcev kompozitov za izdelavo zob iz umetnih smol
D. Popovi}, R. Bobovnik, S. Bolka, M. Vukadinovi}, V. Lazi}, R. Rudolf . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 871
ERRATUM . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 879
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
709–715
ADDITIVE MANUFACTURING: THE FUTURE OF
MANUFACTURING
DODAJALNA (3D) TEHNOLOGIJA: PRIHODNOST PROIZVAJANJA
Sheriff Adefemi Adekanye1, Rasheedat Modupe Mahamood1,2,
Esther Titilayo Akinlabi2, Moses Gbadebo Owolabi3
1University of Ilorin, Department of Mechanical Engineering, Nigeria
2University of Johannesburg, Department of Mechanical Engineering Science, South Africa
3Howard University, Department of Mechanical Engineering, Washington, USA
adekanyefm@gmail.com
Prejem rokopisa – received: 2016-08-26; sprejem za objavo – accepted for publication: 2017-03-16
doi:10.17222/mit.2016.261
Additive manufacturing is an advanced manufacturing method used to fabricate prototypes, tooling as well as functional
products. Additive manufacturing helps us produce complex parts as single-unit objects, which was not possible with the
traditional manufacturing methods. There are different types of additive-manufacturing technologies, including selective laser
melting, laser-metal-deposition process, fused deposition modeling and electron-beam melting. All these additive-manu-
facturing technologies produce three-dimensional (3D) objects by adding materials layer after layer. A 3D object is built directly
from a 3D computer-aided-design (CAD) model of the object. Additive manufacturing is a very promising method for the
aerospace industry, in particular because of its ability to reduce the buy-to-fly ratio. This technology is the technology of the
future because it is going to change the way products are designed and manufactured. In this research, various additive-manu-
facturing technologies are described in detail and some of the research works in this field are also presented. The future research
directions are also highlighted.
Keywords: additive manufacturing, fused deposition modelling, laser metal deposition, selective laser melting, selective laser
sintering
Dodajalni proizvodni proces je napredna tehnologija, ki se uporablja za izdelavo prototipov, orodij in funkcionalnih izdelkov. Z
naprednim proizvodnim procesom lahko izdelamo izdelke ali dele kompliciranih oblik, ki jih ni mo`no izdelati s tradicionalnimi
metodami proizvodnje. Obstajajo razli~ne vrste naprednih dodajalnih tehnologij, kot so na primer: selektivno lasersko
nataljevanje, proces laserskega nana{anja (depozicije) kovin, modeliranje z nana{anjem staljenih kapljic in taljenje z
elektronskim curkom. Vse te napredne dodajalne tehnologije omogo~ajo izdelavo tridimenzionalnih (3D) izdelkov z dodajanjem
materialov plast za plastjo. Tridimenzionalni produkt je izgrajen neposredno s pomo~jo 3D-ra~unalni{ko podprtega modeliranja
(angl. CAD). Napredna dodajalna proizvodnja je zelo obetavna proizvodna metoda, predvsem v letalski in vesoljski industriji,
{e posebej zaradi sposobnosti zmanj{evanja stro{kov izdelave (zmanj{anje razmerja med maso materiala, potrebnega za izdelavo
dolo~enega izdelka in njegovo dejansko maso).
Klju~ne besede: dodajalna proizvodnja, modeliranje z nana{anjem staljenih materialov, laserska depozicija kovin, selektivno
lasersko nataljevanje, selektivno lasersko sintranje
1 INTRODUCTION
The additive-manufacturing process is an advanced
manufacturing process that produces three-dimensional
(3D) objects directly from the 3D computer-aided-design
(CAD) digital information of the parts by adding ma-
terials layer by layer.1–3 This is different from the tradi-
tional manufacturing that involves material removal or an
energy-intensive process such as machining, casting and
forging. With additive manufacturing, a complex-shaped
product can be produced as a single object, which was
not possible in the past.
For the traditional manufacturing process, complex
parts have to be designed based on the ease of manufact-
uring the parts. Complex parts are usually broken down
into smaller parts because of the ease of manufacturing
those parts and the pieces are assembled at the later stage
through various joining processes. This practice does not
only involve labor-intensive processes but also results in
wastage of materials, in addition to the materials re-
moved during shaping and cutting operations; the excess
materials added for joining also contribute to material
wastage. The net weight of a product is also very large
because of the excess materials used during the
assembly. The additive-manufacturing technology, on the
other hand, can make a product directly from the 3D
CAD model of the product by just adding materials layer
by layer. Irrespective of the complexity, any part that can
be drawn using any computer-aided-design software can
be made and as a single-unit object. Additive-manu-
facturing technologies are classified into seven main
groups, namely: powder-bed fusion, directed energy
deposition, sheet lamination, material extrusion, binder
jetting, material jetting and vat photopolymerization.1 A
designer does not need to worry about the manufactur-
ability of a product; he/she is only concerned with the
functionality of the part being designed.
At the inception, the additive manufacturing process
was used to make prototypes because of the limitation in
Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715 709
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.183.183.4:621.039.337:316.422.44 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 51(5)709(2017)
the choice of the materials that could be used on an
additive-manufacturing machine at that time.2–4 There are
a number of industries benefiting from this revolutionary
manufacturing technology,5 such as the aerospace, auto-
mobile and biomedical industries.6–9 In this study, some
of these additive-manufacturing technologies are re-
viewed with the aim of throwing light on these techno-
logies and describing their systems of operation as well
as their advantages, limitations and areas of applications.
The rest of the paper is organized as follows: section
2 of the paper presents the powder-bed-based process.
The extrusion-based process and sheet-lamination pro-
cess are presented in sections 3 and 4, respectively. Di-
rected-energy-deposition process is presented in section
5. Material development in additive manufacturing is
presented in section 6, while conclusions are presented
in section 7.
2 POWDER-BED-BASED PROCESS
This class of additive manufacturing technology uses
the energy from an electron beam or laser beam to fuse
or melt powder materials for the manufacturing of
parts.10,11 The additive-manufacturing technologies in
this class include direct metal laser sintering (DMLS),
electron-beam melting (EBM), selective heat sintering
(SHS), selective laser melting (SLM) and selective laser
sintering (SLS). Each of these technologies is described
in this section.
2.1 Direct metal laser sintering (DMLS)
A schematic diagram of direct laser sintering is
shown in Figure 1 below. This process manufactures
parts by compacting a metal powdered material with a
power source such as laser, without melting but by
binding the material together to create a solid structure
defined by the 3D CAD model of the part.12 The working
principle of this technology makes it capable of design-
ing and producing complex geometry of both internal
and external intricacies. In this process, to avoid the
collision of the recoater blade when moving, the building
and dispenser platforms are lowered by the layer
thickness. The powder metal needed to create a layer of
the material is supplied by the dispenser platform after it
has been ensured that the recoater stands in the right
position. The spreading of the metal powder from the
dispenser to the building platform is done by the re-
coater, moving from the right to the left position, and the
excess metal powder falls into the excess-powder
collector. The scan head moves the laser beam through
the two-dimensional (2D) cross-section, which is
switched on and off during the exposure of designated
areas. The powdered metals are generated, cured and
sintered by the solidified areas through the absorption of
energy.13 This process continues to create layer after
layer until the part is completed. Using this process,
parts have been successfully made from the materials
such as aluminium alloys – AlSi10Mg, stainless steel,
titanium alloys and cobalt chrome superalloy.13 Some of
the advantages of this additive-manufacturing process
are highlighted below.12,13
High speed: parts are produced easily within hours
since special tools are not required.
Complex geometry: this technology allows designs of
both internal and external features.
High quality: it creates parts with high accuracy and
detailed resolution.
It provides parts with better mechanical strength.
In spite of all these advantages, some of the limiting
factors are as follows: the process is expensive and
power intensive, the surface needs to be polished and the
metal support structure removed, and the thermal
post-processing is time consuming.14,15
2.2 Electron-beam melting
Figure 2 shows a schematic presentation of an
electron-beam process. Parts are developed by adding
material layer by layer, following the path described by
the 3D CAD data program.16 Electron beam is generated
within an electron-beam gun; the tungsten filament is
heated to an extremely high temperature so that it
releases electrons; the electrons are accelerated with an
electric field and then focused with electromagnetic
coils.17,18 The electron beam melts each layer of the
S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
710 Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Schematic diagram of the direct metal laser sintering
system14
Figure 2: Schematic drawing of an electron-beam additive-manu-
facturing machine18
metal powder to the desired geometry in a vacuum; the
process in the vacuum eliminates impurities and yields
high-strength properties of the material under
processing.16 The materials used for electron-beam
melting are mainly steel and titanium alloy. The
following advantages are characteristic of this
technology:17
• an increased efficiency of the raw-material use,
• the vacuum eliminates impurities and provides a
good thermal environment for a freeform fabrication,
• it significantly reduces the amount of finishing
operations,
• freedom in design – design for functionality,
• processing of high-melting-point and highly reactive
materials,
• decreased lead times for design and fabrication,
• a high degree of component customization.
Some of the disadvantages of the electron-beam
technology include the following:19
• It requires a vacuum, another system to the machine,
which must be maintained, leading to a high cost.
• It produces X-rays while in operation, which is
hazardous.
2.3 Selective laser sintering
Selective laser sintering is also a powder-bed-based
technology that involves the spreading of powder and the
subsequent compaction. A schematic diagram of the
selective laser sintering is shown in Figure 3. The set-up
is made up of a laser, an automatic powder-layering
apparatus, and a computer system for process control.20
Selective laser sintering uses a substrate for the part
fabrication, which is fixed onto the building platform and
leveled. The sealed building chamber containing oxygen
has a reduced amount of oxygen due to the protective
inert gas such as argon, which is fed into the chamber. A
thin layer of a loose powder is deposited on the substrate
by the layering mechanism. The laser beam scans the
powder-bed surface through the CAD program forming
layers of the material to be produced. The powder
spreading and laser-treatment process are repeated and
the parts are built layer by layer until completion.21–23
This technology allows a variety of materials, the most
common are: wax, paraffin, polymer-metal powders or
various types of steel alloys, polymers, nylon and carbo-
nates.24 The advantages of using selective laser sintering
include:
• parts can be created out of a wide selection of
materials,
• complex geometry is made easy as long as the
non-sintered powder can be removed easily,
• when printing overhanging, unsupported structures,
supports are not needed because the unused powder
provides the necessary support.
Disadvantages of selective laser sintering are:
• the equipment is very expensive,
• the most common problem of this technology is that
the fabricated parts are porous and the surface could
be rough, depending on the materials used,
• thermal distortion occurs on polymer parts and can
cause shrinking and warping of the fabricated parts.
2.4 Selective heat sintering
Figure 4 shows a schematic presentation of the
selective-heat-sintering process, which operates by selec-
tively fusing a thin layer of polymer powder through a
thermal print-head assembly. This assembly operates
bi-directionally and incorporates thermal print heads at
point A as shown in the diagram. The part labelled as B
is the powder-deposition mechanism, C is the layer
heater, D is the chamber for the material built up in an
internal build volume, E is the floor that is a vertically
movable building platform; at F, fresh powder is supplied
via scoops to the powder-deposition mechanism from the
powder containers.25 The heated building platform builds
up the parts, using the layer technology, with the roller
spreading across the plastic powders in layers. The
thermal print head of the apparatus forms a part in its full
cross-section with the already laid powder. The heat at
the building platform sinters the top layer of the powder;
once completed, the process is repeated until a complete
3D object is produced. The complex geometries pro-
S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715 711
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Schematic presentation of selective heat sintering25
Figure 3: Schematic presentation of a selective-laser-sintering
apparatus21
the choice of the materials that could be used on an
additive-manufacturing machine at that time.2–4 There are
a number of industries benefiting from this revolutionary
manufacturing technology,5 such as the aerospace, auto-
mobile and biomedical industries.6–9 In this study, some
of these additive-manufacturing technologies are re-
viewed with the aim of throwing light on these techno-
logies and describing their systems of operation as well
as their advantages, limitations and areas of applications.
The rest of the paper is organized as follows: section
2 of the paper presents the powder-bed-based process.
The extrusion-based process and sheet-lamination pro-
cess are presented in sections 3 and 4, respectively. Di-
rected-energy-deposition process is presented in section
5. Material development in additive manufacturing is
presented in section 6, while conclusions are presented
in section 7.
2 POWDER-BED-BASED PROCESS
This class of additive manufacturing technology uses
the energy from an electron beam or laser beam to fuse
or melt powder materials for the manufacturing of
parts.10,11 The additive-manufacturing technologies in
this class include direct metal laser sintering (DMLS),
electron-beam melting (EBM), selective heat sintering
(SHS), selective laser melting (SLM) and selective laser
sintering (SLS). Each of these technologies is described
in this section.
2.1 Direct metal laser sintering (DMLS)
A schematic diagram of direct laser sintering is
shown in Figure 1 below. This process manufactures
parts by compacting a metal powdered material with a
power source such as laser, without melting but by
binding the material together to create a solid structure
defined by the 3D CAD model of the part.12 The working
principle of this technology makes it capable of design-
ing and producing complex geometry of both internal
and external intricacies. In this process, to avoid the
collision of the recoater blade when moving, the building
and dispenser platforms are lowered by the layer
thickness. The powder metal needed to create a layer of
the material is supplied by the dispenser platform after it
has been ensured that the recoater stands in the right
position. The spreading of the metal powder from the
dispenser to the building platform is done by the re-
coater, moving from the right to the left position, and the
excess metal powder falls into the excess-powder
collector. The scan head moves the laser beam through
the two-dimensional (2D) cross-section, which is
switched on and off during the exposure of designated
areas. The powdered metals are generated, cured and
sintered by the solidified areas through the absorption of
energy.13 This process continues to create layer after
layer until the part is completed. Using this process,
parts have been successfully made from the materials
such as aluminium alloys – AlSi10Mg, stainless steel,
titanium alloys and cobalt chrome superalloy.13 Some of
the advantages of this additive-manufacturing process
are highlighted below.12,13
High speed: parts are produced easily within hours
since special tools are not required.
Complex geometry: this technology allows designs of
both internal and external features.
High quality: it creates parts with high accuracy and
detailed resolution.
It provides parts with better mechanical strength.
In spite of all these advantages, some of the limiting
factors are as follows: the process is expensive and
power intensive, the surface needs to be polished and the
metal support structure removed, and the thermal
post-processing is time consuming.14,15
2.2 Electron-beam melting
Figure 2 shows a schematic presentation of an
electron-beam process. Parts are developed by adding
material layer by layer, following the path described by
the 3D CAD data program.16 Electron beam is generated
within an electron-beam gun; the tungsten filament is
heated to an extremely high temperature so that it
releases electrons; the electrons are accelerated with an
electric field and then focused with electromagnetic
coils.17,18 The electron beam melts each layer of the
S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
710 Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Schematic diagram of the direct metal laser sintering
system14
Figure 2: Schematic drawing of an electron-beam additive-manu-
facturing machine18
metal powder to the desired geometry in a vacuum; the
process in the vacuum eliminates impurities and yields
high-strength properties of the material under
processing.16 The materials used for electron-beam
melting are mainly steel and titanium alloy. The
following advantages are characteristic of this
technology:17
• an increased efficiency of the raw-material use,
• the vacuum eliminates impurities and provides a
good thermal environment for a freeform fabrication,
• it significantly reduces the amount of finishing
operations,
• freedom in design – design for functionality,
• processing of high-melting-point and highly reactive
materials,
• decreased lead times for design and fabrication,
• a high degree of component customization.
Some of the disadvantages of the electron-beam
technology include the following:19
• It requires a vacuum, another system to the machine,
which must be maintained, leading to a high cost.
• It produces X-rays while in operation, which is
hazardous.
2.3 Selective laser sintering
Selective laser sintering is also a powder-bed-based
technology that involves the spreading of powder and the
subsequent compaction. A schematic diagram of the
selective laser sintering is shown in Figure 3. The set-up
is made up of a laser, an automatic powder-layering
apparatus, and a computer system for process control.20
Selective laser sintering uses a substrate for the part
fabrication, which is fixed onto the building platform and
leveled. The sealed building chamber containing oxygen
has a reduced amount of oxygen due to the protective
inert gas such as argon, which is fed into the chamber. A
thin layer of a loose powder is deposited on the substrate
by the layering mechanism. The laser beam scans the
powder-bed surface through the CAD program forming
layers of the material to be produced. The powder
spreading and laser-treatment process are repeated and
the parts are built layer by layer until completion.21–23
This technology allows a variety of materials, the most
common are: wax, paraffin, polymer-metal powders or
various types of steel alloys, polymers, nylon and carbo-
nates.24 The advantages of using selective laser sintering
include:
• parts can be created out of a wide selection of
materials,
• complex geometry is made easy as long as the
non-sintered powder can be removed easily,
• when printing overhanging, unsupported structures,
supports are not needed because the unused powder
provides the necessary support.
Disadvantages of selective laser sintering are:
• the equipment is very expensive,
• the most common problem of this technology is that
the fabricated parts are porous and the surface could
be rough, depending on the materials used,
• thermal distortion occurs on polymer parts and can
cause shrinking and warping of the fabricated parts.
2.4 Selective heat sintering
Figure 4 shows a schematic presentation of the
selective-heat-sintering process, which operates by selec-
tively fusing a thin layer of polymer powder through a
thermal print-head assembly. This assembly operates
bi-directionally and incorporates thermal print heads at
point A as shown in the diagram. The part labelled as B
is the powder-deposition mechanism, C is the layer
heater, D is the chamber for the material built up in an
internal build volume, E is the floor that is a vertically
movable building platform; at F, fresh powder is supplied
via scoops to the powder-deposition mechanism from the
powder containers.25 The heated building platform builds
up the parts, using the layer technology, with the roller
spreading across the plastic powders in layers. The
thermal print head of the apparatus forms a part in its full
cross-section with the already laid powder. The heat at
the building platform sinters the top layer of the powder;
once completed, the process is repeated until a complete
3D object is produced. The complex geometries pro-
S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715 711
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Schematic presentation of selective heat sintering25
Figure 3: Schematic presentation of a selective-laser-sintering
apparatus21
duced are supported by the excess material surrounding
the object.26
This technology uses thermoplastic powder as its
working material. The benefits of this technology are as
follows: it allows inexpensive manufacturing and a good
concept evaluation; the thermal print heads are less
expensive and the overall cost is affordable.
2.5 Selective laser melting
Selective laser melting is similar to the selective-
laser-sintering process, with the main difference being
between full melting and fusing in selective laser sinter-
ing. The process begins with designing a 3D CAD
model; the model is broken down into parts to form a
number of finite layers. For each layer, the laser scan
path is calculated, defining both the boundary contour
and the form of the fill sequence called the raster pattern.
Layers are formed by spreading powder on the build
platform and then the laser melts the powder with the
scanning laser beam; the melted powder solidifies to
form a 3D solid component.27 A schematic diagram of
the process is shown in Figure 5.28
Different materials used with this technology include
steel, titanium, aluminium, cobalt-chromium and nickel
alloys. The benefits of using the selective-laser-melting
technology are as follows:
• it minimizes material wastage and saves costs,
• improved production-development cycle,
• it allows for complex geometry to be made,
• ideal process for a low-volume production,
• improved buy-to-fly ratio,
• functionally graded parts can be produced,
• it allows for fully customized parts that suit indi-
viduals.
The disadvantages of selective laser melting are as
follows:
• it is an expensive and a very slow process,
• tolerances and surface finishes are limited because of
the sticking of the unused powder on the surface of
the produced part.
3 EXTRUSION-BASED SYSTEM
An extrusion-based system uses material extrusion
where the material is heated to the molten state and then
extruded through a nozzle to form parts layer by layer.
The most common technology is fused deposition
modeling (FDM). A schematic diagram of the process is
shown in Figure 6.
Fused deposition modeling produces components by
depositing extruded material in layers through a nozzle.
Fused deposition modeling uses a 3D CAD program to
prepare and build a model; the material is heated to the
molten state and forced through the nozzle under
pressure to build up a component.29 The advantages of
fused deposition modeling are as follows:
• the technology supports complex geometries and
cavities,
• the technology produces high-grade parts,
• it is simple to use.
The limitations of fused deposition modeling are:
• limitation on the materials that can be used,
• limited size,
• the cost of the actual machine.
This technology uses different materials for compo-
nent fabrication, including acrylonitrile butadiene
styrene (ABS), polylactic acid (PLA), polycarbonate
(PC), polyamide (PA), polystyrene (PS), lignin, rubber,
nylon and others.29
4 SHEET-LAMINATION PROCESS
This is an additive-manufacturing technology, which
mainly uses metal sheets or paper for the manufacturing
process. This technology is characterized by a low cost
of production, high strength of the model, possibilities to
S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
712 Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Presentation of an extrusion-based system29Figure 5: Schematic presentation of selective laser melting28
control the range of speed of the process to create the
outcome desired. Apart from the benefits achieved when
using this technology, there are also limitations: the pro-
perties of the outcome product depends on the material
used and the process of fusing the material together
needs improved techniques.30
There are two variants of this process, namely: ultra-
sonic additive manufacturing (UAM) and laminated
object manufacturing (LOM). Ultrasonic additive manu-
facturing (UAM) uses metal as its working material and
employs the layer-by-layer techniques, whereby the
layers are welded together ultrasonically. This process
uses metals like aluminum, copper, stainless steel and
titanium alloy. Laminated object manufacturing (LOM)
incorporates the layer-by-layer lamination technique,
using sheets of paper as the working material and CO2
laser for cutting a sheet that has been joined together by
an adhesive to form layers of CAD-model parts. The
excess material that is not part of the component being
built is sliced into cubes, using a cross-hutch cutting
operation.31
5 DIRECTED ENERGY DEPOSITION
Directed energy deposition is a type of additive
manufacturing that uses the energy from a laser or elec-
tron beam to create a melt pool on the surface of the
substrate, where coaxial powder or wire is deposited,
following the path described by a CAD profile. An
example of this process is laser metal deposition, also
known as laser engineered net shaping (LENS), a pro-
cess with high material-utilization efficiency and cost
efficiency.32–41 The laser, or electron beam, leaves a solid
track of the melted powder. A schematic diagram of this
process is shown in Figure 7. The advantages of directed
energy deposition include the ability to carry out a
high-value repair on a part that was cost prohibitive in
the past; it can also be used to manufacture parts with
functionally graded materials or functionally graded
coatings.39–41 The materials that can be used for this tech-
nology include metals, alloys, ceramics and composites.
Other additive-manufacturing processes are material
jetting, binder jetting and the vat photopolymerization
process. These types of additive manufacturing use non-
metallic based materials and their applications are
limited. For further reading about these and other
additive-manufacturing processes, the reader can consult
references.42–52
A new 3D printing technology was developed and it
will help us create metallic, free-standing 3D structures.
This technology was developed by the researchers at
Harvard University.53 The existing direct ink writing 3D
printing technology is combined with laser annealing to
ensure accurate and stable free-standing 3D structures.
This technology is very useful in the applications where
free-standing 3D microstructures are required, e.g., in
sensors, electronics and also in biomedical applications.
6 MATERIAL DEVELOPMENT FOR ADDITIVE
MANUFACTURING
There is constant demand for light materials with
improved properties, especially in the automobile and
aerospace industries, with the aim of reducing the global
warming. The advent of additive manufacturing is the
beginning of a new dawn in the manufacturing indus-
tries. A number of restrictions were placed on material
development in the past because of thermodynamic
limitations as well as the difficulties in processing some
materials in the bulk form. With additive-manufacturing
technologies, new materials are now being developed
without any restriction. A number of complex metallic
alloys are now being developed and used for the fabri-
cation of important light-weight structures with im-
pressive properties, using the additive-manufacturing
process.54 Complex metallic alloys are intermetallic
compounds with unique thermal and transport properties
that cannot be found in any metal system.55
Complex metallic alloys are needed in a number of
applications because of their excellent properties, but
they cannot be used with the conventional manufacturing
processes because of their brittleness that makes them
difficult to process. These materials have low coeffi-
cients of friction, good corrosion resistance, good wear-
resistance properties and can now be processed using
additive manufacturing. New composite materials with
improved properties are introduced and used in additive-
manufacturing processes, with the functional-part pro-
duction now being commercialized.54 The traditional
manufacturing process presents a lot of challenges when
processing these materials because of the segregation of
constituent materials due to the thermodynamic proper-
ties of individual materials, but additive-manufacturing
technologies can be used to process these materials
without such problems.
S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715 713
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7: Schematic diagram of directed energy deposition41
duced are supported by the excess material surrounding
the object.26
This technology uses thermoplastic powder as its
working material. The benefits of this technology are as
follows: it allows inexpensive manufacturing and a good
concept evaluation; the thermal print heads are less
expensive and the overall cost is affordable.
2.5 Selective laser melting
Selective laser melting is similar to the selective-
laser-sintering process, with the main difference being
between full melting and fusing in selective laser sinter-
ing. The process begins with designing a 3D CAD
model; the model is broken down into parts to form a
number of finite layers. For each layer, the laser scan
path is calculated, defining both the boundary contour
and the form of the fill sequence called the raster pattern.
Layers are formed by spreading powder on the build
platform and then the laser melts the powder with the
scanning laser beam; the melted powder solidifies to
form a 3D solid component.27 A schematic diagram of
the process is shown in Figure 5.28
Different materials used with this technology include
steel, titanium, aluminium, cobalt-chromium and nickel
alloys. The benefits of using the selective-laser-melting
technology are as follows:
• it minimizes material wastage and saves costs,
• improved production-development cycle,
• it allows for complex geometry to be made,
• ideal process for a low-volume production,
• improved buy-to-fly ratio,
• functionally graded parts can be produced,
• it allows for fully customized parts that suit indi-
viduals.
The disadvantages of selective laser melting are as
follows:
• it is an expensive and a very slow process,
• tolerances and surface finishes are limited because of
the sticking of the unused powder on the surface of
the produced part.
3 EXTRUSION-BASED SYSTEM
An extrusion-based system uses material extrusion
where the material is heated to the molten state and then
extruded through a nozzle to form parts layer by layer.
The most common technology is fused deposition
modeling (FDM). A schematic diagram of the process is
shown in Figure 6.
Fused deposition modeling produces components by
depositing extruded material in layers through a nozzle.
Fused deposition modeling uses a 3D CAD program to
prepare and build a model; the material is heated to the
molten state and forced through the nozzle under
pressure to build up a component.29 The advantages of
fused deposition modeling are as follows:
• the technology supports complex geometries and
cavities,
• the technology produces high-grade parts,
• it is simple to use.
The limitations of fused deposition modeling are:
• limitation on the materials that can be used,
• limited size,
• the cost of the actual machine.
This technology uses different materials for compo-
nent fabrication, including acrylonitrile butadiene
styrene (ABS), polylactic acid (PLA), polycarbonate
(PC), polyamide (PA), polystyrene (PS), lignin, rubber,
nylon and others.29
4 SHEET-LAMINATION PROCESS
This is an additive-manufacturing technology, which
mainly uses metal sheets or paper for the manufacturing
process. This technology is characterized by a low cost
of production, high strength of the model, possibilities to
S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
712 Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Presentation of an extrusion-based system29Figure 5: Schematic presentation of selective laser melting28
control the range of speed of the process to create the
outcome desired. Apart from the benefits achieved when
using this technology, there are also limitations: the pro-
perties of the outcome product depends on the material
used and the process of fusing the material together
needs improved techniques.30
There are two variants of this process, namely: ultra-
sonic additive manufacturing (UAM) and laminated
object manufacturing (LOM). Ultrasonic additive manu-
facturing (UAM) uses metal as its working material and
employs the layer-by-layer techniques, whereby the
layers are welded together ultrasonically. This process
uses metals like aluminum, copper, stainless steel and
titanium alloy. Laminated object manufacturing (LOM)
incorporates the layer-by-layer lamination technique,
using sheets of paper as the working material and CO2
laser for cutting a sheet that has been joined together by
an adhesive to form layers of CAD-model parts. The
excess material that is not part of the component being
built is sliced into cubes, using a cross-hutch cutting
operation.31
5 DIRECTED ENERGY DEPOSITION
Directed energy deposition is a type of additive
manufacturing that uses the energy from a laser or elec-
tron beam to create a melt pool on the surface of the
substrate, where coaxial powder or wire is deposited,
following the path described by a CAD profile. An
example of this process is laser metal deposition, also
known as laser engineered net shaping (LENS), a pro-
cess with high material-utilization efficiency and cost
efficiency.32–41 The laser, or electron beam, leaves a solid
track of the melted powder. A schematic diagram of this
process is shown in Figure 7. The advantages of directed
energy deposition include the ability to carry out a
high-value repair on a part that was cost prohibitive in
the past; it can also be used to manufacture parts with
functionally graded materials or functionally graded
coatings.39–41 The materials that can be used for this tech-
nology include metals, alloys, ceramics and composites.
Other additive-manufacturing processes are material
jetting, binder jetting and the vat photopolymerization
process. These types of additive manufacturing use non-
metallic based materials and their applications are
limited. For further reading about these and other
additive-manufacturing processes, the reader can consult
references.42–52
A new 3D printing technology was developed and it
will help us create metallic, free-standing 3D structures.
This technology was developed by the researchers at
Harvard University.53 The existing direct ink writing 3D
printing technology is combined with laser annealing to
ensure accurate and stable free-standing 3D structures.
This technology is very useful in the applications where
free-standing 3D microstructures are required, e.g., in
sensors, electronics and also in biomedical applications.
6 MATERIAL DEVELOPMENT FOR ADDITIVE
MANUFACTURING
There is constant demand for light materials with
improved properties, especially in the automobile and
aerospace industries, with the aim of reducing the global
warming. The advent of additive manufacturing is the
beginning of a new dawn in the manufacturing indus-
tries. A number of restrictions were placed on material
development in the past because of thermodynamic
limitations as well as the difficulties in processing some
materials in the bulk form. With additive-manufacturing
technologies, new materials are now being developed
without any restriction. A number of complex metallic
alloys are now being developed and used for the fabri-
cation of important light-weight structures with im-
pressive properties, using the additive-manufacturing
process.54 Complex metallic alloys are intermetallic
compounds with unique thermal and transport properties
that cannot be found in any metal system.55
Complex metallic alloys are needed in a number of
applications because of their excellent properties, but
they cannot be used with the conventional manufacturing
processes because of their brittleness that makes them
difficult to process. These materials have low coeffi-
cients of friction, good corrosion resistance, good wear-
resistance properties and can now be processed using
additive manufacturing. New composite materials with
improved properties are introduced and used in additive-
manufacturing processes, with the functional-part pro-
duction now being commercialized.54 The traditional
manufacturing process presents a lot of challenges when
processing these materials because of the segregation of
constituent materials due to the thermodynamic proper-
ties of individual materials, but additive-manufacturing
technologies can be used to process these materials
without such problems.
S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715 713
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7: Schematic diagram of directed energy deposition41
A number of new materials with improved properties
and performances are now being developed for the
additive-manufacturing technologies. These new mate-
rials include the high-temperature copper alloy, deve-
loped by NASA, with excellent properties such as
retention of high strength at elevated temperatures, used
in rocket engines. Other materials that present challenges
with the conventional manufacturing processes and are
now successfully processed using the additive-manu-
facturing technologies include titanium aluminide, used
in the fabrication of high-speed gas-turbine blades,
reactor-grade pure niobium and NiTi.
7 CONCLUSIONS
The concept of additive manufacturing has provided
solutions to most engineering problems, especially the
problems faced by product designers in the past when
products needed to be designed and redesigned because
of the limitations imposed by the available manu-
facturing process. Designers can now freely design parts
without having to worry about how they will be made.
This research paper presents a review on various
additive-manufacturing technologies. The aim of this
paper is to provide a clear description of this novel
manufacturing process – the additive-manufacturing
process. Each of the technologies is described in this
paper, including the materials suitable for use in each
technology, advantages, disadvantages. Application areas
and the new material development relating to the addi-
tive-manufacturing technologies are presented. In
summary, additive manufacturing is a reliable techno-
logy that can be used to manufacture prototypes, tooling
and functional parts, as it is less time consuming, allow-
ing an easy fabrication of a complex geometry, an easy
production of customized and personalized parts, with-
out any need for special tools, and low material wastage
that reduces the overall cost of the production. Im-
provements in the computing power and the reduction in
mass-storage costs paved the way for processing the
large amounts of data typical of modern 3D CAD models
within reasonable time frames.
A new 3D printing technology has recently been
developed at Harvard University. This new additive-
manufacturing technology combines the existing direct
ink writing 3D printing technology and the annealing
process. The innovation was born out of the necessity for
flexible and wearable electronic devices such as sensors
and in biomedical applications. The technology uses the
3D printing technology to create complex architectures,
while a printed structure is simultaneously annealed in
midair. This helps to increase the accuracy of the printed
material, also allowing the printed material to be created
in free air without the need for any support structure.
This technology has opened up possibilities of creating
microscopic metallic free-standing 3D structures without
an auxiliary support and in a single manufacturing run.
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MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
A number of new materials with improved properties
and performances are now being developed for the
additive-manufacturing technologies. These new mate-
rials include the high-temperature copper alloy, deve-
loped by NASA, with excellent properties such as
retention of high strength at elevated temperatures, used
in rocket engines. Other materials that present challenges
with the conventional manufacturing processes and are
now successfully processed using the additive-manu-
facturing technologies include titanium aluminide, used
in the fabrication of high-speed gas-turbine blades,
reactor-grade pure niobium and NiTi.
7 CONCLUSIONS
The concept of additive manufacturing has provided
solutions to most engineering problems, especially the
problems faced by product designers in the past when
products needed to be designed and redesigned because
of the limitations imposed by the available manu-
facturing process. Designers can now freely design parts
without having to worry about how they will be made.
This research paper presents a review on various
additive-manufacturing technologies. The aim of this
paper is to provide a clear description of this novel
manufacturing process – the additive-manufacturing
process. Each of the technologies is described in this
paper, including the materials suitable for use in each
technology, advantages, disadvantages. Application areas
and the new material development relating to the addi-
tive-manufacturing technologies are presented. In
summary, additive manufacturing is a reliable techno-
logy that can be used to manufacture prototypes, tooling
and functional parts, as it is less time consuming, allow-
ing an easy fabrication of a complex geometry, an easy
production of customized and personalized parts, with-
out any need for special tools, and low material wastage
that reduces the overall cost of the production. Im-
provements in the computing power and the reduction in
mass-storage costs paved the way for processing the
large amounts of data typical of modern 3D CAD models
within reasonable time frames.
A new 3D printing technology has recently been
developed at Harvard University. This new additive-
manufacturing technology combines the existing direct
ink writing 3D printing technology and the annealing
process. The innovation was born out of the necessity for
flexible and wearable electronic devices such as sensors
and in biomedical applications. The technology uses the
3D printing technology to create complex architectures,
while a printed structure is simultaneously annealed in
midair. This helps to increase the accuracy of the printed
material, also allowing the printed material to be created
in free air without the need for any support structure.
This technology has opened up possibilities of creating
microscopic metallic free-standing 3D structures without
an auxiliary support and in a single manufacturing run.
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facturing and rapid prototyping, CIRP Annals – Manufacturing
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S. A. ADEKANYE et al.: ADDITIVE MANUFACTURING: THE FUTURE OF MANUFACTURING
Materiali in tehnologije / Materials and technology 51 (2017) 5, 709–715 715
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
E. NARED: POMEMBNA OBLETNICA REVIJE MATERIALI IN TEHNOLOGIJE: PETDESET LET IZHAJANJA ...
717–720
POMEMBNA OBLETNICA REVIJE MATERIALI IN TEHNOLOGIJE:
PETDESET LET IZHAJANJA ZNANSTVENE PERIODI^NE
PUBLIKACIJE
AN IMPORTANT ANNIVERSARY OF THE MATERIALS AND
TECHNOLOGY JOURNAL: FIFTY YEARS OF PUBLICATION
Erika Nared
In{titut za kovinske materiale in tehnologije, Lepi pot 11, 1000 Ljubljana
erika.nared@imt.si
Prejem rokopisa – received: 2017-08-11; sprejem za objavo – accepted for publication: 2017-09-04
doi:10.17222/mit.2017.136
Revija Materiali in tehnologije (MIT), ki jo izdaja In{titut za kovinske materiale in tehnologije Ljubljana, v leto{njem letu
praznuje petdeseto obletnico neprekinjenega izhajanja v tiskani obliki. Tudi v elektronski obliki je revija MIT, kot jo skraj{ano
imenujemo, na voljo `e kar precej let in sicer od leta 1996. V ~lanku je predstavljen pregled izhajanja serijske publikacije
Materiali in tehnologije (ISSN 1580-2949) v zadnjih desetih letih; od leta 2007 pa vse do danes. Opisane so nekatere
spremembe v reviji v tem obdobju, njena dinamika izhajanja ter vizija njenega izhajanja v naslednjih letih.
Klju~ne besede: znanstvena periodi~na publikacija, zgodovinski pregled, obletnica izhajanja, uredni{ka politika
The Materials and Technology journal (MIT), published by the Institute of Metals and Technology Ljubljana, is celebrating 50
years of its printed version, with the online archive stretching back to 1996. This article provides an overwiev of the journal
(ISSN 1580-2949), covering the past 10 years. Some of the new changes to the journal are described, together with its vision for
the future.
Keywords: scientific periodical publication, historical overview, anniversary of publication, editorial policy
1 PREDSTAVITEV REVIJE IN NJEN RAZVOJ
Revija Materiali in tehnologije ima zelo dolgo zgo-
dovino izhajanja, saj izhaja `e od leta 1967. @elezarski
zbornik (ISSN 0372-8633), kakor se je imenovala pred
petdesetimi leti, ko je za~ela izhajati, je bil strokovno
glasilo Slovenskih `elezarn in Metalur{kega in{tituta
Ljubljana. ^etrtletnik je takrat tehni~no urejal Edo @agar
in sicer do leta 1980; nasledila sta ga Darko Brada{kja
do leta 1987 in Jana Jamar do leta 1991. Glavni in odgo-
vorni urednik @elezarskega zbornika pa je bil vse do
konca njegovega izhajanja pod tem imenom mag. Jo`a
Arh (Slika 1), torej kar 24 let (1967–1991). @elezarski
zbornik je objavljal prispevke s podro~ja kovinskih in
deloma nekovinskih materialov.1
Po dveh desetletjih je revija zaradi vsebinskih
sprememb dobila novo ime: Kovine zlitine tehnologije
(KZT) (ISSN 1318-0010). Preimenovala se je zaradi vse-
binske raz{iritve, saj je po novem vsebovala prispevke
tako s podro~ja kovinskih materialov kot tudi s podro~ja
anorganskih materialov, polimerov in materialov za
namen vakuumske tehnike. Skladno z menjavo naslova
je dobila tudi novo {tevilko ISSN.1 Glavni in odgovorni
urednik je bil do leta 1994 {e vedno mag. Jo`e Arh,
nasledil ga je mag. Ale{ Lagoja, ki je oblikoval ured-
ni{ko politiko revije v naslednjih 4 letih (Slika 1).
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Materiali in tehnologije / Materials and technology 51 (2017) 5, 717–720 717
UDK 050:669.1:6 (497) ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 51(5)717(2017)
Slika 1 (od leve proti desni): Jo`e Arh – glavni urednik revije od 1967 do 1991, Ale{ Lagoja – glavni urednik revije od 1991 do 1995, Franc
Vodopivec – glavni urednik revije od 1995 do 2011, Matja` Torkar – glavni urednik revije od 2012 do julija 2016 – Paul McGuiness, sedanji
glavni urednik revije
E. NARED: POMEMBNA OBLETNICA REVIJE MATERIALI IN TEHNOLOGIJE: PETDESET LET IZHAJANJA ...
717–720
POMEMBNA OBLETNICA REVIJE MATERIALI IN TEHNOLOGIJE:
PETDESET LET IZHAJANJA ZNANSTVENE PERIODI^NE
PUBLIKACIJE
AN IMPORTANT ANNIVERSARY OF THE MATERIALS AND
TECHNOLOGY JOURNAL: FIFTY YEARS OF PUBLICATION
Erika Nared
In{titut za kovinske materiale in tehnologije, Lepi pot 11, 1000 Ljubljana
erika.nared@imt.si
Prejem rokopisa – received: 2017-08-11; sprejem za objavo – accepted for publication: 2017-09-04
doi:10.17222/mit.2017.136
Revija Materiali in tehnologije (MIT), ki jo izdaja In{titut za kovinske materiale in tehnologije Ljubljana, v leto{njem letu
praznuje petdeseto obletnico neprekinjenega izhajanja v tiskani obliki. Tudi v elektronski obliki je revija MIT, kot jo skraj{ano
imenujemo, na voljo `e kar precej let in sicer od leta 1996. V ~lanku je predstavljen pregled izhajanja serijske publikacije
Materiali in tehnologije (ISSN 1580-2949) v zadnjih desetih letih; od leta 2007 pa vse do danes. Opisane so nekatere
spremembe v reviji v tem obdobju, njena dinamika izhajanja ter vizija njenega izhajanja v naslednjih letih.
Klju~ne besede: znanstvena periodi~na publikacija, zgodovinski pregled, obletnica izhajanja, uredni{ka politika
The Materials and Technology journal (MIT), published by the Institute of Metals and Technology Ljubljana, is celebrating 50
years of its printed version, with the online archive stretching back to 1996. This article provides an overwiev of the journal
(ISSN 1580-2949), covering the past 10 years. Some of the new changes to the journal are described, together with its vision for
the future.
Keywords: scientific periodical publication, historical overview, anniversary of publication, editorial policy
1 PREDSTAVITEV REVIJE IN NJEN RAZVOJ
Revija Materiali in tehnologije ima zelo dolgo zgo-
dovino izhajanja, saj izhaja `e od leta 1967. @elezarski
zbornik (ISSN 0372-8633), kakor se je imenovala pred
petdesetimi leti, ko je za~ela izhajati, je bil strokovno
glasilo Slovenskih `elezarn in Metalur{kega in{tituta
Ljubljana. ^etrtletnik je takrat tehni~no urejal Edo @agar
in sicer do leta 1980; nasledila sta ga Darko Brada{kja
do leta 1987 in Jana Jamar do leta 1991. Glavni in odgo-
vorni urednik @elezarskega zbornika pa je bil vse do
konca njegovega izhajanja pod tem imenom mag. Jo`a
Arh (Slika 1), torej kar 24 let (1967–1991). @elezarski
zbornik je objavljal prispevke s podro~ja kovinskih in
deloma nekovinskih materialov.1
Po dveh desetletjih je revija zaradi vsebinskih
sprememb dobila novo ime: Kovine zlitine tehnologije
(KZT) (ISSN 1318-0010). Preimenovala se je zaradi vse-
binske raz{iritve, saj je po novem vsebovala prispevke
tako s podro~ja kovinskih materialov kot tudi s podro~ja
anorganskih materialov, polimerov in materialov za
namen vakuumske tehnike. Skladno z menjavo naslova
je dobila tudi novo {tevilko ISSN.1 Glavni in odgovorni
urednik je bil do leta 1994 {e vedno mag. Jo`e Arh,
nasledil ga je mag. Ale{ Lagoja, ki je oblikoval ured-
ni{ko politiko revije v naslednjih 4 letih (Slika 1).
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Materiali in tehnologije / Materials and technology 51 (2017) 5, 717–720 717
UDK 050:669.1:6 (497) ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 51(5)717(2017)
Slika 1 (od leve proti desni): Jo`e Arh – glavni urednik revije od 1967 do 1991, Ale{ Lagoja – glavni urednik revije od 1991 do 1995, Franc
Vodopivec – glavni urednik revije od 1995 do 2011, Matja` Torkar – glavni urednik revije od 2012 do julija 2016 – Paul McGuiness, sedanji
glavni urednik revije
Do konca izhajanja revije z imenom Kovine zlitine
tehnologije (1999), je uredni{ko politiko vodil prof. dr.
Franc Vodopivec (Slika 1), ki je bil na mestu urednika
16 let in tako po letih urednikovanja na drugem mestu.
Tehni~no uredni{tvo revije Kovine zlitine tehnologije je
vodila Jana Jamar.
Izdajatelj revije je bil In{titut za kovinske materiale
in tehnologije Ljubljana skupaj s soizdajatelji: ACRONI
Jesenice, IMPOL Slovenska Bistrica, Kemijski in{titut
Ljubljana, Koncern Slovenske `elezarne, Metal Ravne,
Talum Kidri~evo, Fakulteta za strojni{tvo Ljubljana,
Institut Jo`ef Stefan in Slovensko dru{tvo za tribologijo.1
Revija z novim imenom Materiali in tehnologije
(ISSN 1580-2949) je nasledila obe svoji predhodnici na
prelomu tiso~letja in z letom 2000 za~ela izhajati kot
znanstvena serijska publikacija, ki objavlja izvirne in
pregledne znanstvene ~lanke ter strokovne ~lanke, ki
obravnavajo teoreti~na in prakti~na vpra{anja naravo-
slovnih ved in tehnologije na podro~jih kovinskih in
anorganskih materialov, polimerov, vakuumske tehnike,
kompozitnih in gradbenih materialov ter nanomate-
rialov.2
Njen namen je bil raz{iriti njeno interesno podro~je,
da ne bi delovala le na podro~ju kovin. Odgovorni
urednik revije je bil vse do leta 2011 prof. dr. Franc
Vodopivec. Do istega leta je tehni~no urednikovanje
vodila Jana Jamar. Z letom 2012 so se v uredni{kem
odboru zgodile spremembe predvsem na kadrovskem
podro~ju. Uredni{ko delo je vse do svoje upokojitve v
letu 2016 vodil dr. Matja` Torkar (Slika 1), tehni~na
urednica v tem ~asu je bila dr. Danijela A. Skobir
Balanti~.
Leta 2016 je mesto odgovornega in glavnega
urednika revije prevzel dr. Paul John McGuiness (Sli-
ka 1), mesto pomo~nika glavnega in odgovornega ured-
nika doc. dr. Matja` Godec, ~astna glavna urednika sta
prof. dr. Franc Vodopivec in dr. Matja` Torkar. Tehni~no
uredni{tvo vodi Erika Nared, ki na In{titutu za kovinske
materiale in tehnologije vodi tudi specialno knji`nico.
Souredniki revije Materiali in tehnologije so: Igor
Beli~, Jaka Burja, Aleksandra Kocijan, Djordje Man-
drino, Irena Paulin, Danijela A. Skobir Balanti~, Darja
Steiner Petrovi~, Bojan Podgornik iz In{tituta za kovin-
ske materiale in tehnologije (IMT), Bo{tjan Markoli iz
Naravoslovnotehni{ke fakultete (NTF), Sre~o [kapin in
Rok Zaplotnik iz In{tituta Jo`ef Stefan (IJS) ter Ema
@agar iz Kemijskega in{tituta (KI).3 Osve`en je tudi
uredni{ki odbor in zasedba ~lanov mednarodnih pri-
dru`enih ~lanov uredni{kega odbora ter izdajateljskega
sveta.
^lanki revije MIT so indeksirani v bazah podatkov,
kot so: Science Citation Index Expanded, Materials
Science Citation Index in Journal Citation Reports
(Science Edition). Po slednji, bazi podatkov JCR, ima
revija trenutno faktor vpliva 0.436. ^lanki, objavljeni v
reviji, so indeksirani tudi v mednarodnih in doma~ih
sekundarnih virih/bazah: DOAJ (Directory of Open
Access Journals), Google Scholar, SCOPUS, WoS,
COBIB in dLib.si (Digitalna knji`nica Slovenije).
Revija si, tako kot `e ve~ino ~asa svojega izhajanja,
{e vedno prizadeva za vi{jo citiranost, saj bi bila
posledi~no vsekakor bolj zanimiva, tako za slovenske
raziskovalce kot tudi tuje, in bi pritegnila uveljavljene
doma~e in tuje raziskovalce, da bi v njej objavljali svoje
znanstvene ~lanke.3
2 PREGLED IZHAJANJA REVIJE IN OPIS
SPREMEMB V OBDOBJU 2007–2017
Revijo Materiali in tehnologije danes izdaja In{titut
za kovinske materiale in tehnologije Ljubljana (IMT),
skupaj s soizdajatelji: IMPOL - Industrija aluminija Slo-
venska Bistrica, SIJ METAL Ravne in TALUM Kidri-
~evo. Izdajanje revije sofinancira Javna agencija za
raziskovalno dejavnost Republike Slovenije (ARRS),
prej Javna agencija za knjigo (JAK). In{titut za kovinske
materiale in tehnologije sodeluje na javnih razpisih, ki
jih razpi{e ARRS ter tako pridobi sofinanciranje
izdajanja revije.
Periodika izhajanja je 6 {tevilk letno. V vsaki {tevilki
je povpre~no objavljenih od 20 do 25 ~lankov in pri-
spevkov, ki so: ve~inoma izvirni znanstveni ~lanki,
pregledni ~lanki in strokovni ~lanki ter uvodne besede
urednika ob posebnih prilo`nostih. Tak{en obseg omo-
go~a korekten in kvaliteten potek procesa objave
~lankov, od oddaje ~lanka v uredni{tvo, do recenzije,
prevodov in kon~no objave.
V letu 2016 so z novim vodstvom pri{le tudi spre-
membe.4 S {tevilko 6/2016 smo po mnogih letih
brezpla~nega objavljanja ~lankov, za~eli objavo ~lankov
zara~unavati. Eden od razlogov je tudi kr~enje sredstev
za sofinanciranje pri ARRS. Za obi~ajne ~lanke je tako
cena objave 300 EUR, za tiste ~lanke, ki bodo pred-
stavljeni na letni Mednarodni konferenci o materialih in
tehnologijah, ki poteka vsako leto v Portoro`u, pa je
cena za objavo 150 EUR. Druga sprememba je spreje-
manje in objavljanje le izvirnih znanstvenih in pre-
glednih ~lankov.4 Vzrok za odlo~itev uredni{tva, da
uvede objavljanje le znanstvenih prispevkov, je ta, da gre
pri strokovnih ~lankih zgolj za strokovna poro~ila brez
znanstvenega pristopa, in kot tako ti bolj sodijo v
strokovna ali druga interna glasila.
V desetletnem obdobju (2007–2017) je {tevilo oddaje
~lankov nara{~alo in v zadnjih {tirih letih naraslo do
skoraj 400 letno (Slika 2), zato so se ~asovni roki za
objavo podalj{ali skoraj na obdobje enega leta ali ve~,
kar je po mnenju uredni{tva absolutno predolgo.
Zato smo, v pomo~ avtorjem, posodobili Navodila za
avtorje, ki so objavljena tako v tiskani reviji (na zadnjih
straneh vsake {tevilke) kot na spletni strani revije.5
Predstavili smo vzorec oz. predlogo, kako naj bo ~lanek
napisan in, da poleg oddaje ~lanka `elimo od avtorjev
prejeti tudi kontrolni seznam, s katerim avtor potrdi, da
je seznanjen z veljavno politiko uredni{tva. Oddaja
E. NARED: POMEMBNA OBLETNICA REVIJE MATERIALI IN TEHNOLOGIJE: PETDESET LET IZHAJANJA ...
718 Materiali in tehnologije / Materials and technology 51 (2017) 5, 717–720
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
~lankov {e vedno poteka preko e-po{te: mit@imt.si, si pa
uredni{tvo revije prizadeva, da bi v prihodnje v celoti
uveljavili spletno oddajanje ~lankov. Glede na `elje in
zahteve mnogih avtorjev, smo v zadnjem letu uspeli
skraj{ati obdobje ~akanja na objavo ~lankov in ~akalne
vrste iz enega leta na 2 do 3 mesece in tako stopiti
naproti tako avtorjem, kot tudi bralcem oz. ciljni publiki,
ki ima v kraj{em ~asu mo`nost slediti novim spoznanjem
in novostim na podro~ju materialov in tehnologij.
V obdobju od leta 2007 do 2016 je v reviji Materiali
in tehnologije iz{lo 1.071 ~lankov, kar je veliko ve~ kot v
desetletnem obdobju pred letom 2007. V zadnjih letih se
je {tevilo objavljenih ~lankov pove~evalo, saj smo v
uredni{tvo prejemali vedno ve~ ~lankov, katerih avtorji
so `eleli objavo. Nemara je bil najpogostej{i razlog
ravno brezpla~na objava ~lankov. Najve~ji dvig je bil leta
2008, ko je iz{la posebna {tevilka revije, v katero so bili
vklju~eni ~lanki iz druge mednarodne konference o to-
plotni obdelavi in povr{inski obdelavi (2nd International
Conference on the Heat Treatment and Surface Engi-
neering of Tools and Dies), ter ponoven porast med
letoma 2013 in 2014, ko so {e vedno velik dele` ~lankov
predstavljali prispevki udele`encev Mednarodne kon-
ference o materialih in tehnologijah v Portoro`u, katerih
~lanki so bili objavljeni {e v reviji, ~etudi so nekateri
vsebinsko predstavljali zgolj poro~ila o eksperimentih in
ne toliko rezultatov raznih raziskav (Slika 2). Namen
tovrstnih objav je bil med drugim dati prilo`nost tudi
mladim, {e ne uveljavljenim raziskovalcem, da na ta
na~in predstavijo svoja dela.
Interne evidence od leta 2011 dalje ka`ejo, koliko
~lankov je vsakoletno prispelo na naslov uredni{tva
revije v `elji za objavo. Ta {tevilka je dosegla vrh leta
2015. S spremembami v lanskem letu (2016), ko je ured-
ni{ki odbor dolo~il zara~unavanje objave, se je {tevilo
prispelih ~lankov, pri~akovano, postopoma za~elo ni`ati.
2.1 [tevilo objavljenih ~lankov
V prihodnje si v uredni{tvu revije Materiali in
tehnologije `elimo ve~ objav doma~ih raziskovalcev in
strokovnjakov ter tudi tistih iz o`jega evropskega
prostora. V obdobju med 2011 in 2016 so avtorji ~lankov
najpogosteje iz naslednjih dr`av (Slika 3).
V letu 2011 je najve~je {tevilo ~lankov avtorjev iz
Slovenije, ta {tevilka se je v zadnjih letih zelo zni`ala, v
zadnjem ~asu namre~ prevladujejo predvsem objave
avtorjev iz ^e{ke, Tur~ije in Poljske. Pove~uje se torej
{tevilo tujih avtorjev glede na doma~e.
2.2 Pregled revije po letih izhajanja v obdobju od 2007
do 2017
Temeljit pregled revije postopno po letih izhajanja,
od za~etka pa do leta 2007, je mo~ najti v `e objavljeni
literaturi1, zato se bomo na tem mestu dotaknili let, ki so
sledila kasneje, torej v obdobju od 2007 do danes
(Tabela 1). Statisti~na analiza ka`e, da {tevilo ~lankov
na splo{no zelo nara{~a (2007:43 in 2015:159), {tevilo
znanstvenih ~lankov je vi{je kot {tevilo strokovnih, ve~je
je tudi {tevilo ~lankov v angle{kem jeziku.
Z uvedbo sprememb v letu 2016 je uredni{tvo odlo-
~ilo tudi, da bo pri ~lankih {e vedno povzetek v sloven-
skem in angle{kem jeziku, ravno tako tudi naslov ter
klju~ne besede, ki ozna~ujejo vsebino ~lanka v obeh
jezikih Od 4. {tevilke letnika 51 pa podnapisi k slikam in
tabelam ne bodo ve~ v dveh jezikih, pa~ pa le enem – in
sicer v angle{kem jeziku.
Tabela 1: Pregled po letih izhajanja, 2007–2017
Letnik Leto [tevilo zvezkovin {tevilk
Znanstveni in
strokovni ~lanki
41 2007 6 (1–6) 43
42 2008 6 (1–6) 38
43 2009 6 (1–6) 46
44 2010 6 (1–6) 45
45 2011 6 (1–6) 76
46 2012 6 (1–6) 142
47 2013 6 (1–6) 138
48 2014 6 (1–6) 157
49 2015 6 (1–6) 159
50 2016 6 (1–6) 156
51 2017 6 (1–2) 72*
SKUPAJ: 1.072
*{tevilo ~lankov do {tevilke 4 (2017)
E. NARED: POMEMBNA OBLETNICA REVIJE MATERIALI IN TEHNOLOGIJE: PETDESET LET IZHAJANJA ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 717–720 719
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Slika 2: Grafikon s prikazom {tevila ~lankov prispelih za objavo po
letih (rde~a barva) in {tevilo dejansko objavljenih (modra barva)
Slika 3: Dr`ave avtorjev objavljenih ~lankov med leti 2011–2016
Do konca izhajanja revije z imenom Kovine zlitine
tehnologije (1999), je uredni{ko politiko vodil prof. dr.
Franc Vodopivec (Slika 1), ki je bil na mestu urednika
16 let in tako po letih urednikovanja na drugem mestu.
Tehni~no uredni{tvo revije Kovine zlitine tehnologije je
vodila Jana Jamar.
Izdajatelj revije je bil In{titut za kovinske materiale
in tehnologije Ljubljana skupaj s soizdajatelji: ACRONI
Jesenice, IMPOL Slovenska Bistrica, Kemijski in{titut
Ljubljana, Koncern Slovenske `elezarne, Metal Ravne,
Talum Kidri~evo, Fakulteta za strojni{tvo Ljubljana,
Institut Jo`ef Stefan in Slovensko dru{tvo za tribologijo.1
Revija z novim imenom Materiali in tehnologije
(ISSN 1580-2949) je nasledila obe svoji predhodnici na
prelomu tiso~letja in z letom 2000 za~ela izhajati kot
znanstvena serijska publikacija, ki objavlja izvirne in
pregledne znanstvene ~lanke ter strokovne ~lanke, ki
obravnavajo teoreti~na in prakti~na vpra{anja naravo-
slovnih ved in tehnologije na podro~jih kovinskih in
anorganskih materialov, polimerov, vakuumske tehnike,
kompozitnih in gradbenih materialov ter nanomate-
rialov.2
Njen namen je bil raz{iriti njeno interesno podro~je,
da ne bi delovala le na podro~ju kovin. Odgovorni
urednik revije je bil vse do leta 2011 prof. dr. Franc
Vodopivec. Do istega leta je tehni~no urednikovanje
vodila Jana Jamar. Z letom 2012 so se v uredni{kem
odboru zgodile spremembe predvsem na kadrovskem
podro~ju. Uredni{ko delo je vse do svoje upokojitve v
letu 2016 vodil dr. Matja` Torkar (Slika 1), tehni~na
urednica v tem ~asu je bila dr. Danijela A. Skobir
Balanti~.
Leta 2016 je mesto odgovornega in glavnega
urednika revije prevzel dr. Paul John McGuiness (Sli-
ka 1), mesto pomo~nika glavnega in odgovornega ured-
nika doc. dr. Matja` Godec, ~astna glavna urednika sta
prof. dr. Franc Vodopivec in dr. Matja` Torkar. Tehni~no
uredni{tvo vodi Erika Nared, ki na In{titutu za kovinske
materiale in tehnologije vodi tudi specialno knji`nico.
Souredniki revije Materiali in tehnologije so: Igor
Beli~, Jaka Burja, Aleksandra Kocijan, Djordje Man-
drino, Irena Paulin, Danijela A. Skobir Balanti~, Darja
Steiner Petrovi~, Bojan Podgornik iz In{tituta za kovin-
ske materiale in tehnologije (IMT), Bo{tjan Markoli iz
Naravoslovnotehni{ke fakultete (NTF), Sre~o [kapin in
Rok Zaplotnik iz In{tituta Jo`ef Stefan (IJS) ter Ema
@agar iz Kemijskega in{tituta (KI).3 Osve`en je tudi
uredni{ki odbor in zasedba ~lanov mednarodnih pri-
dru`enih ~lanov uredni{kega odbora ter izdajateljskega
sveta.
^lanki revije MIT so indeksirani v bazah podatkov,
kot so: Science Citation Index Expanded, Materials
Science Citation Index in Journal Citation Reports
(Science Edition). Po slednji, bazi podatkov JCR, ima
revija trenutno faktor vpliva 0.436. ^lanki, objavljeni v
reviji, so indeksirani tudi v mednarodnih in doma~ih
sekundarnih virih/bazah: DOAJ (Directory of Open
Access Journals), Google Scholar, SCOPUS, WoS,
COBIB in dLib.si (Digitalna knji`nica Slovenije).
Revija si, tako kot `e ve~ino ~asa svojega izhajanja,
{e vedno prizadeva za vi{jo citiranost, saj bi bila
posledi~no vsekakor bolj zanimiva, tako za slovenske
raziskovalce kot tudi tuje, in bi pritegnila uveljavljene
doma~e in tuje raziskovalce, da bi v njej objavljali svoje
znanstvene ~lanke.3
2 PREGLED IZHAJANJA REVIJE IN OPIS
SPREMEMB V OBDOBJU 2007–2017
Revijo Materiali in tehnologije danes izdaja In{titut
za kovinske materiale in tehnologije Ljubljana (IMT),
skupaj s soizdajatelji: IMPOL - Industrija aluminija Slo-
venska Bistrica, SIJ METAL Ravne in TALUM Kidri-
~evo. Izdajanje revije sofinancira Javna agencija za
raziskovalno dejavnost Republike Slovenije (ARRS),
prej Javna agencija za knjigo (JAK). In{titut za kovinske
materiale in tehnologije sodeluje na javnih razpisih, ki
jih razpi{e ARRS ter tako pridobi sofinanciranje
izdajanja revije.
Periodika izhajanja je 6 {tevilk letno. V vsaki {tevilki
je povpre~no objavljenih od 20 do 25 ~lankov in pri-
spevkov, ki so: ve~inoma izvirni znanstveni ~lanki,
pregledni ~lanki in strokovni ~lanki ter uvodne besede
urednika ob posebnih prilo`nostih. Tak{en obseg omo-
go~a korekten in kvaliteten potek procesa objave
~lankov, od oddaje ~lanka v uredni{tvo, do recenzije,
prevodov in kon~no objave.
V letu 2016 so z novim vodstvom pri{le tudi spre-
membe.4 S {tevilko 6/2016 smo po mnogih letih
brezpla~nega objavljanja ~lankov, za~eli objavo ~lankov
zara~unavati. Eden od razlogov je tudi kr~enje sredstev
za sofinanciranje pri ARRS. Za obi~ajne ~lanke je tako
cena objave 300 EUR, za tiste ~lanke, ki bodo pred-
stavljeni na letni Mednarodni konferenci o materialih in
tehnologijah, ki poteka vsako leto v Portoro`u, pa je
cena za objavo 150 EUR. Druga sprememba je spreje-
manje in objavljanje le izvirnih znanstvenih in pre-
glednih ~lankov.4 Vzrok za odlo~itev uredni{tva, da
uvede objavljanje le znanstvenih prispevkov, je ta, da gre
pri strokovnih ~lankih zgolj za strokovna poro~ila brez
znanstvenega pristopa, in kot tako ti bolj sodijo v
strokovna ali druga interna glasila.
V desetletnem obdobju (2007–2017) je {tevilo oddaje
~lankov nara{~alo in v zadnjih {tirih letih naraslo do
skoraj 400 letno (Slika 2), zato so se ~asovni roki za
objavo podalj{ali skoraj na obdobje enega leta ali ve~,
kar je po mnenju uredni{tva absolutno predolgo.
Zato smo, v pomo~ avtorjem, posodobili Navodila za
avtorje, ki so objavljena tako v tiskani reviji (na zadnjih
straneh vsake {tevilke) kot na spletni strani revije.5
Predstavili smo vzorec oz. predlogo, kako naj bo ~lanek
napisan in, da poleg oddaje ~lanka `elimo od avtorjev
prejeti tudi kontrolni seznam, s katerim avtor potrdi, da
je seznanjen z veljavno politiko uredni{tva. Oddaja
E. NARED: POMEMBNA OBLETNICA REVIJE MATERIALI IN TEHNOLOGIJE: PETDESET LET IZHAJANJA ...
718 Materiali in tehnologije / Materials and technology 51 (2017) 5, 717–720
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
~lankov {e vedno poteka preko e-po{te: mit@imt.si, si pa
uredni{tvo revije prizadeva, da bi v prihodnje v celoti
uveljavili spletno oddajanje ~lankov. Glede na `elje in
zahteve mnogih avtorjev, smo v zadnjem letu uspeli
skraj{ati obdobje ~akanja na objavo ~lankov in ~akalne
vrste iz enega leta na 2 do 3 mesece in tako stopiti
naproti tako avtorjem, kot tudi bralcem oz. ciljni publiki,
ki ima v kraj{em ~asu mo`nost slediti novim spoznanjem
in novostim na podro~ju materialov in tehnologij.
V obdobju od leta 2007 do 2016 je v reviji Materiali
in tehnologije iz{lo 1.071 ~lankov, kar je veliko ve~ kot v
desetletnem obdobju pred letom 2007. V zadnjih letih se
je {tevilo objavljenih ~lankov pove~evalo, saj smo v
uredni{tvo prejemali vedno ve~ ~lankov, katerih avtorji
so `eleli objavo. Nemara je bil najpogostej{i razlog
ravno brezpla~na objava ~lankov. Najve~ji dvig je bil leta
2008, ko je iz{la posebna {tevilka revije, v katero so bili
vklju~eni ~lanki iz druge mednarodne konference o to-
plotni obdelavi in povr{inski obdelavi (2nd International
Conference on the Heat Treatment and Surface Engi-
neering of Tools and Dies), ter ponoven porast med
letoma 2013 in 2014, ko so {e vedno velik dele` ~lankov
predstavljali prispevki udele`encev Mednarodne kon-
ference o materialih in tehnologijah v Portoro`u, katerih
~lanki so bili objavljeni {e v reviji, ~etudi so nekateri
vsebinsko predstavljali zgolj poro~ila o eksperimentih in
ne toliko rezultatov raznih raziskav (Slika 2). Namen
tovrstnih objav je bil med drugim dati prilo`nost tudi
mladim, {e ne uveljavljenim raziskovalcem, da na ta
na~in predstavijo svoja dela.
Interne evidence od leta 2011 dalje ka`ejo, koliko
~lankov je vsakoletno prispelo na naslov uredni{tva
revije v `elji za objavo. Ta {tevilka je dosegla vrh leta
2015. S spremembami v lanskem letu (2016), ko je ured-
ni{ki odbor dolo~il zara~unavanje objave, se je {tevilo
prispelih ~lankov, pri~akovano, postopoma za~elo ni`ati.
2.1 [tevilo objavljenih ~lankov
V prihodnje si v uredni{tvu revije Materiali in
tehnologije `elimo ve~ objav doma~ih raziskovalcev in
strokovnjakov ter tudi tistih iz o`jega evropskega
prostora. V obdobju med 2011 in 2016 so avtorji ~lankov
najpogosteje iz naslednjih dr`av (Slika 3).
V letu 2011 je najve~je {tevilo ~lankov avtorjev iz
Slovenije, ta {tevilka se je v zadnjih letih zelo zni`ala, v
zadnjem ~asu namre~ prevladujejo predvsem objave
avtorjev iz ^e{ke, Tur~ije in Poljske. Pove~uje se torej
{tevilo tujih avtorjev glede na doma~e.
2.2 Pregled revije po letih izhajanja v obdobju od 2007
do 2017
Temeljit pregled revije postopno po letih izhajanja,
od za~etka pa do leta 2007, je mo~ najti v `e objavljeni
literaturi1, zato se bomo na tem mestu dotaknili let, ki so
sledila kasneje, torej v obdobju od 2007 do danes
(Tabela 1). Statisti~na analiza ka`e, da {tevilo ~lankov
na splo{no zelo nara{~a (2007:43 in 2015:159), {tevilo
znanstvenih ~lankov je vi{je kot {tevilo strokovnih, ve~je
je tudi {tevilo ~lankov v angle{kem jeziku.
Z uvedbo sprememb v letu 2016 je uredni{tvo odlo-
~ilo tudi, da bo pri ~lankih {e vedno povzetek v sloven-
skem in angle{kem jeziku, ravno tako tudi naslov ter
klju~ne besede, ki ozna~ujejo vsebino ~lanka v obeh
jezikih Od 4. {tevilke letnika 51 pa podnapisi k slikam in
tabelam ne bodo ve~ v dveh jezikih, pa~ pa le enem – in
sicer v angle{kem jeziku.
Tabela 1: Pregled po letih izhajanja, 2007–2017
Letnik Leto [tevilo zvezkovin {tevilk
Znanstveni in
strokovni ~lanki
41 2007 6 (1–6) 43
42 2008 6 (1–6) 38
43 2009 6 (1–6) 46
44 2010 6 (1–6) 45
45 2011 6 (1–6) 76
46 2012 6 (1–6) 142
47 2013 6 (1–6) 138
48 2014 6 (1–6) 157
49 2015 6 (1–6) 159
50 2016 6 (1–6) 156
51 2017 6 (1–2) 72*
SKUPAJ: 1.072
*{tevilo ~lankov do {tevilke 4 (2017)
E. NARED: POMEMBNA OBLETNICA REVIJE MATERIALI IN TEHNOLOGIJE: PETDESET LET IZHAJANJA ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 717–720 719
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Slika 2: Grafikon s prikazom {tevila ~lankov prispelih za objavo po
letih (rde~a barva) in {tevilo dejansko objavljenih (modra barva)
Slika 3: Dr`ave avtorjev objavljenih ~lankov med leti 2011–2016
2.3 Pregled revije po tipologiji
[tevilo ~lankov, ki spadajo v skupino izvirnih znan-
stvenih ~lankov se je z leti vi{alo, najve~ jih je bilo v letu
2016, najmanj pa v letu 2007. [tevilo strokovnih ~lan-
kov, ki so z letom 2016, po odlo~itvi uredni{tva ukinjeni,
je obdr`alo konstantno {tevilko, od 30 do 40 letno.
Tabela 2: Pregled prispevkov v letih od 2007 do 2017 po tipologiji
Leto Tipologija
1.01
(izvirni znan-
stveni ~lanek)
1.02
(pregledni
znanstveni
~lanek)
1.04
(strokovni
~lanek)
1.20
(predgovor,
spremna
beseda)
2007 39 4 1 0
2008 30 6 8 0
2009 38 3 9 0
2010 43 8 7 0
2011 79 4 14 0
2012 81 0 30 3
2013 98 6 30 0
2014 118 2 39 0
2015 114 6 35 0
2016 125 2 30 2
2017
(do {t. 3)* 46 2 10 1
*podatki pridobljeni pred izidom 4. {tevilke
3 DOSTOPNOST REVIJE
Revija Materiali in tehnologije je `e v devetdesetih
letih prej{njega stoletja, ko se je imenovala {e Kovine
zlitine tehnologije, stopila naproti ideji, da bi bila bolj
odprta in la`je dostopna {ir{i javnosti. V skladu z `eljami
in dejanji uredni{tva in navsezadnje te`njami {ir{e
javnosti ter splo{nih trendov tistega ~asa, je v letu 1998
postala dostopna na svetovnem spletu v celotnem bese-
dilu in jo je bilo sprva mo~ najti na spletni strani
Centralne tehni{ke knji`nice (repozitorij CTK), kasneje
pa na spletni strani www.imt.si, kjer je revija dostopna {e
danes.
Revija (njeni predhodnici @elezarski zbornik in Ko-
vine zlitine tehnologije) je bila v Narodni in univerzitetni
knji`nici (NUK) vklju~ena v projekt digitalizacije in je
danes na voljo v digitalni obliki tako na spletnem portalu
pri Digitalni knji`nici Slovenije pri NUK, kot tudi na
spletni strani revije MIT. Ob petdeseti obletnici izhajanja
revije smo se v uredni{tvu odlo~ili, da na spletni strani
revije uredimo celoten arhiv vseh {tevilk revije, tudi {te-
vilk njenih predhodnic, ki so doslej manjkale: @elezarski
zbornik (ISSN 0372-8633) in Kovine zlitine tehnologije
(ISSN 1318-0010). Revijo Materiali in tehnologije je
mogo~e prelistati kadarkoli in od koderkoli (~e le imate v
bli`ini internet), in sicer od prve {tevilke @elezarskega
zbornika iz leta 1967 pa do vseh ~lankov revije MIT do
danes. Ve~ino {tevilk revije v tiskani obliki hranimo tudi
v knji`nici IMT. Skupno je bilo tako pripravljenih, in v
spletni arhiv dodanih, 123 datotek; 102 {tevilki @elezar-
skega zbornika in 21 {tevilk revije KZT.6
4 ZAKLJU^EK
Za~etki revije in objava ~lankov samo iz podro~ja
`elezarstva in metalurgije pred petdesetimi leti, ko je
industrija jekla in `elezarstva predstavljala najve~ji dele`
tovrstne dejavnosti v gospodarstvu takratne Jugoslavije
in kasneje dr`ave Slovenije, so bili med drugim ravno
tako podlaga za nadaljnje raziskovanje in uveljavljanje
metalurgije kot panoge, ki je do danes po~asi, a vztrajno
pre{la v skoraj vse pore industrije in je v dana{njem ~asu
razvoja tehnologije na izredno visoki ravni. Danes tako z
visoko tehnolo{kimi principi in najnovej{imi tehnolo-
gijami na podro~ju tako kovinskih, kot tudi drugih
materialov, dosegamo zavidljive rezultate. Napredni ma-
teriali, napredne proizvodne tehnologije, nanotehnologije
in biotehnologije so klju~na podro~ja, identificirana v
strategiji pametne specializacije, ki bodo omogo~ala
evropski in slovenski industriji ohraniti mednarodno
globalno konkuren~nost in izkoristiti nove trge. Z njimi
sta metalurgija in kemijska industrija nelo~ljivo povezani
in sta del opredeljenih klju~nih tehnologij. Pri tem razvoj
novih materialov in tehnologij pomeni vstop novih, do
sedaj neznanih mo`nosti, na tr`i{~e, kjer imajo ra-
ziskave, razvoj in inovacije zelo pomembno vlogo.7 Del
raziskav in preizkusov, je predstavljen tudi skozi prispev-
ke in ~lanke v reviji MIT in tudi v drugih medijih,8 ki s
tem {iri nova dognanja, spoznanja in znanja na {ir{i krog
raziskovalcev, znanstvenikov in drugo potencialno publi-
ko. Revija Materiali in tehnologije je danes sicer pri-
znana tudi izven meja Slovenije, vendar pa si v prihodnje
{e vedno `elimo, da bi v reviji objavljalo ve~je {tevilo
strokovnjakov in raziskovalcev iz Slovenije in evrop-
skega prostora.
5 LITERATURA
1 N. Jamar, J. Jamar, Zgodovina znanstvene serijske publikacije
Materiali in tehnologije/Materials and Technology = Historical over-
view of the scientific journal Materiali in tehnologije/Materials and
Technology, Mater. Tehnol., 41 (2007) 1, 13–19
2 http://mit.imt.si/Revija/index-slo.html, 28.7.2017
3 http://mit.imt.si/Revija/information-slo.html, 28.7.2017
4 P. McGuiness, Predgovor urednika=Editor’s preface, Mater. Tehnol.,
50 (2016) 5, str. 639–640
5 http://mit.imt.si/Revija/authors-slo.html, 28.07.2017
6 http://mit.imt.si/Revija/archive.html, 29.07.2017
7 https://www.gzs.si/Novice/ArticleId/58642/srip-matpro, 29.7.2017
8 M. Godec, Najve~ji hit je 3D tisk, Glas gospodarstva, Naj materiali,
pano`na {tevilka, april (2017), 30–31
E. NARED: POMEMBNA OBLETNICA REVIJE MATERIALI IN TEHNOLOGIJE: PETDESET LET IZHAJANJA ...
720 Materiali in tehnologije / Materials and technology 51 (2017) 5, 717–720
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
I. SLATKOVSKÝ et al.: INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263 AFTER SPECIAL HEAT TREATMENT
721–727
INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263
AFTER SPECIAL HEAT TREATMENT
PREISKAVA MEJ ZRN V ZLITINI 263 PO POSEBNI TOPLOTNI
OBDELAVI
Ivan Slatkovský, Mária Dománková, Martin Sahul
Slovak University of Technology Bratislava, Faculty of Materials Science and Technology, Institute of Materials Science,
Bottová 25, 917 24, Trnava, Slovak Republic
ivan.slatkovsky@stuba.sk
Prejem rokopisa – received: 2016-06-20; sprejem za objavo – accepted for publication: 2017-01-24
doi:10.17222/mit.2016.115
Alloy 263 is well known for its very good creep resistance and also for its weldability. These kinds of properties are appreciated
in the power-plant industry where Alloy 263 is used for shafts in a high-pressure circle. One of the possible ways to improve the
properties of superalloys, including Alloy 263, is through the effect of the grain-boundary serration (GBS) which, as research
indicates, is associated with the improvement of the creep resistance that can lead to an increased efficiency of coal power
plants. Grain-boundary serration was observed in different kinds of superalloys although the formation mechanism of serration
has not been clearly explained yet. Some researchers reported that the formation of serration is associated with the change in the
character of the precipitates at grain boundaries. This paper deals with an investigation of the grain boundaries in Alloy 263
using two different kinds of heat treatment. To form serrated grain boundaries in Material A (MA), slow controlled cooling from
the temperature of solution annealing to 800 °C was carried out. Standard heat treatment of Alloy 263 was performed on
material B (MB). Experimental techniques of scanning electron microscopy (SEM) and transmission electron microscopy
(TEM), including electron diffraction, were used to analyze the microstructure, determine the character of the grain boundaries
and identify the secondary particles at the grain boundaries.
Keywords: Alloy 263, grain-boundary serration, precipitates
Zlitina 263 je znana po zelo dobri odpornosti proti lezenju in tudi po dobri varivosti. Te vrste lastnosti so zelo cenjene v
termoelektrarnah, kjer se zlitina 263 uporablja za gredi v visokotla~nem delu turbin. Eden od mo`nih na~inov za izbolj{anje
lastnosti superzlitin, vklju~no z zlitino 263, je tvorba (nastanek) nazob~anih kristalnih mej (angl. GBS). Preiskava je pokazala,
da je ta fenomen povezan z izbolj{anjem odpornosti proti lezenju, kar lahko vodi k ve~ji u~inkovitosti termoelektrarn na
premog. U~inek nazob~anosti mej kristalnih zrn so opazili pri razli~nih vrstah superzlitin, ~eprav mehanizem tvorbe {e ni
popolnoma pojasnjen. Nekateri raziskovalci so ugotovili, da je tvorba nazob~anosti povezana s spremembo lastnosti izlo~kov na
mejah med zrni. Prispevek se ukvarja s preiskavo mej med kristalnimi zrni v zlitini 263 z dvema razli~nima vrstama toplotne
obdelave. Nastanek nazob~anih mej kristalnih zrn materiala A (MA), je bil povzro~en s po~asnim kontroliranim ohlajanjem iz
temperature raztopnega `arjenja, ki je bila 800 ° C. Standardna toplotna obdelava se je izvajala za zlitino 263 - material B (MB).
Za analizo mikrostrukture so uporabili vrsti~ni elektronski mikroskop (SEM) in presevno elektronsko mikroskopijo (TEM),
vklju~no z elektronsko difrakcijo. Na ta na~in so dolo~ili lastnosti mej kristalnih zrn in sekundarnih delcev na mejah zrn.
Klju~ne besede: zlitina 263, nazob~anost mej kristalnih zrn, izlo~ki
1 INTRODUCTION
It is a well-known fact that a reduction in the CO2
emissions produced by coal power plants is one of the
major goals for the countries all over the world. A
possible way to reduce the CO2 emissions is to improve
thermal efficiencies through super critical and ultra-
super critical technologies in power plants, where
efficiencies above 40 % could be reached. Hence, to
achieve this kind of efficiency, the materials like nickel-
based superalloys are considered to be used (Fig-
ure 1).1–3
To prolong the life time of parts when the tempera-
ture exceeds 700 °C, scientists are looking for new paths
of material processing. In the case of heat treatment, one
of the possible ways is a modification of grain boun-
daries (GBs) when the character of the boundaries
changes from straight to zigzag.
The phenomenon of grain-boundary serration (GBS)
was observed in the nickel-based superalloys and in
austenitic stainless steels. The formation of serrations on
grain boundaries (GBs) has not been fully described yet.
However, it is a known fact that in the case of nickel-
based superalloys, the GBS is closely related to the slow
controlled cooling from the temperature of the solution
annealing.
Early studies of serration in the nickel-based alloys
were focused on the interaction between the grain
boundary and ’ phase.4,5 Until now, researchers have
found that the serration is associated not only with the
presence of the ’ phase, but also with other precipitates
like M23C6 carbide, the  phase or the ’’ phase on the
GBs that were observed in different kinds of superalloys.
The formation of these secondary particles could be
highly related to the serrated grain boundaries and their
formation.6–13 Recently, contemporary authors have
Materiali in tehnologije / Materials and technology 51 (2017) 5, 721–727 721
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 669.018:620.1:620.193 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)721(2017)
2.3 Pregled revije po tipologiji
[tevilo ~lankov, ki spadajo v skupino izvirnih znan-
stvenih ~lankov se je z leti vi{alo, najve~ jih je bilo v letu
2016, najmanj pa v letu 2007. [tevilo strokovnih ~lan-
kov, ki so z letom 2016, po odlo~itvi uredni{tva ukinjeni,
je obdr`alo konstantno {tevilko, od 30 do 40 letno.
Tabela 2: Pregled prispevkov v letih od 2007 do 2017 po tipologiji
Leto Tipologija
1.01
(izvirni znan-
stveni ~lanek)
1.02
(pregledni
znanstveni
~lanek)
1.04
(strokovni
~lanek)
1.20
(predgovor,
spremna
beseda)
2007 39 4 1 0
2008 30 6 8 0
2009 38 3 9 0
2010 43 8 7 0
2011 79 4 14 0
2012 81 0 30 3
2013 98 6 30 0
2014 118 2 39 0
2015 114 6 35 0
2016 125 2 30 2
2017
(do {t. 3)* 46 2 10 1
*podatki pridobljeni pred izidom 4. {tevilke
3 DOSTOPNOST REVIJE
Revija Materiali in tehnologije je `e v devetdesetih
letih prej{njega stoletja, ko se je imenovala {e Kovine
zlitine tehnologije, stopila naproti ideji, da bi bila bolj
odprta in la`je dostopna {ir{i javnosti. V skladu z `eljami
in dejanji uredni{tva in navsezadnje te`njami {ir{e
javnosti ter splo{nih trendov tistega ~asa, je v letu 1998
postala dostopna na svetovnem spletu v celotnem bese-
dilu in jo je bilo sprva mo~ najti na spletni strani
Centralne tehni{ke knji`nice (repozitorij CTK), kasneje
pa na spletni strani www.imt.si, kjer je revija dostopna {e
danes.
Revija (njeni predhodnici @elezarski zbornik in Ko-
vine zlitine tehnologije) je bila v Narodni in univerzitetni
knji`nici (NUK) vklju~ena v projekt digitalizacije in je
danes na voljo v digitalni obliki tako na spletnem portalu
pri Digitalni knji`nici Slovenije pri NUK, kot tudi na
spletni strani revije MIT. Ob petdeseti obletnici izhajanja
revije smo se v uredni{tvu odlo~ili, da na spletni strani
revije uredimo celoten arhiv vseh {tevilk revije, tudi {te-
vilk njenih predhodnic, ki so doslej manjkale: @elezarski
zbornik (ISSN 0372-8633) in Kovine zlitine tehnologije
(ISSN 1318-0010). Revijo Materiali in tehnologije je
mogo~e prelistati kadarkoli in od koderkoli (~e le imate v
bli`ini internet), in sicer od prve {tevilke @elezarskega
zbornika iz leta 1967 pa do vseh ~lankov revije MIT do
danes. Ve~ino {tevilk revije v tiskani obliki hranimo tudi
v knji`nici IMT. Skupno je bilo tako pripravljenih, in v
spletni arhiv dodanih, 123 datotek; 102 {tevilki @elezar-
skega zbornika in 21 {tevilk revije KZT.6
4 ZAKLJU^EK
Za~etki revije in objava ~lankov samo iz podro~ja
`elezarstva in metalurgije pred petdesetimi leti, ko je
industrija jekla in `elezarstva predstavljala najve~ji dele`
tovrstne dejavnosti v gospodarstvu takratne Jugoslavije
in kasneje dr`ave Slovenije, so bili med drugim ravno
tako podlaga za nadaljnje raziskovanje in uveljavljanje
metalurgije kot panoge, ki je do danes po~asi, a vztrajno
pre{la v skoraj vse pore industrije in je v dana{njem ~asu
razvoja tehnologije na izredno visoki ravni. Danes tako z
visoko tehnolo{kimi principi in najnovej{imi tehnolo-
gijami na podro~ju tako kovinskih, kot tudi drugih
materialov, dosegamo zavidljive rezultate. Napredni ma-
teriali, napredne proizvodne tehnologije, nanotehnologije
in biotehnologije so klju~na podro~ja, identificirana v
strategiji pametne specializacije, ki bodo omogo~ala
evropski in slovenski industriji ohraniti mednarodno
globalno konkuren~nost in izkoristiti nove trge. Z njimi
sta metalurgija in kemijska industrija nelo~ljivo povezani
in sta del opredeljenih klju~nih tehnologij. Pri tem razvoj
novih materialov in tehnologij pomeni vstop novih, do
sedaj neznanih mo`nosti, na tr`i{~e, kjer imajo ra-
ziskave, razvoj in inovacije zelo pomembno vlogo.7 Del
raziskav in preizkusov, je predstavljen tudi skozi prispev-
ke in ~lanke v reviji MIT in tudi v drugih medijih,8 ki s
tem {iri nova dognanja, spoznanja in znanja na {ir{i krog
raziskovalcev, znanstvenikov in drugo potencialno publi-
ko. Revija Materiali in tehnologije je danes sicer pri-
znana tudi izven meja Slovenije, vendar pa si v prihodnje
{e vedno `elimo, da bi v reviji objavljalo ve~je {tevilo
strokovnjakov in raziskovalcev iz Slovenije in evrop-
skega prostora.
5 LITERATURA
1 N. Jamar, J. Jamar, Zgodovina znanstvene serijske publikacije
Materiali in tehnologije/Materials and Technology = Historical over-
view of the scientific journal Materiali in tehnologije/Materials and
Technology, Mater. Tehnol., 41 (2007) 1, 13–19
2 http://mit.imt.si/Revija/index-slo.html, 28.7.2017
3 http://mit.imt.si/Revija/information-slo.html, 28.7.2017
4 P. McGuiness, Predgovor urednika=Editor’s preface, Mater. Tehnol.,
50 (2016) 5, str. 639–640
5 http://mit.imt.si/Revija/authors-slo.html, 28.07.2017
6 http://mit.imt.si/Revija/archive.html, 29.07.2017
7 https://www.gzs.si/Novice/ArticleId/58642/srip-matpro, 29.7.2017
8 M. Godec, Najve~ji hit je 3D tisk, Glas gospodarstva, Naj materiali,
pano`na {tevilka, april (2017), 30–31
E. NARED: POMEMBNA OBLETNICA REVIJE MATERIALI IN TEHNOLOGIJE: PETDESET LET IZHAJANJA ...
720 Materiali in tehnologije / Materials and technology 51 (2017) 5, 717–720
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
I. SLATKOVSKÝ et al.: INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263 AFTER SPECIAL HEAT TREATMENT
721–727
INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263
AFTER SPECIAL HEAT TREATMENT
PREISKAVA MEJ ZRN V ZLITINI 263 PO POSEBNI TOPLOTNI
OBDELAVI
Ivan Slatkovský, Mária Dománková, Martin Sahul
Slovak University of Technology Bratislava, Faculty of Materials Science and Technology, Institute of Materials Science,
Bottová 25, 917 24, Trnava, Slovak Republic
ivan.slatkovsky@stuba.sk
Prejem rokopisa – received: 2016-06-20; sprejem za objavo – accepted for publication: 2017-01-24
doi:10.17222/mit.2016.115
Alloy 263 is well known for its very good creep resistance and also for its weldability. These kinds of properties are appreciated
in the power-plant industry where Alloy 263 is used for shafts in a high-pressure circle. One of the possible ways to improve the
properties of superalloys, including Alloy 263, is through the effect of the grain-boundary serration (GBS) which, as research
indicates, is associated with the improvement of the creep resistance that can lead to an increased efficiency of coal power
plants. Grain-boundary serration was observed in different kinds of superalloys although the formation mechanism of serration
has not been clearly explained yet. Some researchers reported that the formation of serration is associated with the change in the
character of the precipitates at grain boundaries. This paper deals with an investigation of the grain boundaries in Alloy 263
using two different kinds of heat treatment. To form serrated grain boundaries in Material A (MA), slow controlled cooling from
the temperature of solution annealing to 800 °C was carried out. Standard heat treatment of Alloy 263 was performed on
material B (MB). Experimental techniques of scanning electron microscopy (SEM) and transmission electron microscopy
(TEM), including electron diffraction, were used to analyze the microstructure, determine the character of the grain boundaries
and identify the secondary particles at the grain boundaries.
Keywords: Alloy 263, grain-boundary serration, precipitates
Zlitina 263 je znana po zelo dobri odpornosti proti lezenju in tudi po dobri varivosti. Te vrste lastnosti so zelo cenjene v
termoelektrarnah, kjer se zlitina 263 uporablja za gredi v visokotla~nem delu turbin. Eden od mo`nih na~inov za izbolj{anje
lastnosti superzlitin, vklju~no z zlitino 263, je tvorba (nastanek) nazob~anih kristalnih mej (angl. GBS). Preiskava je pokazala,
da je ta fenomen povezan z izbolj{anjem odpornosti proti lezenju, kar lahko vodi k ve~ji u~inkovitosti termoelektrarn na
premog. U~inek nazob~anosti mej kristalnih zrn so opazili pri razli~nih vrstah superzlitin, ~eprav mehanizem tvorbe {e ni
popolnoma pojasnjen. Nekateri raziskovalci so ugotovili, da je tvorba nazob~anosti povezana s spremembo lastnosti izlo~kov na
mejah med zrni. Prispevek se ukvarja s preiskavo mej med kristalnimi zrni v zlitini 263 z dvema razli~nima vrstama toplotne
obdelave. Nastanek nazob~anih mej kristalnih zrn materiala A (MA), je bil povzro~en s po~asnim kontroliranim ohlajanjem iz
temperature raztopnega `arjenja, ki je bila 800 ° C. Standardna toplotna obdelava se je izvajala za zlitino 263 - material B (MB).
Za analizo mikrostrukture so uporabili vrsti~ni elektronski mikroskop (SEM) in presevno elektronsko mikroskopijo (TEM),
vklju~no z elektronsko difrakcijo. Na ta na~in so dolo~ili lastnosti mej kristalnih zrn in sekundarnih delcev na mejah zrn.
Klju~ne besede: zlitina 263, nazob~anost mej kristalnih zrn, izlo~ki
1 INTRODUCTION
It is a well-known fact that a reduction in the CO2
emissions produced by coal power plants is one of the
major goals for the countries all over the world. A
possible way to reduce the CO2 emissions is to improve
thermal efficiencies through super critical and ultra-
super critical technologies in power plants, where
efficiencies above 40 % could be reached. Hence, to
achieve this kind of efficiency, the materials like nickel-
based superalloys are considered to be used (Fig-
ure 1).1–3
To prolong the life time of parts when the tempera-
ture exceeds 700 °C, scientists are looking for new paths
of material processing. In the case of heat treatment, one
of the possible ways is a modification of grain boun-
daries (GBs) when the character of the boundaries
changes from straight to zigzag.
The phenomenon of grain-boundary serration (GBS)
was observed in the nickel-based superalloys and in
austenitic stainless steels. The formation of serrations on
grain boundaries (GBs) has not been fully described yet.
However, it is a known fact that in the case of nickel-
based superalloys, the GBS is closely related to the slow
controlled cooling from the temperature of the solution
annealing.
Early studies of serration in the nickel-based alloys
were focused on the interaction between the grain
boundary and ’ phase.4,5 Until now, researchers have
found that the serration is associated not only with the
presence of the ’ phase, but also with other precipitates
like M23C6 carbide, the  phase or the ’’ phase on the
GBs that were observed in different kinds of superalloys.
The formation of these secondary particles could be
highly related to the serrated grain boundaries and their
formation.6–13 Recently, contemporary authors have
Materiali in tehnologije / Materials and technology 51 (2017) 5, 721–727 721
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 669.018:620.1:620.193 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)721(2017)
found that in some alloys, in the early stages of the
formation, GBS occurred in the absence of the adjacent
coarse ’ particle or M23C6 carbide.12–14 Furthermore, H.
U. Hong et al.15 observed in their research that the
serrations on GBs were formed in the material after a
long solution annealing (2000 min) without any slow
controlled cooling. Also, no secondary particles were
observed.
As reported by the authors of references,10–13 GBS
leads to a morphological change of the carbides on GBs,
from granular to planar. A growth of secondary particles
on the serrated grain boundaries was observed mostly at
a low angle and on special GBs. The authors of referen-
ces10,11 reported that the formation of planar M23C6
carbides is orientated in the {111} plane. Also, the
orientation relationship between the carbides and the
matrix was observed.
Therefore, the purpose of this study is to investigate
the effect of special heat treatment on GBs and identify
the secondary phases formed at the GBs in Alloy 263.
2 EXPERIMENTAL PART
In order to investigate GBS, commercially available
hot-rolled nickel-based superalloy 263 was used in the
study. The chemical composition of the examined alloy
is given in Table 1.
A two-stage special heat-treatment method was
designed to form a GBS in material A, based on the
previous studies.6–13 The solution heat treatment followed
by slow cooling at a carefully controlled rate (until the
aging temperature was reached) was found to be
necessary to generate serrated GB structures.
Material A was solution annealed at 1150 °C for
80 min and cooled slowly to 800 °C at a cooling rate of
3 °C/min, followed by water quenching. After that,
precipitation hardening at 800 °C for 4 h was performed
and followed by air cooling.
On material B, the standard two-stage heat treatment
of Alloy 263 (solution annealing at 1150 °C for 80 min
followed by water quenching and precipitation hardening
at 800 °C for 4 h, cooled in the air) was applied.
For a SEM analysis, the samples were sliced, mecha-
nically grinded and polished. The etching solution for
the SEM observation was a solution containing 5 g
FeCl3, 15 mL HCl, 2 mL HNO3 and 60 mL ethanol used
for 10–15 s in order to reveal the GB configuration and
carbides. A JEOL 7600 F scanning electron microscope
was used for the observation of the GBs.
Thin foils were prepared for a detailed grain-boun-
dary observation using mechanical grinding to a thick-
ness of about 0.1 mm, and then electrolytically etched.
Etching was done on TENUPOL 5 in a solution of
perchloric acid and methanol (1:9). The temperature
during etching was -30 °C and the voltage was 25 V. For
the detailed observation, a transmission electron micro-
scope, JEOL 200 CX with an accelerating voltage of 200
kV, was used. To identify secondary particles at the GBs,
an electron-diffraction analysis was applied.
Extraction replicas were prepared in the following
route: samples were first prepared using standard
metallographic techniques. After that, they were etched
with the above-mentioned solution for 15 s. A thin
carbon film was sputtered onto the etched surface of the
samples. The carbon film (replica) was subsequently
electrolytically extracted, using 8 % perchloric acid, at a
bias of 10 V.
3 RESULTS
Figure 2 shows the microstructure of material A after
the special heat treatment. Light microscopy revealed a
polyhedral grain, the presence of twins as well as the
grain-size heterogeneity in material A. Two types of
grain-boundary morphology were observed with light
microscopy. Straight boundaries with a percentage share
of approximately 38 % of the total of 254 observed grain
boundaries were located in the structure of material A.
Serrated boundaries, as the second type, covered
approximately 39 % of the noticed boundaries. In some
cases, the type of boundary could not be determined
owing to the state of the boundary.
As a reference sample, material B (Figure 3) was
processed with the conventional heat treatment. In
material B, we also observed the base microstructure
I. SLATKOVSKÝ et al.: INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263 AFTER SPECIAL HEAT TREATMENT
722 Materiali in tehnologije / Materials and technology 51 (2017) 5, 721–727
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Table 1: Nominal chemical composition of Alloy 263
Chemical-composition limits
Weight % Ni Cr Co Mo C Al Ti Al+Ti Mn
263 Bal 19–21 19–21 5.6–6.1 0.04–0.08 0.60 max 1.90–2.40 2.40–2.80 0.60 max
Weight % Si Fe Cu B Pb S Ag Bi
263 0.40 max 0.07 max 0.020 max 0.005 max 0.0020 0.007 max 0.005 max 0.0001 max
Figure 1: Net efficiency development in the case of hard coal-fired
power plants considering different structural alloys3
with polyhedral grains with the heterogeneity of the
grain size and twins similar to those in material A. As
expected, no serrated boundaries were observed in
material B after the conventional heat treatment.
The grain-boundary details taken with SEM indicate
the presence of precipitates on the serrated and straight
grain boundaries in material A as well as in material B
(Figures 4 and 5). Results of the EDX analysis taken
from material A (Figure 6) reveal the eventual presence
of carbon-rich secondary particles, which could possibly
indicate the presence of carbides at the grain boundaries.
The occurrence of the other elements at the grain boun-
daries was not significant. Comparable results were also
noticed for material B. To confirm the presence of the
secondary particles formed at the boundaries and to
identify the chemical nature of these particles, ED and
EDX using the TEM were taken for materials A and B,
and the results are shown below.
The authors of reference16 predicted and identified
typical kinds of the secondary particles for this alloy,
allowing us to expect mainly the presence of the MC,
M6C and M23C6 carbides.
Figures 7 to 9 summarize the identified phases on the
replicas for material A. Electron diffraction spectra as
well as the EDX analysis of the extracted precipitates
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Figure 5: SEM, material B: a) straight grain-boundary triple point and
b) detail of precipitatesFigure 3: Material B
Figure 4: SEM, material A: a) serrated grain-boundary triple point
and b) detail of precipitates
Figure 2: Material A
found that in some alloys, in the early stages of the
formation, GBS occurred in the absence of the adjacent
coarse ’ particle or M23C6 carbide.12–14 Furthermore, H.
U. Hong et al.15 observed in their research that the
serrations on GBs were formed in the material after a
long solution annealing (2000 min) without any slow
controlled cooling. Also, no secondary particles were
observed.
As reported by the authors of references,10–13 GBS
leads to a morphological change of the carbides on GBs,
from granular to planar. A growth of secondary particles
on the serrated grain boundaries was observed mostly at
a low angle and on special GBs. The authors of referen-
ces10,11 reported that the formation of planar M23C6
carbides is orientated in the {111} plane. Also, the
orientation relationship between the carbides and the
matrix was observed.
Therefore, the purpose of this study is to investigate
the effect of special heat treatment on GBs and identify
the secondary phases formed at the GBs in Alloy 263.
2 EXPERIMENTAL PART
In order to investigate GBS, commercially available
hot-rolled nickel-based superalloy 263 was used in the
study. The chemical composition of the examined alloy
is given in Table 1.
A two-stage special heat-treatment method was
designed to form a GBS in material A, based on the
previous studies.6–13 The solution heat treatment followed
by slow cooling at a carefully controlled rate (until the
aging temperature was reached) was found to be
necessary to generate serrated GB structures.
Material A was solution annealed at 1150 °C for
80 min and cooled slowly to 800 °C at a cooling rate of
3 °C/min, followed by water quenching. After that,
precipitation hardening at 800 °C for 4 h was performed
and followed by air cooling.
On material B, the standard two-stage heat treatment
of Alloy 263 (solution annealing at 1150 °C for 80 min
followed by water quenching and precipitation hardening
at 800 °C for 4 h, cooled in the air) was applied.
For a SEM analysis, the samples were sliced, mecha-
nically grinded and polished. The etching solution for
the SEM observation was a solution containing 5 g
FeCl3, 15 mL HCl, 2 mL HNO3 and 60 mL ethanol used
for 10–15 s in order to reveal the GB configuration and
carbides. A JEOL 7600 F scanning electron microscope
was used for the observation of the GBs.
Thin foils were prepared for a detailed grain-boun-
dary observation using mechanical grinding to a thick-
ness of about 0.1 mm, and then electrolytically etched.
Etching was done on TENUPOL 5 in a solution of
perchloric acid and methanol (1:9). The temperature
during etching was -30 °C and the voltage was 25 V. For
the detailed observation, a transmission electron micro-
scope, JEOL 200 CX with an accelerating voltage of 200
kV, was used. To identify secondary particles at the GBs,
an electron-diffraction analysis was applied.
Extraction replicas were prepared in the following
route: samples were first prepared using standard
metallographic techniques. After that, they were etched
with the above-mentioned solution for 15 s. A thin
carbon film was sputtered onto the etched surface of the
samples. The carbon film (replica) was subsequently
electrolytically extracted, using 8 % perchloric acid, at a
bias of 10 V.
3 RESULTS
Figure 2 shows the microstructure of material A after
the special heat treatment. Light microscopy revealed a
polyhedral grain, the presence of twins as well as the
grain-size heterogeneity in material A. Two types of
grain-boundary morphology were observed with light
microscopy. Straight boundaries with a percentage share
of approximately 38 % of the total of 254 observed grain
boundaries were located in the structure of material A.
Serrated boundaries, as the second type, covered
approximately 39 % of the noticed boundaries. In some
cases, the type of boundary could not be determined
owing to the state of the boundary.
As a reference sample, material B (Figure 3) was
processed with the conventional heat treatment. In
material B, we also observed the base microstructure
I. SLATKOVSKÝ et al.: INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263 AFTER SPECIAL HEAT TREATMENT
722 Materiali in tehnologije / Materials and technology 51 (2017) 5, 721–727
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Table 1: Nominal chemical composition of Alloy 263
Chemical-composition limits
Weight % Ni Cr Co Mo C Al Ti Al+Ti Mn
263 Bal 19–21 19–21 5.6–6.1 0.04–0.08 0.60 max 1.90–2.40 2.40–2.80 0.60 max
Weight % Si Fe Cu B Pb S Ag Bi
263 0.40 max 0.07 max 0.020 max 0.005 max 0.0020 0.007 max 0.005 max 0.0001 max
Figure 1: Net efficiency development in the case of hard coal-fired
power plants considering different structural alloys3
with polyhedral grains with the heterogeneity of the
grain size and twins similar to those in material A. As
expected, no serrated boundaries were observed in
material B after the conventional heat treatment.
The grain-boundary details taken with SEM indicate
the presence of precipitates on the serrated and straight
grain boundaries in material A as well as in material B
(Figures 4 and 5). Results of the EDX analysis taken
from material A (Figure 6) reveal the eventual presence
of carbon-rich secondary particles, which could possibly
indicate the presence of carbides at the grain boundaries.
The occurrence of the other elements at the grain boun-
daries was not significant. Comparable results were also
noticed for material B. To confirm the presence of the
secondary particles formed at the boundaries and to
identify the chemical nature of these particles, ED and
EDX using the TEM were taken for materials A and B,
and the results are shown below.
The authors of reference16 predicted and identified
typical kinds of the secondary particles for this alloy,
allowing us to expect mainly the presence of the MC,
M6C and M23C6 carbides.
Figures 7 to 9 summarize the identified phases on the
replicas for material A. Electron diffraction spectra as
well as the EDX analysis of the extracted precipitates
I. SLATKOVSKÝ et al.: INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263 AFTER SPECIAL HEAT TREATMENT
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Figure 5: SEM, material B: a) straight grain-boundary triple point and
b) detail of precipitatesFigure 3: Material B
Figure 4: SEM, material A: a) serrated grain-boundary triple point
and b) detail of precipitates
Figure 2: Material A
extracted at the grain boundaries confirmed the presence
of carbides. The carbides contained minor elements such
as Mo, Cr and Ni (Table 2).
The M6C ((Co,Ni)3Mo3C) carbide was not confirmed
by electron diffraction (confirmed only by the EDX anal-
I. SLATKOVSKÝ et al.: INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263 AFTER SPECIAL HEAT TREATMENT
724 Materiali in tehnologije / Materials and technology 51 (2017) 5, 721–727
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Figure 8: a) Extracted small particles of irregular shapes at the grain
boundary, material B and b) point-diffraction spectra of MC carbide
Figure 6: EDX map, carbon-rich grain boundary, material A
Figure 9: Typical EDX spectra of precipitates in material A: a)
carbide MC and b) carbide M6C
Figure 7: a) Detail of a precipitate at the grain boundary, material A
and b) point-diffraction spectra of M23C6 carbide
ysis); a slight difference between the lattice parameters
of the M23C6 and M6C (aM23C6 = 0.1065 nm aM6C = 0.1085
nm) carbides makes the identification difficult.
Table 2: Approximate chemical composition of carbides in Material A
M23C6 MC M6C
Avr. (w/%)  Avr. (w/%)  Avr. (w/%) 
Mo 28.6 ±6.1 64.7 ±6.3 22.9 ±5.0
Ti 4.6 ±3.6 26.0 ±5.0 3.3 ±1.3
Cr 66.8 ±9.6 8.8 ±4.3 27.5 ±3.6
Ni 46.3 ±6.8
Details of the serrated boundaries and precipitates in
Material A, observed with TEM on the foils, can be seen
in Figures 10a to 10b. As shown in the figures, the
asymmetry of the serration was observed on some parts
of the serrated boundaries when the irregularity of the
serration exhibited a high difference in  at the observed
parts of the boundary (on the edges,  = 800–900 nm; in
the middle of the boundary,  = 150–200 nm).
On the serrated grain boundary documented in
Figure 10b, secondary particles of different sizes were
noticed. The shape of the particles copied the serration of
the boundary. Using electron diffraction, the particles
were identified as the M23C6 carbide and the matrix as
the -phase (Figure 11). Electron diffraction also
revealed the existence of the orientation relationship
between the matrix and carbide: {200} || {200}M23C6.
Figure 12a documents a straight grain boundary with
discrete secondary particles in material A. Along the full
length of the boundary, the particles of triangular or
rectangular shapes were observed. As shown in Figure
12b, this kind of particles does not copy the shape of the
boundary as in the case of the precipitates at the serrated
boundaries. The size of the precipitates was approxi-
mately 300 nm. The particles were identified, with elec-
tron diffraction (Figure 13), as M23C6 carbide and the
matrix as phase . However, in the case of the straight
grain boundary, no orientation relationship between the
matrix and the precipitate was spotted.
The precipitates extracted at the grain boundary in
material B as M23C6 carbide were identified on the
replicas using electron diffraction, as documented below
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Figure 12: Detail of straight grain boundary, material A, bright field;
b) detail of precipitates at the grain boundary, material A, dark field;
reflection (2 -2 4)M23C6 was used for the dark-field display
Figure 10: Detail of serrated grain boundary, material A, bright field;
b) detail of precipitates at the grain boundary, material A, dark field;
reflection (2 0 0) M23C6 was used for the dark-field display
Figure 11: Point diffraction spectra – phase  and M23C6 carbide
extracted at the grain boundaries confirmed the presence
of carbides. The carbides contained minor elements such
as Mo, Cr and Ni (Table 2).
The M6C ((Co,Ni)3Mo3C) carbide was not confirmed
by electron diffraction (confirmed only by the EDX anal-
I. SLATKOVSKÝ et al.: INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263 AFTER SPECIAL HEAT TREATMENT
724 Materiali in tehnologije / Materials and technology 51 (2017) 5, 721–727
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: a) Extracted small particles of irregular shapes at the grain
boundary, material B and b) point-diffraction spectra of MC carbide
Figure 6: EDX map, carbon-rich grain boundary, material A
Figure 9: Typical EDX spectra of precipitates in material A: a)
carbide MC and b) carbide M6C
Figure 7: a) Detail of a precipitate at the grain boundary, material A
and b) point-diffraction spectra of M23C6 carbide
ysis); a slight difference between the lattice parameters
of the M23C6 and M6C (aM23C6 = 0.1065 nm aM6C = 0.1085
nm) carbides makes the identification difficult.
Table 2: Approximate chemical composition of carbides in Material A
M23C6 MC M6C
Avr. (w/%)  Avr. (w/%)  Avr. (w/%) 
Mo 28.6 ±6.1 64.7 ±6.3 22.9 ±5.0
Ti 4.6 ±3.6 26.0 ±5.0 3.3 ±1.3
Cr 66.8 ±9.6 8.8 ±4.3 27.5 ±3.6
Ni 46.3 ±6.8
Details of the serrated boundaries and precipitates in
Material A, observed with TEM on the foils, can be seen
in Figures 10a to 10b. As shown in the figures, the
asymmetry of the serration was observed on some parts
of the serrated boundaries when the irregularity of the
serration exhibited a high difference in  at the observed
parts of the boundary (on the edges,  = 800–900 nm; in
the middle of the boundary,  = 150–200 nm).
On the serrated grain boundary documented in
Figure 10b, secondary particles of different sizes were
noticed. The shape of the particles copied the serration of
the boundary. Using electron diffraction, the particles
were identified as the M23C6 carbide and the matrix as
the -phase (Figure 11). Electron diffraction also
revealed the existence of the orientation relationship
between the matrix and carbide: {200} || {200}M23C6.
Figure 12a documents a straight grain boundary with
discrete secondary particles in material A. Along the full
length of the boundary, the particles of triangular or
rectangular shapes were observed. As shown in Figure
12b, this kind of particles does not copy the shape of the
boundary as in the case of the precipitates at the serrated
boundaries. The size of the precipitates was approxi-
mately 300 nm. The particles were identified, with elec-
tron diffraction (Figure 13), as M23C6 carbide and the
matrix as phase . However, in the case of the straight
grain boundary, no orientation relationship between the
matrix and the precipitate was spotted.
The precipitates extracted at the grain boundary in
material B as M23C6 carbide were identified on the
replicas using electron diffraction, as documented below
I. SLATKOVSKÝ et al.: INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263 AFTER SPECIAL HEAT TREATMENT
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Figure 12: Detail of straight grain boundary, material A, bright field;
b) detail of precipitates at the grain boundary, material A, dark field;
reflection (2 -2 4)M23C6 was used for the dark-field display
Figure 10: Detail of serrated grain boundary, material A, bright field;
b) detail of precipitates at the grain boundary, material A, dark field;
reflection (2 0 0) M23C6 was used for the dark-field display
Figure 11: Point diffraction spectra – phase  and M23C6 carbide
in Figure 14. Other precipitates were not identified,
which does not exclude their presence in Material B.
From the results of the EDX analysis (Figure 15,
Table 3) of the observed precipitates in Material B, the
presence of MC and M6C carbides is also possible.
Figures 16a to 16b document details of straight
boundaries and precipitates at the grain boundary in
material B on the foils. It can be noticed that the
secondary particles extracted at the grain boundaries
have a polyhedral or rectangular shape, and none of the
observed particles copies the shape of the boundary. The
size of these precipitates is in range of 100–200 nm.
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Figure 13: Point diffraction spectra – phase  and M23C6 carbide
Figure 16: Detail of straight grain boundary, material B, bright field;
b) detail of precipitates at the grain boundary, material B, dark field;
reflection (5 -1 -3)M23C6 was used for the dark-field display
Figure 14: Extracted small particle of an irregular shape at the grain
boundary, material A; b) point diffraction spectra of M23C6 carbide
Figure 15: Typical EDX spectra of precipitates in material B: a)
M23C6 carbide, b) M6C carbide
Table 3: Approximate chemical composition of carbides in material B
M23C6 MC M6C
Avr. (w/%)  Avr. (w/%)  Avr. (w/%) 
Mo 27.3 ±6.4 62.9 ±5.6 24.6 ±3.2
Ti 4.3 ±4.9 25.0 ±4.7 4.1 ±1.8
Cr 61.7 ±9.4 10.5 ±3.5 26.0 ±3.1
Ni 42.6 ±5.1
From the diffraction spectra (Figure 17), secondary
particles at the grain boundaries were identified as M23C6
carbide. No relationship between the matrix and preci-
pitate was observed in all the diffraction spectra in
material B.
4 CONCLUSION
To study the serration and precipitates at the grain
boundaries and identify secondary particles in Alloy 263,
both SEM and TEM were used in the research. To form
serrated grain boundaries, material A was heat treated
under special conditions, based on the studies from the
literature. Material A was compared with material B,
which underwent the conventional heat treatment based
on the material list for Alloy 263. The experimental
results can be summarized as follows:
• Serrated and straight grain boundaries were observed
in material A. Approximately 39 % of the observed
boundaries were serrated. In material B, only straight
grain boundaries were observed.
• In both cases, precipitates at the grain boundaries
were identified as carbide M23C6 using electron
diffraction. In material A, carbide MC was also ob-
served.
• The EDX analysis revealed three types of carbides,
based on their chemical compositions, in both ma-
terials: MC, M6C and M23C6.
• Electron diffraction also revealed the orientation
relationship between the matrix and the precipitate at
the serrated grain boundaries: {200} {200}M23C6.
No orientation relationship was observed at the
straight grain boundaries.
Acknowledgments
The authors wish to thank the financial support of the
Slovak Republic Scientific Grant Agency (VEGA)
within Grant No. 1/0402/13.
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MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 17: Point diffraction spectra – phase  and M23C6 carbide
in Figure 14. Other precipitates were not identified,
which does not exclude their presence in Material B.
From the results of the EDX analysis (Figure 15,
Table 3) of the observed precipitates in Material B, the
presence of MC and M6C carbides is also possible.
Figures 16a to 16b document details of straight
boundaries and precipitates at the grain boundary in
material B on the foils. It can be noticed that the
secondary particles extracted at the grain boundaries
have a polyhedral or rectangular shape, and none of the
observed particles copies the shape of the boundary. The
size of these precipitates is in range of 100–200 nm.
I. SLATKOVSKÝ et al.: INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263 AFTER SPECIAL HEAT TREATMENT
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Figure 13: Point diffraction spectra – phase  and M23C6 carbide
Figure 16: Detail of straight grain boundary, material B, bright field;
b) detail of precipitates at the grain boundary, material B, dark field;
reflection (5 -1 -3)M23C6 was used for the dark-field display
Figure 14: Extracted small particle of an irregular shape at the grain
boundary, material A; b) point diffraction spectra of M23C6 carbide
Figure 15: Typical EDX spectra of precipitates in material B: a)
M23C6 carbide, b) M6C carbide
Table 3: Approximate chemical composition of carbides in material B
M23C6 MC M6C
Avr. (w/%)  Avr. (w/%)  Avr. (w/%) 
Mo 27.3 ±6.4 62.9 ±5.6 24.6 ±3.2
Ti 4.3 ±4.9 25.0 ±4.7 4.1 ±1.8
Cr 61.7 ±9.4 10.5 ±3.5 26.0 ±3.1
Ni 42.6 ±5.1
From the diffraction spectra (Figure 17), secondary
particles at the grain boundaries were identified as M23C6
carbide. No relationship between the matrix and preci-
pitate was observed in all the diffraction spectra in
material B.
4 CONCLUSION
To study the serration and precipitates at the grain
boundaries and identify secondary particles in Alloy 263,
both SEM and TEM were used in the research. To form
serrated grain boundaries, material A was heat treated
under special conditions, based on the studies from the
literature. Material A was compared with material B,
which underwent the conventional heat treatment based
on the material list for Alloy 263. The experimental
results can be summarized as follows:
• Serrated and straight grain boundaries were observed
in material A. Approximately 39 % of the observed
boundaries were serrated. In material B, only straight
grain boundaries were observed.
• In both cases, precipitates at the grain boundaries
were identified as carbide M23C6 using electron
diffraction. In material A, carbide MC was also ob-
served.
• The EDX analysis revealed three types of carbides,
based on their chemical compositions, in both ma-
terials: MC, M6C and M23C6.
• Electron diffraction also revealed the orientation
relationship between the matrix and the precipitate at
the serrated grain boundaries: {200} {200}M23C6.
No orientation relationship was observed at the
straight grain boundaries.
Acknowledgments
The authors wish to thank the financial support of the
Slovak Republic Scientific Grant Agency (VEGA)
within Grant No. 1/0402/13.
5 REFERENCES
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Romanosky, Defining the materials issues and research needs for
ultra-supercritical steam turbines, Proc. 4th Inter. Conf. Advances in
Materials Technology for Fossil Power Plants, Hilton Head Island,
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2 R. J. Campbell, Increasing the Efficiency of Existing Coal-Fired
Power Plants, https://fas.org/sgp/crs/misc/R43343.pdf, 14.06.2017
3 G. Stein-Brzozowska, D. M. Flórez, J. Maier, G. Scheffknecht,
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plants: Fireside corrosion. Laboratory studies and power plant expo-
sures, Fuel, 108, (2013), 521–533, doi:10.1016/j.fuel.2012.11.081
4 A. K. Koul, G. H. Gessinger, On the mechanism of serrated grain
boundary formation in Ni-based superalloys, Acta Metallurgica, 31,
(1983) 7, 1061–1069, doi:10.1016/0001-6160(83)90202-X
5 H. L. Danflou, M. Marty, A. Walder, Formation of serrated grain
boundaries and their effect on the mechanical properties in a P/M
nickel base superalloy, Proc. 7th Inter. Symp. on Superalloys, Seven
Springs Mountain, USA, 1992, 63–72, doi:10.7449/1992/ Super-
alloys_1992_63_72
6 R. J. Mitchell, H. Y. Li, Z. W. Huang, On the formation of serrated
grain boundaries and fan type structures in an advanced polycry-
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03.008
7 C. L. Qiu, P. Andrews, On the formation of irregular-shaped gamma
prime and serrated grain boundaries in a nickel-based superalloy
during continuous cooling, Materials Characterization, 76 (2013),
28–34, doi:10.1016/j.matchar.2012.11.012
8 A. C. Yeh, K. W. Lu, C. M. Kuo, H. Y. Bor, C. N. Wei, Effect of
Serrated Grain Boundaries on the Creep Property of Inconel 718
Superalloy, Materials Science and Engineering A, 530 (2011),
525–529, doi:10.1016/j.msea.2011.10.014
9 D. H. Ping, Y. F. Gu, C. Y. Cui, H. Harada, Grain boundary segre-
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2007.01.090
10 L. Jiang, R. Hu, H. Kou, J. Li, G. Bai, H. Fu, The effect of M23C6
carbides on the formation of grain boundary serrations in a wrought
Ni-based superalloy, Materials Science and Engineering A, 536
(2012), 37–44, doi:10.1016/j.msea.2011.11.060
11 Y. S. Lim, D. J. Kim, S. S. Hwang, H. P. Kim, S. W. Kim, M23C6
precipitation behavior and grain boundary serration in Ni-based
Alloy 690, Materials Characterization, 96 (2014), 28–39,
doi:10.1016/j.matchar.2014.07.008
12 H. U. Hong, I. S. Kim, B. G. Choi, M. Y. Kim, C. Y. Jo, The effect of
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13 H. U. Hong, I. S. Kim, B. G. Choi, Y. S. Yoo, C. Y. Jo, On the role of
grain boundary serration in simulated weld heat-affected zone
liquation of a wrought nickel-based superalloy, Metallurgical and
Materials Transactions A, 43 (2012) 1, 173–181, doi:10.1007/
s11661-011-0837-2
14 H. U. Hong, I. S. Kim, B. G. Choi, Y. S. Yoo, C. Y. Jo, On the
Mechanism of Serrated Grain Boundary Formation in Ni-Based
Superalloys with Low '? Volume Fraction, Proc. 12th Inter. Symp.
on Superalloys, Seven Springs Mountain, USA, 2012, 53–61,
doi:10.1002/9781118516430.ch6
15 H. U. Hong, F. H. Latief, T. Blanc, I. S. Kim, B. G. Choi, C. Y. Jo, J.
H. Lee, Influence of chromium content on microstructure and grain
boundary serration formation in a ternary Ni-xCr-0.1C model alloy,
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16 J. C. Zhao, V. Ravikumar, A. M. Beltran, Phase Precipitation and
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I. SLATKOVSKÝ et al.: INVESTIGATION OF GRAIN BOUNDARIES IN ALLOY 263 AFTER SPECIAL HEAT TREATMENT
Materiali in tehnologije / Materials and technology 51 (2017) 5, 721–727 727
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 17: Point diffraction spectra – phase  and M23C6 carbide
J. PTA^INOVÁ et al.: FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL AFTER SUB-ZERO TREATMENT ...
729–733
FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL
AFTER SUB-ZERO TREATMENT FOR 17 H AND DOUBLE
TEMPERING
LOMNA @ILAVOST LEDEBURITNEGA JEKLA VANADIS 6 PO
TOPLOTNI OBDELAVI S 17-URNIM PODHLAJEVANJEM IN
DVOJNIM POPU[^ANJEM
Jana Pta~inová1, Peter Jur~i1, Ivo Dlouhý2
1Institute of Materials Science, Faculty of Materials Science and Technology Trnava, Jána Bottu 25, 917 24 Trnava, Slovakia
2Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Zizkova 22, 61662 Brno, Czech Republic
jana.ptacinova@stuba.sk
Prejem rokopisa – received: 2016-06-20; sprejem za objavo – accepted for publication: 2017-03-01
doi:10.17222/mit.2016.118
The influence of sub-zero treatment on the fracture toughness of Cr-V ledeburitic steel Vanadis 6 was examined, in comparison
with the same material processed without the sub-zero period. The microstructure of the material consists of the matrix and
several types of carbides – eutectic carbides (ECs), secondary carbides (SCs) and small globular carbides (SGCs). Small
amounts of retained austenite were also present in the microstructure, but only in the untempered or low-temperature tempered
steel. Sub-zero treatment increases the amount of small globular carbides. On the other hand, tempering results in a decrease in
the population density of these particles. The fracture toughness of conventionally heat-treated steel firstly increases with the
tempering, but then it decreases rapidly when the steel is tempered at the temperature of secondary hardening. In the case of the
sub-zero treated material, the fracture toughness is correspondingly lower when the material is tempered at low temperatures,
but it becomes slightly higher in the temperature range normally used for secondary hardening. Generally, one can say that the
fracture toughness follows well the values of the hardness of the material, except in the narrow temperature range in the case of
the sub-zero treated steel, where a "window" for a simultaneous enhancement of hardness and toughness exists.
Keywords: ledeburitic steel, sub-zero treatment, fracture toughness, carbides, fracture surface
Avtorji so preu~evali vpliv standardne toplotne obdelave v kombinaciji s podhlajevanjem na lomno `ilavost Cr-V ledeburitnega
jekla Vanadis 6 v primerjavi z enakim materialom brez podhlajevanja. Mikrostruktura materiala sestoji iz matrice in ve~ vrst
karbidov – evtekti~nih karbidov (angl. ECs), sekundarnih karbidov (angl. SCs), in manj{ih globularnih karbidov (angl. SGCs).
Majhne koli~ine zaostalega avstenita so bile prisotne v mikrostrukturi, vendar le v nepopu{~enem ali v nizkotemperaturno-
popu{~enenem jeklu. Podhlajevanje pove~uje koli~ino manj{ih globularnih karbidov. Po drugi strani pa se s popu{~anjem
zmanj{uje populacijska gostota teh delcev. Lomna `ilavost konvencionalno toplotno obdelanega jekla se najprej pove~uje s
povi{evanjem temperature popu{~anja vendar se pri~ne hitro zmanj{evati (zni`evati), ko je prekora~ena temperatura sekun-
darnega utrjevanja. Pri toplotni obdelavi materiala s podhlajevanjem, se lomna `ilavost posledi~no zni`uje, ko je le-to
popu{~eno na ni`jih temperaturah, vendar se rahlo pove~uje v obmo~ju temperature, ki se navadno uporablja za sekundarno
utrjevanje. V splo{nem je mo`no re~i, da lomna `ilavost sledi vrednostim trdote materiala, razen v primeru toplotno obdelanega
jekla v kombinaciji s podhlajevanjem, kjer v zelo ozkem temperaturnem intervalu obstaja "temperaturno okno". Tam dose`emo
isto~asno pove~anje `ilavosti in trdote.
Klju~ne besede: ledeburitno jeklo, podhlajevanje, lomna `ilavost, karbidi, lomna povr{ina
1 INTRODUCTION
In modern tooling, it is necessary that the tool mate-
rials have a good wear performance and high hardness,
accompanied with at least acceptable toughness. Other-
wise the tools might fail very early, before any wear
damage can occur.
The toughness of steels quantifies their resistance to
the initiation of brittle fracture under given conditions.
For real tools made of Cr- and Cr-V ledeburitic steels,
the toughness determines their capability to resist either
the chipping or a total failure. The toughness, being
expressed by the three-point bending strength, b, of
brittle ledeburitic steels is the highest when the material
is soft-annealed and it decreases as a consequence of the
application of heat treatment. However, the higher the
austenitizing temperature (TA), the lower is the toughness
and, at a constant TA, the tempering first results in an in-
crease in the toughness, which is followed by its
decrease at the maximum secondary hardening.1 Hence,
the application of proper heat treatment is a compromise
between the requirements for the maximum hardness and
at least acceptable toughness.
There is a shortage of literature pertaining to the
effect of the sub-zero treatment (SZT) on the toughness.
D. N. Collins and J. Dormer reported that the toughness
of the AISI D2 steel decreased with an application of the
SZT up to the temperature of –70 °C, which was
followed by a moderate increase in the toughness when a
lower processing temperature was used for the SZT.2,3 It
should be noticed here that the samples were low-tem-
perature tempered at 200 °C.
Materiali in tehnologije / Materials and technology 51 (2017) 5, 729–733 729
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.178.2:669.15-194.58:691.714:62-712 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)729(2017)
J. PTA^INOVÁ et al.: FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL AFTER SUB-ZERO TREATMENT ...
729–733
FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL
AFTER SUB-ZERO TREATMENT FOR 17 H AND DOUBLE
TEMPERING
LOMNA @ILAVOST LEDEBURITNEGA JEKLA VANADIS 6 PO
TOPLOTNI OBDELAVI S 17-URNIM PODHLAJEVANJEM IN
DVOJNIM POPU[^ANJEM
Jana Pta~inová1, Peter Jur~i1, Ivo Dlouhý2
1Institute of Materials Science, Faculty of Materials Science and Technology Trnava, Jána Bottu 25, 917 24 Trnava, Slovakia
2Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Zizkova 22, 61662 Brno, Czech Republic
jana.ptacinova@stuba.sk
Prejem rokopisa – received: 2016-06-20; sprejem za objavo – accepted for publication: 2017-03-01
doi:10.17222/mit.2016.118
The influence of sub-zero treatment on the fracture toughness of Cr-V ledeburitic steel Vanadis 6 was examined, in comparison
with the same material processed without the sub-zero period. The microstructure of the material consists of the matrix and
several types of carbides – eutectic carbides (ECs), secondary carbides (SCs) and small globular carbides (SGCs). Small
amounts of retained austenite were also present in the microstructure, but only in the untempered or low-temperature tempered
steel. Sub-zero treatment increases the amount of small globular carbides. On the other hand, tempering results in a decrease in
the population density of these particles. The fracture toughness of conventionally heat-treated steel firstly increases with the
tempering, but then it decreases rapidly when the steel is tempered at the temperature of secondary hardening. In the case of the
sub-zero treated material, the fracture toughness is correspondingly lower when the material is tempered at low temperatures,
but it becomes slightly higher in the temperature range normally used for secondary hardening. Generally, one can say that the
fracture toughness follows well the values of the hardness of the material, except in the narrow temperature range in the case of
the sub-zero treated steel, where a "window" for a simultaneous enhancement of hardness and toughness exists.
Keywords: ledeburitic steel, sub-zero treatment, fracture toughness, carbides, fracture surface
Avtorji so preu~evali vpliv standardne toplotne obdelave v kombinaciji s podhlajevanjem na lomno `ilavost Cr-V ledeburitnega
jekla Vanadis 6 v primerjavi z enakim materialom brez podhlajevanja. Mikrostruktura materiala sestoji iz matrice in ve~ vrst
karbidov – evtekti~nih karbidov (angl. ECs), sekundarnih karbidov (angl. SCs), in manj{ih globularnih karbidov (angl. SGCs).
Majhne koli~ine zaostalega avstenita so bile prisotne v mikrostrukturi, vendar le v nepopu{~enem ali v nizkotemperaturno-
popu{~enenem jeklu. Podhlajevanje pove~uje koli~ino manj{ih globularnih karbidov. Po drugi strani pa se s popu{~anjem
zmanj{uje populacijska gostota teh delcev. Lomna `ilavost konvencionalno toplotno obdelanega jekla se najprej pove~uje s
povi{evanjem temperature popu{~anja vendar se pri~ne hitro zmanj{evati (zni`evati), ko je prekora~ena temperatura sekun-
darnega utrjevanja. Pri toplotni obdelavi materiala s podhlajevanjem, se lomna `ilavost posledi~no zni`uje, ko je le-to
popu{~eno na ni`jih temperaturah, vendar se rahlo pove~uje v obmo~ju temperature, ki se navadno uporablja za sekundarno
utrjevanje. V splo{nem je mo`no re~i, da lomna `ilavost sledi vrednostim trdote materiala, razen v primeru toplotno obdelanega
jekla v kombinaciji s podhlajevanjem, kjer v zelo ozkem temperaturnem intervalu obstaja "temperaturno okno". Tam dose`emo
isto~asno pove~anje `ilavosti in trdote.
Klju~ne besede: ledeburitno jeklo, podhlajevanje, lomna `ilavost, karbidi, lomna povr{ina
1 INTRODUCTION
In modern tooling, it is necessary that the tool mate-
rials have a good wear performance and high hardness,
accompanied with at least acceptable toughness. Other-
wise the tools might fail very early, before any wear
damage can occur.
The toughness of steels quantifies their resistance to
the initiation of brittle fracture under given conditions.
For real tools made of Cr- and Cr-V ledeburitic steels,
the toughness determines their capability to resist either
the chipping or a total failure. The toughness, being
expressed by the three-point bending strength, b, of
brittle ledeburitic steels is the highest when the material
is soft-annealed and it decreases as a consequence of the
application of heat treatment. However, the higher the
austenitizing temperature (TA), the lower is the toughness
and, at a constant TA, the tempering first results in an in-
crease in the toughness, which is followed by its
decrease at the maximum secondary hardening.1 Hence,
the application of proper heat treatment is a compromise
between the requirements for the maximum hardness and
at least acceptable toughness.
There is a shortage of literature pertaining to the
effect of the sub-zero treatment (SZT) on the toughness.
D. N. Collins and J. Dormer reported that the toughness
of the AISI D2 steel decreased with an application of the
SZT up to the temperature of –70 °C, which was
followed by a moderate increase in the toughness when a
lower processing temperature was used for the SZT.2,3 It
should be noticed here that the samples were low-tem-
perature tempered at 200 °C.
Materiali in tehnologije / Materials and technology 51 (2017) 5, 729–733 729
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.178.2:669.15-194.58:691.714:62-712 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)729(2017)
The fracture toughness (KIC) (the resistance against
the crack propagation) of the ledeburitic steels can be
quantified using either pre-cracked (bend) specimens or
chevron-notch technique.4–6 The fracture toughness of
ledeburitic steels is very low, whereas the KIC values
usually follow the three-point bending strength.
The effect of the SZT on the fracture toughness was
reported by several groups of investigators. D. Das et al.,
for instance, pointed out that the SZT carried out at tem-
peratures of (–75, –125 and –196) °C led to a decrease in
the fracture toughness compared to conventionally
heat-treated samples when low-temperature tempered at
210 °C for 2 h.6 The variations in the fracture toughness
were attributed to the reduction of the retained-austenite
amount and to the increase in the population density of
SGCs in the microstructure after the SZT. In our recent
paper, it was demonstrated that the effect of the SZT on
the toughness is marginal, but the effect of this process-
ing on KIC is rather positive when the material is
tempered at the temperature normally used for secondary
hardening.7
The main goal of the current investigations is to
characterize the variations in the toughness and fracture
toughness of powder-metallurgy (PM) ledeburitic steel
Vanadis 6, as a result of different heat-treatment sche-
dules (austenitizing, sub-zero treatment, tempering) and
to relate them to the microstructural alterations, due to
these treatments, like the reduction of the retained-auste-
nite amount, increase in the carbide count and others.8–10
2 MATERIAL AND EXPERIMENTAL METHODS
2.1 Material and processing
The experimental material was tool steel Vanadis 6
with nominally 2.1 % C, 1.0 % Si, 0.4 % Mn, 6.8 % Cr,
1.5 % Mo, 5.4 % V and Fe as balance, made with PM.
The conventional heat treatment (CHT) consisted of the
following steps: heating up the material to the desired TA
(1050 °C) in a vacuum furnace, holding it at the tem-
perature for 30 min and nitrogen gas quenching (5 bar).
One set of samples was processed without the inclusion
of the SZT between quenching and tempering, while the
other samples were subjected to the SZT carried out at
the temperature of -196 °C for 17 h. Double tempering
(2 h + 2 h) was performed at temperatures in the range of
100–600 °C.
2.2 Experimental methods
Metallographic samples were prepared in the stan-
dard way and etched with a picric-acid reagent for the
SEM observation. The microstructure was recorded
using a JEOL JSM 7600 F device equipped with an EDS
detector (Oxford Instruments), at an acceleration voltage
of 15 kV. Details of the categorization of the carbides
have been published recently.10,11 Macro-hardness
measurements were completed using the Vickers (HV10)
method. Five measurements were made on the metallo-
graphic specimens processed with any combination of
heat-treatment parameters, and both the mean values and
the standard deviations were then calculated. For the
fracture-toughness determination, pre-cracked specimens
predetermined for bending, with dimensions of 10 mm ×
10 mm × 55 mm were used. For the pre-cracking of the
samples, four bending fixtures and a resonance freq-
uency machine (Cracktronic 8024) were used, which
allowed the fatigue-crack initiation and further propa-
gation through the actual frequency of the cycling to be
controlled.
In addition, the crack development was checked on
both sides of a sample using digital long-distance micro-
scopic methods. Both the pre-crack preparation and bend
tests were carried out at room temperature according to
the ISO12137 standard.12 For the test, an Instron 8862
machine was used and a loading rate of 0.1 mm/min was
applied. Specimen deflection was measured by means of
an inductive transducer integrated directly into the
loading axis. In total, five samples were tested for each
investigated condition. Fracture-surface morphology was
investigated with a scanning electron microscope JEOL
JSM 7600F with an EDS detector (Oxford Instruments).
The topography of fracture surfaces was studied with a
confocal laser scanning microscope Zeiss LSM 700.
Three-dimensional topographical resolution was
achieved using the ZEN 2009 software.
3 RESULTS AND DISCUSSION
3.1 Microstructure
SEM micrographs, Figure 1, show an example of the
microstructure evolution of the material after the
sub-zero treatment in liquid nitrogen (for 17 h) and
tempering at different temperatures. The microstructure
after the sub-zero treatment consists of a matrix made of
martensite, a small amount of retained austenite and
carbides (Figure 1a). As expected, the character of the
matrix microstructure changes with increasing tempering
temperature (Figures 1b to 1f). Due to the tempering,
the martensite becomes more sensitive to the etching
agent (the so-called tempered martensite). This is a com-
monly known fact and it is related to the precipitation of
carbides during tempering. Retained austenite is trans-
formed into martensite. The portion of carbides is also
changed due to the tempering. The volume fractions of
both eutectic carbides and secondary carbides are
invariant over the range of the heat-treatment parameters
used in the experiments. The population density of the
small globular carbides increased with the application of
the sub-zero treatment, but rapidly decreased with the in-
creasing temperature of the tempering (Figure 2).
3.2 Hardness characteristics
The bulk hardness of the non-SZT and SZT Vanadis
6 steel, as a function of the tempering temperature, is
J. PTA^INOVÁ et al.: FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL AFTER SUB-ZERO TREATMENT ...
730 Materiali in tehnologije / Materials and technology 51 (2017) 5, 729–733
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
shown in Table 1. The as-quenched hardness of the
conventionally heat-treated steel that was not tempered
was 838±4.33 HV10. The hardness of the SZT steel
soaked in liquid nitrogen for 17 h and not tempered was
correspondingly higher, e.g., 917±7.58 HV10. These
results show that the as-quenched bulk hardness of the
Vanadis 6 steel was improved due to the sub-zero treat-
ment. The hardness of all the samples then decreases
with the increasing tempering temperature. Here it is
interesting that the hardness of the SZT steel remains
higher up to the tempering temperature of 450 °C and
then it drops more intensely than that of the convention-
ally quenched samples. As a result, the hardness of the
specimens tempered in the range of the temperatures
commonly used for secondary hardening is lower for the
SZT steel than for the steel achieved after conventional
quenching and tempering.
3.3 Evaluation of the SZT effect on the fracture tough-
ness
The fracture toughness vs. heat treatment schedules
are presented in Figure 3a (CHT) and Figure 3b (SZT).
In the case of the CHT, the fracture toughness of the
untempered material was measured to be 16.4 MPa m1/2
and it firstly increased, due to the tempering, to 19.6
MPa m1/2 with a subsequent slight decrease; however, the
decrease in the fracture toughness markedly accelerated,
at the temperature normally used for secondary harden-
ing, to a value of 14.8 MPa m1/2. For the material
subjected to the SZT in liquid nitrogen for 17 h, the
fracture toughness before tempering was 13.3 MPa m1/2
(i.e., lower than that of the CHT steel). Then, the fracture
toughness increased with the tempering temperature to
more than 15 MPa m1/2 and just slightly decreased at a
tempering temperature of 450 °C. It is interesting that
the KIC values were higher for the SZT steel than for the
CHT steel at the temperature of secondary hardening
(530 °C).
Table 1: Hardness of the Vanadis 6 ledeburitic steel
Tempering of samples
CHT SZT
HV10 ± HV10 ±
untempered 838 4.3 917 7.6
tempered at 170 °C 784 3.7 870 5.4
tempered at 330 °C 736 7 851 15.7
tempered at 450 °C 724 3 816 7.8
tempered at 530 °C 758 5.9 701 3
The behaviour of fracture toughness vs. heat treat-
ment schedules can be classified as logical because one
can expect higher KIC at lower hardness and lower KIC at
elevated hardness. In other words, the application of the
SZT decreases KIC when the steel is low-temperature
tempered. This fact should be considered by toolmakers
and users of tools in all the cases when they expect an
increase in the wear resistance due to a high hardness
resulting from the sub-zero treatment. A very interesting
fact was found for the material tempered at the
temperature of secondary hardening – KIC was higher for
the SZT material. This is in good agreement with the
recently published results where a similar tendency was
found for the same SZT steel treated for 4 h and 10 h.7
One can say that this kind of KIC behaviour can be
expected because of the lower hardness of the SZT steel
when tempered in this temperature range; however, in
the mentioned paper, it was also found that the increase
of KIC is accompanied with a better wear performance of
the material, due to the presence of a higher amount of
carbides. The enhanced number of carbides, compared to
the material after the conventional quenching, was also
identified in the current work. Hence, one can expect that
a sub-zero treated material would have a better wear
performance, too. These results are also interesting from
J. PTA^INOVÁ et al.: FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL AFTER SUB-ZERO TREATMENT ...
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MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Amount of carbides in the Vanadis 6 steel after conventional
quenching, SZT for 17 h and SZT and subsequent tempering at
different temperatures
Figure 1: SEM micrographs showing the microstructure of Vanadis 6
ledeburitic steel: a) after quenching, subsequent SZT and b) tempering
at 170 °C, c) 330 °C, d) 450 °C, e) 530 °C, f) 600 °C
The fracture toughness (KIC) (the resistance against
the crack propagation) of the ledeburitic steels can be
quantified using either pre-cracked (bend) specimens or
chevron-notch technique.4–6 The fracture toughness of
ledeburitic steels is very low, whereas the KIC values
usually follow the three-point bending strength.
The effect of the SZT on the fracture toughness was
reported by several groups of investigators. D. Das et al.,
for instance, pointed out that the SZT carried out at tem-
peratures of (–75, –125 and –196) °C led to a decrease in
the fracture toughness compared to conventionally
heat-treated samples when low-temperature tempered at
210 °C for 2 h.6 The variations in the fracture toughness
were attributed to the reduction of the retained-austenite
amount and to the increase in the population density of
SGCs in the microstructure after the SZT. In our recent
paper, it was demonstrated that the effect of the SZT on
the toughness is marginal, but the effect of this process-
ing on KIC is rather positive when the material is
tempered at the temperature normally used for secondary
hardening.7
The main goal of the current investigations is to
characterize the variations in the toughness and fracture
toughness of powder-metallurgy (PM) ledeburitic steel
Vanadis 6, as a result of different heat-treatment sche-
dules (austenitizing, sub-zero treatment, tempering) and
to relate them to the microstructural alterations, due to
these treatments, like the reduction of the retained-auste-
nite amount, increase in the carbide count and others.8–10
2 MATERIAL AND EXPERIMENTAL METHODS
2.1 Material and processing
The experimental material was tool steel Vanadis 6
with nominally 2.1 % C, 1.0 % Si, 0.4 % Mn, 6.8 % Cr,
1.5 % Mo, 5.4 % V and Fe as balance, made with PM.
The conventional heat treatment (CHT) consisted of the
following steps: heating up the material to the desired TA
(1050 °C) in a vacuum furnace, holding it at the tem-
perature for 30 min and nitrogen gas quenching (5 bar).
One set of samples was processed without the inclusion
of the SZT between quenching and tempering, while the
other samples were subjected to the SZT carried out at
the temperature of -196 °C for 17 h. Double tempering
(2 h + 2 h) was performed at temperatures in the range of
100–600 °C.
2.2 Experimental methods
Metallographic samples were prepared in the stan-
dard way and etched with a picric-acid reagent for the
SEM observation. The microstructure was recorded
using a JEOL JSM 7600 F device equipped with an EDS
detector (Oxford Instruments), at an acceleration voltage
of 15 kV. Details of the categorization of the carbides
have been published recently.10,11 Macro-hardness
measurements were completed using the Vickers (HV10)
method. Five measurements were made on the metallo-
graphic specimens processed with any combination of
heat-treatment parameters, and both the mean values and
the standard deviations were then calculated. For the
fracture-toughness determination, pre-cracked specimens
predetermined for bending, with dimensions of 10 mm ×
10 mm × 55 mm were used. For the pre-cracking of the
samples, four bending fixtures and a resonance freq-
uency machine (Cracktronic 8024) were used, which
allowed the fatigue-crack initiation and further propa-
gation through the actual frequency of the cycling to be
controlled.
In addition, the crack development was checked on
both sides of a sample using digital long-distance micro-
scopic methods. Both the pre-crack preparation and bend
tests were carried out at room temperature according to
the ISO12137 standard.12 For the test, an Instron 8862
machine was used and a loading rate of 0.1 mm/min was
applied. Specimen deflection was measured by means of
an inductive transducer integrated directly into the
loading axis. In total, five samples were tested for each
investigated condition. Fracture-surface morphology was
investigated with a scanning electron microscope JEOL
JSM 7600F with an EDS detector (Oxford Instruments).
The topography of fracture surfaces was studied with a
confocal laser scanning microscope Zeiss LSM 700.
Three-dimensional topographical resolution was
achieved using the ZEN 2009 software.
3 RESULTS AND DISCUSSION
3.1 Microstructure
SEM micrographs, Figure 1, show an example of the
microstructure evolution of the material after the
sub-zero treatment in liquid nitrogen (for 17 h) and
tempering at different temperatures. The microstructure
after the sub-zero treatment consists of a matrix made of
martensite, a small amount of retained austenite and
carbides (Figure 1a). As expected, the character of the
matrix microstructure changes with increasing tempering
temperature (Figures 1b to 1f). Due to the tempering,
the martensite becomes more sensitive to the etching
agent (the so-called tempered martensite). This is a com-
monly known fact and it is related to the precipitation of
carbides during tempering. Retained austenite is trans-
formed into martensite. The portion of carbides is also
changed due to the tempering. The volume fractions of
both eutectic carbides and secondary carbides are
invariant over the range of the heat-treatment parameters
used in the experiments. The population density of the
small globular carbides increased with the application of
the sub-zero treatment, but rapidly decreased with the in-
creasing temperature of the tempering (Figure 2).
3.2 Hardness characteristics
The bulk hardness of the non-SZT and SZT Vanadis
6 steel, as a function of the tempering temperature, is
J. PTA^INOVÁ et al.: FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL AFTER SUB-ZERO TREATMENT ...
730 Materiali in tehnologije / Materials and technology 51 (2017) 5, 729–733
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
shown in Table 1. The as-quenched hardness of the
conventionally heat-treated steel that was not tempered
was 838±4.33 HV10. The hardness of the SZT steel
soaked in liquid nitrogen for 17 h and not tempered was
correspondingly higher, e.g., 917±7.58 HV10. These
results show that the as-quenched bulk hardness of the
Vanadis 6 steel was improved due to the sub-zero treat-
ment. The hardness of all the samples then decreases
with the increasing tempering temperature. Here it is
interesting that the hardness of the SZT steel remains
higher up to the tempering temperature of 450 °C and
then it drops more intensely than that of the convention-
ally quenched samples. As a result, the hardness of the
specimens tempered in the range of the temperatures
commonly used for secondary hardening is lower for the
SZT steel than for the steel achieved after conventional
quenching and tempering.
3.3 Evaluation of the SZT effect on the fracture tough-
ness
The fracture toughness vs. heat treatment schedules
are presented in Figure 3a (CHT) and Figure 3b (SZT).
In the case of the CHT, the fracture toughness of the
untempered material was measured to be 16.4 MPa m1/2
and it firstly increased, due to the tempering, to 19.6
MPa m1/2 with a subsequent slight decrease; however, the
decrease in the fracture toughness markedly accelerated,
at the temperature normally used for secondary harden-
ing, to a value of 14.8 MPa m1/2. For the material
subjected to the SZT in liquid nitrogen for 17 h, the
fracture toughness before tempering was 13.3 MPa m1/2
(i.e., lower than that of the CHT steel). Then, the fracture
toughness increased with the tempering temperature to
more than 15 MPa m1/2 and just slightly decreased at a
tempering temperature of 450 °C. It is interesting that
the KIC values were higher for the SZT steel than for the
CHT steel at the temperature of secondary hardening
(530 °C).
Table 1: Hardness of the Vanadis 6 ledeburitic steel
Tempering of samples
CHT SZT
HV10 ± HV10 ±
untempered 838 4.3 917 7.6
tempered at 170 °C 784 3.7 870 5.4
tempered at 330 °C 736 7 851 15.7
tempered at 450 °C 724 3 816 7.8
tempered at 530 °C 758 5.9 701 3
The behaviour of fracture toughness vs. heat treat-
ment schedules can be classified as logical because one
can expect higher KIC at lower hardness and lower KIC at
elevated hardness. In other words, the application of the
SZT decreases KIC when the steel is low-temperature
tempered. This fact should be considered by toolmakers
and users of tools in all the cases when they expect an
increase in the wear resistance due to a high hardness
resulting from the sub-zero treatment. A very interesting
fact was found for the material tempered at the
temperature of secondary hardening – KIC was higher for
the SZT material. This is in good agreement with the
recently published results where a similar tendency was
found for the same SZT steel treated for 4 h and 10 h.7
One can say that this kind of KIC behaviour can be
expected because of the lower hardness of the SZT steel
when tempered in this temperature range; however, in
the mentioned paper, it was also found that the increase
of KIC is accompanied with a better wear performance of
the material, due to the presence of a higher amount of
carbides. The enhanced number of carbides, compared to
the material after the conventional quenching, was also
identified in the current work. Hence, one can expect that
a sub-zero treated material would have a better wear
performance, too. These results are also interesting from
J. PTA^INOVÁ et al.: FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL AFTER SUB-ZERO TREATMENT ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 729–733 731
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Amount of carbides in the Vanadis 6 steel after conventional
quenching, SZT for 17 h and SZT and subsequent tempering at
different temperatures
Figure 1: SEM micrographs showing the microstructure of Vanadis 6
ledeburitic steel: a) after quenching, subsequent SZT and b) tempering
at 170 °C, c) 330 °C, d) 450 °C, e) 530 °C, f) 600 °C
the point of view of industrial practice. They indicate
that it is possible to increase the wear performance,
along the toughness, of the material in a certain tem-
pering-temperature range.
3.4 Fracture-surface morphology
The micro-mechanics of fracture propagation is
demonstrated through representative SEM micrographs
showing how the fractures appear in the cases of the steel
that was sub-zero treated at –196 °C for 17 h and not
tempered (the lowest KIC) and the steel that was sub-zero
treated at –196 °C for 17 h and tempered at 170 °C
(higher KIC). The fracture surface of the sample that was
not tempered and the one that was tempered at 170 °C
are shown in Figure 4.
At first glance, there is not any difference between
the mentioned fracture surfaces; they appear flat and
shiny, Figures 4a and 4c. However, a more detailed
observation, at higher magnifications, makes it clear that
the fracture surface of the sample after low-temperature
tempering contains an enhanced number of sites with
micro-plastic deformation (PLD), located mainly at the
carbide/matrix interfaces, Figure 4b. On the other hand,
there are very few such sites on the fracture surface of
the untempered specimen, Figure 4d, and the fracture
surface contains a number of cleavage facettes/regions
(CL).
The topographies of the fracture surfaces of the CHT
Vanadis 6 steel tempered at 170 °C and sub-zero treated,
and the untempered one, expressed with surface rough-
ness Ra are shown in Table 2 and Figure 5. The average
roughness of the fracture surface of the conventionally
quenched steel tempered at 170 °C was 4.2703±0.248 mm.
J. PTA^INOVÁ et al.: FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL AFTER SUB-ZERO TREATMENT ...
732 Materiali in tehnologije / Materials and technology 51 (2017) 5, 729–733
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Representative SEM micrographs of the fracture surfaces of
the specimens processed with: a), b) SZT at –196 °C for 17 h and tem-
pered at 170 °C and c), d) SZT at –196 °C for 17 h and not tempered
Figure 3: Hardness HV10 and fracture toughness KIC of Vanadis 6
steel in dependence on the heat treatment applied: a) conventional
quenching, b) sub-zero treatment
Figure 5: Confocal-microscope micrographs showing the fracture surfaces of KIC specimens: a) CHT and tempered at 170 °C, b) SZT and not
tempered
The roughness of the sub-zero treated and untempered
Vanadis 6 steel was lower, i.e., 2.918±0.323 μm. At this
point, it should be noted that in this stage of
experimental efforts, only the materials with the highest
KIC and the one with the lowest KIC were checked with
confocal microscopy. However, the measurements gave
unambiguous differences in the roughness, where the
higher roughness corresponded to the higher KIC, i.e., to
the conventionally quenched and low-temperature tem-
pered sample.
Table 2: Roughness of the fracture surface of Vanadis 6 ledeburitic
steel
No. of measurement Tempered at 170 °C SZT anduntempered
1. 4.136 2.591
2. 4.750 3.507
3. 3.925 2.656
Ra (μm) 4.2703±0.248 2.918±0.323
4 CONCLUSIONS
The main conclusions based on the presented results
are as follows:
• Sub-zero treatment reduces the amount of retained
austenite and increases the population density of
small globular carbides in the microstructure.
• The amount of small globular carbides decreases
with the application of tempering treatment. The
higher the tempering temperature, the more signifi-
cant is the reduction of the carbide count.
• The hardness of the sub-zero treated material is
higher than that of the conventionally quenched one.
Also, this tendency is preserved when steel is low-
temperature tempered.
• On the other hand, the hardness of the conventionally
quenched steel becomes higher than that of the SZT
one when tempered at the temperature of secondary
hardening.
• The fracture toughness is generally higher for the
conventionally quenched steel Vanadis 6 except in
the case when the material is tempered to the se-
condary hardness. One can conclude that KIC follows
the reciprocal value of the hardness, i.e., the harder
the material the lower is the fracture toughness.
• The differences in KIC are reflected in the topo-
graphies of the fracture surfaces, represented by the
surface roughness – the higher the fracture toughness,
the higher is the roughness of the fracture surface.
Acknowledgements
This paper is the result of implementing project CE
for development and application of advanced diagnostic
methods in processing of metallic and non-metallic
materials, ITMS: 26220120048, supported by the Re-
search & Development Operational Programme funded
by the ERDF. This research was supported by the grant
project VEGA 1/0735/14.
5 REFERENCES
1 P. Jur~i, Cr-V Ledeburitic Cold-Work Tool Steels, Mater. Tehnol., 45
(2011) 5, 383–394
2 D. N. Collins, Cryogenic treatment of tool steels, Advanced Ma-
terials and Processes, 12 (1998), 24–29
3 D. N. Collins, J. Dormer, Deep Cryogenic Treatment of a D2 Cold-
Work Tool Steel, Heat Treatment of Metals, 24 (1997) 3, 71–74
4 H. Berns, C. Bröckmann, Fracture of Hot Formed Ledeburitic Chro-
mium Steels, Engineering Fracture Mechanics, 58 (1997), 311–325,
doi:10.1016/S0013-7944(97)00118-5
5 I. Dlouhý, M. Holzmann, J. Valka, Using chevron notched specimens
for determining fracture toughness of bearing steels, Kovove mate-
riály/Metallic materials, 32 (1994), 3–13
6 D. Das, R. Sarkar, A. K. Dutta, K. K. Ray, Influence of sub-zero
treatments on fracture toughness of AISI D2 steel, Materials Science
and Engineering, A 528 (2010), 589–603, doi:10.1016/j.msea.2010.
09.057
7 J. Sobotová, P. Jur~i, I. Dlouhý, The effect of subzero treatment on
microstructure, fracture toughness and wear resistance of Vanadis 6
tool steel, Materials Science and Engineering, A 652 (2016),
192–204, doi:10.1016/j.msea.2015.11.078
8 K. Amini, A. Akhbarizadeh, S. Javadpour, Investigating the effect of
holding duration on the microstructure of 1.2080 tool steel during the
deep cryogenic treatment, Vacuum, 86 (2012), 1534–1540,
doi:10.1016/j.vacuum.2012.02.013
9 D. Das, A. K. Dutta, K. K. Ray, Sub-zero treatments of AISI D2
steel: Part I, Microstructure and hardness, Materials Science and
Engineering, A 527 (2010), 2182–2193, doi:10.1016/j.msea.2009.
10.070
10 P. Jur~i, M. Dománková, ¼. ^aplovi~, J. Pta~inová, J. Sobotová, P.
Salabová, O. Prikner, B. [u{tar{i~, D. Jenko, Microstructure and
hardness of sub-zero treated and untempered P/M Vanadis 6 ledebu-
ritic tool steel, Vacuum, 111 (2015), 92–101, doi:10.1016/j.vacuum.
2014.10.004
11 P. Bílek, J. Sobotová, P. Jur~i, Evaluation of the Microstructural
Changes in Cr-V Ledeburitic Tool Steel Depending on the Austeni-
tization Temperature, Mater. Tehnol., 45 (2011), 489–493
12 ^SN EN ISO 12137: 2010 – Metallic materials – determination of
plane strain fracture toughness, European Committee for Standardi-
zation, Brussels
J. PTA^INOVÁ et al.: FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL AFTER SUB-ZERO TREATMENT ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 729–733 733
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
the point of view of industrial practice. They indicate
that it is possible to increase the wear performance,
along the toughness, of the material in a certain tem-
pering-temperature range.
3.4 Fracture-surface morphology
The micro-mechanics of fracture propagation is
demonstrated through representative SEM micrographs
showing how the fractures appear in the cases of the steel
that was sub-zero treated at –196 °C for 17 h and not
tempered (the lowest KIC) and the steel that was sub-zero
treated at –196 °C for 17 h and tempered at 170 °C
(higher KIC). The fracture surface of the sample that was
not tempered and the one that was tempered at 170 °C
are shown in Figure 4.
At first glance, there is not any difference between
the mentioned fracture surfaces; they appear flat and
shiny, Figures 4a and 4c. However, a more detailed
observation, at higher magnifications, makes it clear that
the fracture surface of the sample after low-temperature
tempering contains an enhanced number of sites with
micro-plastic deformation (PLD), located mainly at the
carbide/matrix interfaces, Figure 4b. On the other hand,
there are very few such sites on the fracture surface of
the untempered specimen, Figure 4d, and the fracture
surface contains a number of cleavage facettes/regions
(CL).
The topographies of the fracture surfaces of the CHT
Vanadis 6 steel tempered at 170 °C and sub-zero treated,
and the untempered one, expressed with surface rough-
ness Ra are shown in Table 2 and Figure 5. The average
roughness of the fracture surface of the conventionally
quenched steel tempered at 170 °C was 4.2703±0.248 mm.
J. PTA^INOVÁ et al.: FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL AFTER SUB-ZERO TREATMENT ...
732 Materiali in tehnologije / Materials and technology 51 (2017) 5, 729–733
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Representative SEM micrographs of the fracture surfaces of
the specimens processed with: a), b) SZT at –196 °C for 17 h and tem-
pered at 170 °C and c), d) SZT at –196 °C for 17 h and not tempered
Figure 3: Hardness HV10 and fracture toughness KIC of Vanadis 6
steel in dependence on the heat treatment applied: a) conventional
quenching, b) sub-zero treatment
Figure 5: Confocal-microscope micrographs showing the fracture surfaces of KIC specimens: a) CHT and tempered at 170 °C, b) SZT and not
tempered
The roughness of the sub-zero treated and untempered
Vanadis 6 steel was lower, i.e., 2.918±0.323 μm. At this
point, it should be noted that in this stage of
experimental efforts, only the materials with the highest
KIC and the one with the lowest KIC were checked with
confocal microscopy. However, the measurements gave
unambiguous differences in the roughness, where the
higher roughness corresponded to the higher KIC, i.e., to
the conventionally quenched and low-temperature tem-
pered sample.
Table 2: Roughness of the fracture surface of Vanadis 6 ledeburitic
steel
No. of measurement Tempered at 170 °C SZT anduntempered
1. 4.136 2.591
2. 4.750 3.507
3. 3.925 2.656
Ra (μm) 4.2703±0.248 2.918±0.323
4 CONCLUSIONS
The main conclusions based on the presented results
are as follows:
• Sub-zero treatment reduces the amount of retained
austenite and increases the population density of
small globular carbides in the microstructure.
• The amount of small globular carbides decreases
with the application of tempering treatment. The
higher the tempering temperature, the more signifi-
cant is the reduction of the carbide count.
• The hardness of the sub-zero treated material is
higher than that of the conventionally quenched one.
Also, this tendency is preserved when steel is low-
temperature tempered.
• On the other hand, the hardness of the conventionally
quenched steel becomes higher than that of the SZT
one when tempered at the temperature of secondary
hardening.
• The fracture toughness is generally higher for the
conventionally quenched steel Vanadis 6 except in
the case when the material is tempered to the se-
condary hardness. One can conclude that KIC follows
the reciprocal value of the hardness, i.e., the harder
the material the lower is the fracture toughness.
• The differences in KIC are reflected in the topo-
graphies of the fracture surfaces, represented by the
surface roughness – the higher the fracture toughness,
the higher is the roughness of the fracture surface.
Acknowledgements
This paper is the result of implementing project CE
for development and application of advanced diagnostic
methods in processing of metallic and non-metallic
materials, ITMS: 26220120048, supported by the Re-
search & Development Operational Programme funded
by the ERDF. This research was supported by the grant
project VEGA 1/0735/14.
5 REFERENCES
1 P. Jur~i, Cr-V Ledeburitic Cold-Work Tool Steels, Mater. Tehnol., 45
(2011) 5, 383–394
2 D. N. Collins, Cryogenic treatment of tool steels, Advanced Ma-
terials and Processes, 12 (1998), 24–29
3 D. N. Collins, J. Dormer, Deep Cryogenic Treatment of a D2 Cold-
Work Tool Steel, Heat Treatment of Metals, 24 (1997) 3, 71–74
4 H. Berns, C. Bröckmann, Fracture of Hot Formed Ledeburitic Chro-
mium Steels, Engineering Fracture Mechanics, 58 (1997), 311–325,
doi:10.1016/S0013-7944(97)00118-5
5 I. Dlouhý, M. Holzmann, J. Valka, Using chevron notched specimens
for determining fracture toughness of bearing steels, Kovove mate-
riály/Metallic materials, 32 (1994), 3–13
6 D. Das, R. Sarkar, A. K. Dutta, K. K. Ray, Influence of sub-zero
treatments on fracture toughness of AISI D2 steel, Materials Science
and Engineering, A 528 (2010), 589–603, doi:10.1016/j.msea.2010.
09.057
7 J. Sobotová, P. Jur~i, I. Dlouhý, The effect of subzero treatment on
microstructure, fracture toughness and wear resistance of Vanadis 6
tool steel, Materials Science and Engineering, A 652 (2016),
192–204, doi:10.1016/j.msea.2015.11.078
8 K. Amini, A. Akhbarizadeh, S. Javadpour, Investigating the effect of
holding duration on the microstructure of 1.2080 tool steel during the
deep cryogenic treatment, Vacuum, 86 (2012), 1534–1540,
doi:10.1016/j.vacuum.2012.02.013
9 D. Das, A. K. Dutta, K. K. Ray, Sub-zero treatments of AISI D2
steel: Part I, Microstructure and hardness, Materials Science and
Engineering, A 527 (2010), 2182–2193, doi:10.1016/j.msea.2009.
10.070
10 P. Jur~i, M. Dománková, ¼. ^aplovi~, J. Pta~inová, J. Sobotová, P.
Salabová, O. Prikner, B. [u{tar{i~, D. Jenko, Microstructure and
hardness of sub-zero treated and untempered P/M Vanadis 6 ledebu-
ritic tool steel, Vacuum, 111 (2015), 92–101, doi:10.1016/j.vacuum.
2014.10.004
11 P. Bílek, J. Sobotová, P. Jur~i, Evaluation of the Microstructural
Changes in Cr-V Ledeburitic Tool Steel Depending on the Austeni-
tization Temperature, Mater. Tehnol., 45 (2011), 489–493
12 ^SN EN ISO 12137: 2010 – Metallic materials – determination of
plane strain fracture toughness, European Committee for Standardi-
zation, Brussels
J. PTA^INOVÁ et al.: FRACTURE TOUGHNESS OF LEDEBURITIC VANADIS 6 STEEL AFTER SUB-ZERO TREATMENT ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 729–733 733
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
735–743
ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL
OXIDES MgxZn1-xAl2O4 BY FIRST-PRINCIPLES CALCULATIONS
ELEKTRONSKE IN OPTI^NE LASTNOSTI SPINELNIH OKSIDOV
MgxZn1-xAl2O4, IZPELJANE IZ TEORETI^NIH OSNOV
Chao Xiang1, Jianxiong Zhang1, Yun Lu2, Dong Tian3, Cheng Peng1
1Yangtze Normai University, School of Mechanical and Engineering, Fuling 408000, China
2Chiba University, Institute of Material Science and Engineering, Chiba 2790000, Japan
3Kunming University of Science and Technology, Faculty of Science, Kunming 650093, China
1254618608@qq.com
Prejem rokopisa – received: 2016-10-07; sprejem za objavo – accepted for publication: 2017-02-10
doi:10.17222/mit.2016.296
The structural, electronic and optical properties of perfect MgxZn1-xAl2O4 oxides have been studied by first-principles
calculations within the generalized gradient approximation of the density functional theory. It is interesting to note that a linear
increase of cell volume (V) with increasing doping amount (x) occurs. The band gap increases in the series from 3.851 eV to
5.079 eV, which is in agreement with theoretical and experimental values. In addition, a blue shift of the absorption shoulder is
observed in the UV region with the increase of x, as predicted by the imaginary part 2() of the dielectric function at zero
frequency as well as bandgap. This can be explained by the threshold of the electronic transition from O-2p to the empty Mg-3p
electron states due to the substitution of Zn with Mg. The real part 1() of the dielectric function located at zero frequency has
a square fit relationship with refractive index n(0), which is 1.71–1.77 from x=0 to x=1. The energy-loss function shows that the
replacement of Zn by Mg is responsible for a decrease in the intensity of the sharp peaks. The reflectivity shows that a higher
coefficient of reflectivity (R(0)) at zero frequency corresponds to a smaller bandgap.
Keywords: electronic transitions, dielectric function, refractive index, adsorption shoulder
Preiskovali smo strukturne, elektronske in opti~ne lastnosti idealnih MgxZn1-xAl2O4 oksidov, izpeljane iz teoreti~nih osnov
znotraj posplo{ene gradientne aproksimacije funkcionalne teorije gostote. Opazili smo linearno pove~anje celi~nega volumna
(V) z nara{~ajo~o koncentracijo Mg (x). [irina prepovedanega pasu nara{~a od 3.851 eV to 5.079 eV v skladu s teoreti~nimi in
eksperimentalnimi vrednostmi. Poleg tega ob nara{~anju dele`a Mg opazimo modri premik dodatnega absorpcijskega vrha v UV
obmo~ju, kakor napovedujeta vrednosti imaginarnega dela 2() dielektri~ne funkcije pri frekvenci 0 in prepovedanega pasu. To
je mogo~e pojasniti s pragom elektronskega prehoda elektronov iz O-2p v prazen Mg-3p zaradi nadomestitve Zn z Mg. Kvadrat
realnega dela 1() dielektri~ne funkcije pri frekvenci 0 se ujema z lomnim koli~nikom n(0), ki je 1.71–1.77 za x=0 do x=1.
Funkcija izgube energije, ka`e, da zamenjava Zn z Mg povzro~a zmanj{anje intenzitete ostrih vrhov. Reflektivnost ka`e, da vi{ji
koeficient refleksije (R(0)) pri frekvenci 0 odgovarja manj{i {irini prepovedanega pasu.
Klju~ne besede: prehodi elektronov, dielektri~na funkcija, lomni koli~nik, dodatni absorpcijski vrh
1 INTRODUCTION
Spinal oxides with the general chemical formula
AB2O4 have a close-packed, face-centered-cubic struc-
ture (space group Fd3m) characterized by two
symmetrically distinct polyhedra: a tetrahedron and an
octahedron. They are widely used in various fields such
as catalysis, gas sensor, semiconductor, biomedical,
catalyst carrier, as well as electroluminescent displays
owing to their catalytic, physical, structural, electronic
and optical properties.1,2 Among them, MgAl2O4 and
ZnAl2O4 have high-temperature resistance,3,4 and they
are highly reflective for wavelengths in the ultraviolet
(UV) region, which make them candidate materials for
reflective optical coating in aerospace applications.5 In
particular, MgAl2O4 is one of the potential candidates for
the full wave band transparent window materials with
high transmittance in IR- and visible-wavelength even
extending to microwave ranges,6,7 and it also can be used
as lamps and lasers,8 transparent ceramic material for
high-temperature,9 transparent armor and glass.10
Similarly, ZnAl2O4 can be used as a ceramic material
similar to MgAl2O4. It can be used as a transparent
conductor, optical material and dielectric material,11,12
and it is suitable for UV photoelectronic applications.13
Simultaneously, much work has been done on the
structural, electronic and optical properties of MgAl2O4
and ZnAl2O4 over the past few years.14–27 The effect of
point vacancies on the spectral properties of MgAl2O4
has been studied by S. L. Jiang et al.14 They revealed that
the absorption peak at 5.3 eV is attributed to the neutral
oxygen vacancy Vo
0 , while two peaks at 3.2 eV and 4.75
eV are attributed to the 1+ charged oxygen vacancy Vo
1 + .
A related mechanism of transparency in MgAl2O4 nano-
ceramics prepared by sintering under high pressure and
low temperature has been studied by J. Zhang et al.15,
who suggested that the decrease in the transparency with
increasing temperature (>700 °C) is therefore a result of
the light scattering at large pores. The low-temperature,
high-pressure preparation of transparent nanocrystalline
MgAl2O4 ceramics has been investigated by T. C. Lu et
Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743 735
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:621.3.011.5:535.327 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)735(2017)
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
735–743
ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL
OXIDES MgxZn1-xAl2O4 BY FIRST-PRINCIPLES CALCULATIONS
ELEKTRONSKE IN OPTI^NE LASTNOSTI SPINELNIH OKSIDOV
MgxZn1-xAl2O4, IZPELJANE IZ TEORETI^NIH OSNOV
Chao Xiang1, Jianxiong Zhang1, Yun Lu2, Dong Tian3, Cheng Peng1
1Yangtze Normai University, School of Mechanical and Engineering, Fuling 408000, China
2Chiba University, Institute of Material Science and Engineering, Chiba 2790000, Japan
3Kunming University of Science and Technology, Faculty of Science, Kunming 650093, China
1254618608@qq.com
Prejem rokopisa – received: 2016-10-07; sprejem za objavo – accepted for publication: 2017-02-10
doi:10.17222/mit.2016.296
The structural, electronic and optical properties of perfect MgxZn1-xAl2O4 oxides have been studied by first-principles
calculations within the generalized gradient approximation of the density functional theory. It is interesting to note that a linear
increase of cell volume (V) with increasing doping amount (x) occurs. The band gap increases in the series from 3.851 eV to
5.079 eV, which is in agreement with theoretical and experimental values. In addition, a blue shift of the absorption shoulder is
observed in the UV region with the increase of x, as predicted by the imaginary part 2() of the dielectric function at zero
frequency as well as bandgap. This can be explained by the threshold of the electronic transition from O-2p to the empty Mg-3p
electron states due to the substitution of Zn with Mg. The real part 1() of the dielectric function located at zero frequency has
a square fit relationship with refractive index n(0), which is 1.71–1.77 from x=0 to x=1. The energy-loss function shows that the
replacement of Zn by Mg is responsible for a decrease in the intensity of the sharp peaks. The reflectivity shows that a higher
coefficient of reflectivity (R(0)) at zero frequency corresponds to a smaller bandgap.
Keywords: electronic transitions, dielectric function, refractive index, adsorption shoulder
Preiskovali smo strukturne, elektronske in opti~ne lastnosti idealnih MgxZn1-xAl2O4 oksidov, izpeljane iz teoreti~nih osnov
znotraj posplo{ene gradientne aproksimacije funkcionalne teorije gostote. Opazili smo linearno pove~anje celi~nega volumna
(V) z nara{~ajo~o koncentracijo Mg (x). [irina prepovedanega pasu nara{~a od 3.851 eV to 5.079 eV v skladu s teoreti~nimi in
eksperimentalnimi vrednostmi. Poleg tega ob nara{~anju dele`a Mg opazimo modri premik dodatnega absorpcijskega vrha v UV
obmo~ju, kakor napovedujeta vrednosti imaginarnega dela 2() dielektri~ne funkcije pri frekvenci 0 in prepovedanega pasu. To
je mogo~e pojasniti s pragom elektronskega prehoda elektronov iz O-2p v prazen Mg-3p zaradi nadomestitve Zn z Mg. Kvadrat
realnega dela 1() dielektri~ne funkcije pri frekvenci 0 se ujema z lomnim koli~nikom n(0), ki je 1.71–1.77 za x=0 do x=1.
Funkcija izgube energije, ka`e, da zamenjava Zn z Mg povzro~a zmanj{anje intenzitete ostrih vrhov. Reflektivnost ka`e, da vi{ji
koeficient refleksije (R(0)) pri frekvenci 0 odgovarja manj{i {irini prepovedanega pasu.
Klju~ne besede: prehodi elektronov, dielektri~na funkcija, lomni koli~nik, dodatni absorpcijski vrh
1 INTRODUCTION
Spinal oxides with the general chemical formula
AB2O4 have a close-packed, face-centered-cubic struc-
ture (space group Fd3m) characterized by two
symmetrically distinct polyhedra: a tetrahedron and an
octahedron. They are widely used in various fields such
as catalysis, gas sensor, semiconductor, biomedical,
catalyst carrier, as well as electroluminescent displays
owing to their catalytic, physical, structural, electronic
and optical properties.1,2 Among them, MgAl2O4 and
ZnAl2O4 have high-temperature resistance,3,4 and they
are highly reflective for wavelengths in the ultraviolet
(UV) region, which make them candidate materials for
reflective optical coating in aerospace applications.5 In
particular, MgAl2O4 is one of the potential candidates for
the full wave band transparent window materials with
high transmittance in IR- and visible-wavelength even
extending to microwave ranges,6,7 and it also can be used
as lamps and lasers,8 transparent ceramic material for
high-temperature,9 transparent armor and glass.10
Similarly, ZnAl2O4 can be used as a ceramic material
similar to MgAl2O4. It can be used as a transparent
conductor, optical material and dielectric material,11,12
and it is suitable for UV photoelectronic applications.13
Simultaneously, much work has been done on the
structural, electronic and optical properties of MgAl2O4
and ZnAl2O4 over the past few years.14–27 The effect of
point vacancies on the spectral properties of MgAl2O4
has been studied by S. L. Jiang et al.14 They revealed that
the absorption peak at 5.3 eV is attributed to the neutral
oxygen vacancy Vo
0 , while two peaks at 3.2 eV and 4.75
eV are attributed to the 1+ charged oxygen vacancy Vo
1 + .
A related mechanism of transparency in MgAl2O4 nano-
ceramics prepared by sintering under high pressure and
low temperature has been studied by J. Zhang et al.15,
who suggested that the decrease in the transparency with
increasing temperature (>700 °C) is therefore a result of
the light scattering at large pores. The low-temperature,
high-pressure preparation of transparent nanocrystalline
MgAl2O4 ceramics has been investigated by T. C. Lu et
Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743 735
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:621.3.011.5:535.327 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)735(2017)
al.,16 indicating that the nanoceramics are highly
transparent even though their relative densities are all
less than 99 %, owing to the low or negligible light
scattering from the nanosized grains and pores. The
optical properties of ZnAl2O4 nanomaterials obtained by
the hydrothermal method have been investigated by
Miron and co-workers,17 demonstrating that the band gap
is determined from the absorbance spectra, and it de-
pends strongly on the temperature used for further
heating the samples. A first-principles study on struc-
tural, electronic and optical properties of spinel oxides
ZnAl2O4, ZnGa2O4 and ZnIn2O4 has been carried out by
F. Zerarga and co-workers,18 implying that the peaks and
structures in the optical spectra are assigned to interband
transitions. The fabrication of transparent polycrystalline
ZnAl2O4 – a new optical bulk ceramic – has been inve-
stigated by Goldstein and co-workers19, who suggested
that specimens have a high transparency (ILT78\%;
 = 800 nm; t = 2 mm). Plus, the differences in struc-
tural, electronic and optical performance between alumi-
num spinel MgAl2O4 and ZnAl2O4 have been
presented.28–30 To the best of our knowledge, the struc-
ture, electronic and optical properties of MgxZn1-xAl2O4
(0<x<1), expressed quantitatively, which is that Zn2+ of
ZnAl2O4 is replaced by Mg2+, has not been investigated.
The main aim of this investigation is to study the effect
of the replacement of Zn by Mg on the structural,
electronic and optical properties of ZnAl2O4.
2 MODEL AND COMPUTATIONAL DETAILS
In our calculations a 56-atom unit cell was modeled
for the investigation of spinel MgxZn1-xAl2O4 in Figure 1.
All the calculations were performed using the Cam-
bridge Serial Total Energy Package (CASTEP) prog-
ram,31 based on density functional theory (DFT).
Electron-ion interaction for cations and anions was
described by the ultra-soft pseudo-potential. The Mg-3s
and -3p, Zn-3d and -4s, Al-3s and -3p, as well as O-2s
and -2p states were treated as valence electrons. The
exchange-correlation potential was calculated by the
generalized gradient approximation (GGA) developed by
the Perdew, Burke, and Ernzerhof functional (PBE).32
The Engel-Vosko scheme (GGA+U) was applied to
electronic properties calculations.33 The on-site Coulomb
Hubbard interaction U = 1 eV is used to treat correctly
the localized 3d electrons in the MgxZn1-xAl2O4. The unit
cell was optimized using the Broyden, Fletcher, Gold-
farb, Shanno (BFGS) method,34 and convergence tole-
rance of energy charge, maximum force and stress were
1×10–5 eV, 0.03 eV/nm, and 0.05 GPa, respectively.18,35
The plane-wave basis set cutoff energy was set to
340 eV.18,35,36 The Brillouin-zone integration was imple-
mented within Monkhorst-Pack scheme using a 4×4×4
mesh to optimize structures and calculate the electronic
and optical properties.
3 RESULTS AND DISCUSSION
3.1 Structure and electronic properties
To obtain the stable structure of MgxZn1-xAl2O4
spinel, the internal coordinates and lattice parameters of
crystals were relaxed during geometry optimization.
Structure parameters such as average bond lengths, cell
volumes and lattice constants are presented, in compari-
son with the experimental data and values available from
other calculations. As can be seen from Tables 1 and 2,
the calculated equilibrium lattice constants a0 of both
ZnAl2O4 and MgAl2O4 are larger than the experimental
values with less than 2 % deviation. Plus, substitution of
Zn with Mg causes a slight increase in the lattice con-
stant of ZnAl2O4. This is in agreement with the previous
report.28 In addition, one can see that the substitution of
Zn with Mg is responsible for the formation of the
pseudo-cubic, spinel-type structure. The average bond
lengths of dAl-O increases with increasing doping con-
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
736 Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Calculated average bond length (nm) of MgxZn1-xAl2O4
Figure 1: MgxZn1-xAl2O4 crystal structure. The gray, pink, green and
red represent Zn, Al, Mg and O atoms, respectively.
centration (from x=0 to x=1) in Figure 2. In contrast,
both dZn-O and dMg-O are nearly unchanged. This suggests
that the increase of lattice constants is mostly due to an
increase of dAl-O. The minimum dAl-O in MgAl2O4 is
0.1938 nm, which indicates that the strength of the
chemical bond (Al-O) of MgAl2O4 is stronger than that
of MgxZn1-xAl2O4. The cell volume (V) is listed in Table
2. The volume-x curve is plotted, where x is the concen-
tration of Mg. It is interesting to note that the unit-cell
volume (V) changes linearly with the doping amount x
(Figure 3).
Table 1: Calculated structure parameters for MgxZn1-xAl2O4 (x=0,
x=1), experimental measurement and other available theoretical values
ZnAl2O4 MgAl2O4
Present Expt. Calc. Present Expt. Calc.
a0
(nm) 0.8194
0.8091137 0.803218
0.8208
0.8083643 0.802728
0.808638
0.799528 0.8084444 0.8107425
0.819339
0.80645
0.814627
m
0.8021540 0.8093414
0.801841
0.8246
0.80242
The calculated band gaps are presented in Figure 4.
From the figure it is easy to see that the bottom of the
conduction band and the top of the valence band for
MgxZn1-xAl2O4 occurs at  point with a direct band gap.
Among them, ZnAl2O4 (x=0) has a band gap value of
3.851 eV, which is very close to the experimental value.48
The MgAl2O4 (x=1) has a band gap value of 5.079 eV,
which is underestimated with respect to the experimental
value of 7.8 eV.49 A similar issue has been addressed in
previous theoretical calculations for MgAl2O4 such as
5.30 eV14, 5.1 eV27 and 5.36 eV.28 This is mostly due to
the fact that the calculated band gaps are related to DFT
limitations, not considering the discontinuity in the
exchange-correlation potential,50 as mentioned in 35. The
calculations also show that the band-gap value increases
due to changing the cation Zn with Mg. The band gap is
4.078 for x=0.25, 4.349 for x=0.50 and 4.638 eV for
x=0.75.
As discussed above, the band gap of optimized
MgxZn1-xAl2O4 is different (from x=0 to x=1). This is
attributed to the difference in the density of states
(DOS). To further elucidate the nature of the electronic
band structure, we calculated the total density of states
(DOS) and partial density of states (PDOS) of O, Zn, Mg
and Al for the spinel oxides. As can be seen from Fig-
ure 5, the Fermi level is set to zero. It is mostly created
by Zn-3d and O-2p states for x=0, with small con-
tributions coming from Al-3p, and it is mostly created by
Mg-3p and O-2p states for x=1, with small contributions
of Al-3p. The character of the upper valence band is
localized between -6.2 eV and 0 eV, which is mainly due
to hybridization between the Zn-3d and O-2p states and
due less to hybridization of O-2p with Zn-3p, Zn-4s,
Mg-3s, Mg-3p, Al-3s and Al-3p as well. The lower
valence band in the energy range -18.8 eV to -16.3 eV is
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743 737
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Calculated the unit cell volume V (nm3) as a function of
composition x
Table 2: Calculated structure parameters for MgxZn1-xAl2O4
0.000 0.250 0.500 0.750 1000 Al2O3
a(nm) 0.8194 0.8198 0.8202 0.8205 0.8208
b(nm) 0.8194 0.8197 0.8202 0.8205 0.8208
c(nm) 0.8194 0.8197 0.8199 0.8203 0.8208
dZn-O(nm) 0.1980 0.201918 0.1979 0.1978 0.1979 0.203218
dAl-O(nm) 0.1938 0.196318 0.188b 0.1939 0.1940 0.1942 0.1943 0.1941925 0.19247
dMg-O(nm) 0.1980 0.1980 0.1978 0.1978 0.1956925
Volume (nm3) 550.25 0.55092 0.55164 0.55237 0.55312
Figure 4: Calculated electronic band structure of MgxZn1-xAl2O4
(x=0, x=0.25, x=0.5, x=0.75, x=1)
al.,16 indicating that the nanoceramics are highly
transparent even though their relative densities are all
less than 99 %, owing to the low or negligible light
scattering from the nanosized grains and pores. The
optical properties of ZnAl2O4 nanomaterials obtained by
the hydrothermal method have been investigated by
Miron and co-workers,17 demonstrating that the band gap
is determined from the absorbance spectra, and it de-
pends strongly on the temperature used for further
heating the samples. A first-principles study on struc-
tural, electronic and optical properties of spinel oxides
ZnAl2O4, ZnGa2O4 and ZnIn2O4 has been carried out by
F. Zerarga and co-workers,18 implying that the peaks and
structures in the optical spectra are assigned to interband
transitions. The fabrication of transparent polycrystalline
ZnAl2O4 – a new optical bulk ceramic – has been inve-
stigated by Goldstein and co-workers19, who suggested
that specimens have a high transparency (ILT78\%;
 = 800 nm; t = 2 mm). Plus, the differences in struc-
tural, electronic and optical performance between alumi-
num spinel MgAl2O4 and ZnAl2O4 have been
presented.28–30 To the best of our knowledge, the struc-
ture, electronic and optical properties of MgxZn1-xAl2O4
(0<x<1), expressed quantitatively, which is that Zn2+ of
ZnAl2O4 is replaced by Mg2+, has not been investigated.
The main aim of this investigation is to study the effect
of the replacement of Zn by Mg on the structural,
electronic and optical properties of ZnAl2O4.
2 MODEL AND COMPUTATIONAL DETAILS
In our calculations a 56-atom unit cell was modeled
for the investigation of spinel MgxZn1-xAl2O4 in Figure 1.
All the calculations were performed using the Cam-
bridge Serial Total Energy Package (CASTEP) prog-
ram,31 based on density functional theory (DFT).
Electron-ion interaction for cations and anions was
described by the ultra-soft pseudo-potential. The Mg-3s
and -3p, Zn-3d and -4s, Al-3s and -3p, as well as O-2s
and -2p states were treated as valence electrons. The
exchange-correlation potential was calculated by the
generalized gradient approximation (GGA) developed by
the Perdew, Burke, and Ernzerhof functional (PBE).32
The Engel-Vosko scheme (GGA+U) was applied to
electronic properties calculations.33 The on-site Coulomb
Hubbard interaction U = 1 eV is used to treat correctly
the localized 3d electrons in the MgxZn1-xAl2O4. The unit
cell was optimized using the Broyden, Fletcher, Gold-
farb, Shanno (BFGS) method,34 and convergence tole-
rance of energy charge, maximum force and stress were
1×10–5 eV, 0.03 eV/nm, and 0.05 GPa, respectively.18,35
The plane-wave basis set cutoff energy was set to
340 eV.18,35,36 The Brillouin-zone integration was imple-
mented within Monkhorst-Pack scheme using a 4×4×4
mesh to optimize structures and calculate the electronic
and optical properties.
3 RESULTS AND DISCUSSION
3.1 Structure and electronic properties
To obtain the stable structure of MgxZn1-xAl2O4
spinel, the internal coordinates and lattice parameters of
crystals were relaxed during geometry optimization.
Structure parameters such as average bond lengths, cell
volumes and lattice constants are presented, in compari-
son with the experimental data and values available from
other calculations. As can be seen from Tables 1 and 2,
the calculated equilibrium lattice constants a0 of both
ZnAl2O4 and MgAl2O4 are larger than the experimental
values with less than 2 % deviation. Plus, substitution of
Zn with Mg causes a slight increase in the lattice con-
stant of ZnAl2O4. This is in agreement with the previous
report.28 In addition, one can see that the substitution of
Zn with Mg is responsible for the formation of the
pseudo-cubic, spinel-type structure. The average bond
lengths of dAl-O increases with increasing doping con-
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
736 Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Calculated average bond length (nm) of MgxZn1-xAl2O4
Figure 1: MgxZn1-xAl2O4 crystal structure. The gray, pink, green and
red represent Zn, Al, Mg and O atoms, respectively.
centration (from x=0 to x=1) in Figure 2. In contrast,
both dZn-O and dMg-O are nearly unchanged. This suggests
that the increase of lattice constants is mostly due to an
increase of dAl-O. The minimum dAl-O in MgAl2O4 is
0.1938 nm, which indicates that the strength of the
chemical bond (Al-O) of MgAl2O4 is stronger than that
of MgxZn1-xAl2O4. The cell volume (V) is listed in Table
2. The volume-x curve is plotted, where x is the concen-
tration of Mg. It is interesting to note that the unit-cell
volume (V) changes linearly with the doping amount x
(Figure 3).
Table 1: Calculated structure parameters for MgxZn1-xAl2O4 (x=0,
x=1), experimental measurement and other available theoretical values
ZnAl2O4 MgAl2O4
Present Expt. Calc. Present Expt. Calc.
a0
(nm) 0.8194
0.8091137 0.803218
0.8208
0.8083643 0.802728
0.808638
0.799528 0.8084444 0.8107425
0.819339
0.80645
0.814627
m
0.8021540 0.8093414
0.801841
0.8246
0.80242
The calculated band gaps are presented in Figure 4.
From the figure it is easy to see that the bottom of the
conduction band and the top of the valence band for
MgxZn1-xAl2O4 occurs at  point with a direct band gap.
Among them, ZnAl2O4 (x=0) has a band gap value of
3.851 eV, which is very close to the experimental value.48
The MgAl2O4 (x=1) has a band gap value of 5.079 eV,
which is underestimated with respect to the experimental
value of 7.8 eV.49 A similar issue has been addressed in
previous theoretical calculations for MgAl2O4 such as
5.30 eV14, 5.1 eV27 and 5.36 eV.28 This is mostly due to
the fact that the calculated band gaps are related to DFT
limitations, not considering the discontinuity in the
exchange-correlation potential,50 as mentioned in 35. The
calculations also show that the band-gap value increases
due to changing the cation Zn with Mg. The band gap is
4.078 for x=0.25, 4.349 for x=0.50 and 4.638 eV for
x=0.75.
As discussed above, the band gap of optimized
MgxZn1-xAl2O4 is different (from x=0 to x=1). This is
attributed to the difference in the density of states
(DOS). To further elucidate the nature of the electronic
band structure, we calculated the total density of states
(DOS) and partial density of states (PDOS) of O, Zn, Mg
and Al for the spinel oxides. As can be seen from Fig-
ure 5, the Fermi level is set to zero. It is mostly created
by Zn-3d and O-2p states for x=0, with small con-
tributions coming from Al-3p, and it is mostly created by
Mg-3p and O-2p states for x=1, with small contributions
of Al-3p. The character of the upper valence band is
localized between -6.2 eV and 0 eV, which is mainly due
to hybridization between the Zn-3d and O-2p states and
due less to hybridization of O-2p with Zn-3p, Zn-4s,
Mg-3s, Mg-3p, Al-3s and Al-3p as well. The lower
valence band in the energy range -18.8 eV to -16.3 eV is
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743 737
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Calculated the unit cell volume V (nm3) as a function of
composition x
Table 2: Calculated structure parameters for MgxZn1-xAl2O4
0.000 0.250 0.500 0.750 1000 Al2O3
a(nm) 0.8194 0.8198 0.8202 0.8205 0.8208
b(nm) 0.8194 0.8197 0.8202 0.8205 0.8208
c(nm) 0.8194 0.8197 0.8199 0.8203 0.8208
dZn-O(nm) 0.1980 0.201918 0.1979 0.1978 0.1979 0.203218
dAl-O(nm) 0.1938 0.196318 0.188b 0.1939 0.1940 0.1942 0.1943 0.1941925 0.19247
dMg-O(nm) 0.1980 0.1980 0.1978 0.1978 0.1956925
Volume (nm3) 550.25 0.55092 0.55164 0.55237 0.55312
Figure 4: Calculated electronic band structure of MgxZn1-xAl2O4
(x=0, x=0.25, x=0.5, x=0.75, x=1)
mostly attributed to the O 2s states which split into two:
a sharp peak at -17.8 eV and a lower part 2.3 eV wide
with double peaks for x=0. A similar issue appears for
x=1, from which we can see that the lower valence band
located at approximately -15.4 eV – -18.1 eV is mostly
due to the O 2s states, which split into two: a sharp peak
at -16.1 eV and a lower double peak with a width of 2.1
eV. This is in agreement with other theoretical calcula-
tions.14,25 Moreover, we note that Zn-4s and Al-3p states
play the main role in the conduction band above the
Fermi level (between 3.8 eV and 8.2 eV) for x=0, while
conduction band is mainly created by Mg-3p states for
x=1.
3.2 Optical properties
The common response from both electronic and ionic
polarization are taken into account in the condition of a
low-frequency electric field, although only electronic
polarization is considered as the dielectric function, and
phonon contributions to the optical is not taken into
account, when calculating optical properties in CASTEP
package. It is well known that the optical properties are
determined considering only direct interband transitions.
One of the main optical characteristics of a solid is its
dielectric function, which can express other optical pro-
perties and describe the optical response of the medium.
The dielectric function, (), is given in Equation (1):
() = 	() + 
() (1)
Here 	() and 
() are the real and imaginary parts
of the dielectric function, respectively. The complex
2() is defined as Equation (2):39

() =
2 2
0
2e
V
u r E Ek
c
k v c
k
v
k
c
k
vπ

  
 ⋅ − −∑
, ,
) (2)
where V is the unit cell volume, u is the vector defining
the polarization of the incident electric field, c and v are
the valence band and conduction band, respectively, r is
the momentum operator, and both Ek
c and Ek
v are eigen-
states. This expression is similar to Fermi’s golden rule
for time-dependent perturbations. The 
() can be
thought of as detailing the real transitions between the
unoccupied and occupied electronic states. The ima-
ginary and real parts of the dielectric function are linked
by a Kramers-Kronig transform because the dielectric
function describes a causal response. Furthermore, the
Kramers-Kronig transform is also used to obtain the
	(), which is:18
 
  
 



1 2 2
0
1
2
( )
' ( ' )
( ' ) ( )
'= +
−
∞
∫π P d (3)
Both imaginary parts and the real of the dielectric
function can also be used to calculate important optical
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
738 Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Calculated total (DOS) and partial densities (PDOS) of states for MgxZn1-xAl2O4 (x=0, x=0.25, x=0.5, x=0.75, x=1)
functions, such as the refractive index n() and the
reflectivity R():18
n( )
( ) ( )





   


	






= +
+⎡
⎣
⎢
⎤
⎦
⎥
2
(4)
R( )





=
−
+
⎡
⎣⎢
⎤
⎦⎥
1
1
2
(5)
The peak of the imaginary part of the dielectric
function corresponds to electronic transitions from the
valence to conduction bands, depending on the electronic
transition energy of the conduction and valence band,
i.e., the band-gap energies. The instrumental smearing of
0.3 eV is used to simulate the broadening effects. The
dielectric functions for various x are displayed in Figure
6. It can be seen that the first critical point starts at (3.49,
3.61, 3.90, 4.22 and 4.63) eV from x=0 to x=1. These are
known as the edge of the optical absorption. Further-
more, the different widths of the non-zero part of 
()
are related with different widths of the adsorption
spectra, as shown in Figure 6, suggesting that absorption
region for MgxZn1-xAl2O4(x=0) is wider than that of
MgxZn1-xAl2O4}(x=1). Other peak locates at (7.38, 10.87,
14.63 and 15.54) eV for x=0, (7.49, 7.37, 11.65, 12.83
and 14.22) eV for x=0.25, (6.83, 7.84, 10.31, 12.81 and
15.36) eV for x=0.5, (5.81, 10.16, 12.85 and 15.85) eV
for x=0.75 and (6.25, 7.69, 10.01, 12.91 and 16.16) eV
for x=1, in comparison with other theoretical values
listed in Table 3. These peaks of origins and features on
the basis of decomposing each spectrum to its individual
pair contribution from each pair of valence and conduc-
tion bands are obtained by the analysis of 
().18 The
calculated real part of 	() in infinite wavelength (i.e.,
the limit of zero energy) for x=0 is 3.15 and for x=1 is
2.93, which are in agreement with previous theoretical
values such as 2.6691 for ZnAl2O418 and 3.112 for
MgAl2O4.25 In addition, the 	() in infinite wavelength
is 3.09 for x=0.25, 3.03 for x=0.5, and 2.98 for x=0.75,
signaling a negative trend with the increase of the doping
amount x. These can also indicate that the real part 	()
of the limit of zero energy has a square fit relationship
with the refractive index n localized zero energy. The
	() reaches a maximum value of 4.22 at about 6.15 eV
for x=0, 4.28 at 8.88 eV for x=0.25, 4.30 at 8.82 eV for
x=0.5, 4.29 at 8.84 eV for x=0.75 and 4.39 at 8.73 for
x=1.
Table 3: The imaginary part (2()) of dielectric function for
MgxZn1-xAl2O4 (x=0, x=1).
ZnAl2O4 MgAl2O4
Present Calc.18 Calc.40 Present Calc.23 Calc.25
Threshol
d Peak
position
3.49 4.31 ~4.25 4.63 7.4
7.38 7.82 ~8.01 6.25 6.27
10.87 10.98 7.69 7.44
14.63 14.08 10.01 9.42
15.54 12.91 10.19
The refractive index and the extinction coefficient are
shown in Figure 7. One can see that the static refractive
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743 739
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7: Refractive index of Mg doped ions at various doping con-
centration
Figure 6: Dielectric function of Mg doped ions at different doping
concentration
mostly attributed to the O 2s states which split into two:
a sharp peak at -17.8 eV and a lower part 2.3 eV wide
with double peaks for x=0. A similar issue appears for
x=1, from which we can see that the lower valence band
located at approximately -15.4 eV – -18.1 eV is mostly
due to the O 2s states, which split into two: a sharp peak
at -16.1 eV and a lower double peak with a width of 2.1
eV. This is in agreement with other theoretical calcula-
tions.14,25 Moreover, we note that Zn-4s and Al-3p states
play the main role in the conduction band above the
Fermi level (between 3.8 eV and 8.2 eV) for x=0, while
conduction band is mainly created by Mg-3p states for
x=1.
3.2 Optical properties
The common response from both electronic and ionic
polarization are taken into account in the condition of a
low-frequency electric field, although only electronic
polarization is considered as the dielectric function, and
phonon contributions to the optical is not taken into
account, when calculating optical properties in CASTEP
package. It is well known that the optical properties are
determined considering only direct interband transitions.
One of the main optical characteristics of a solid is its
dielectric function, which can express other optical pro-
perties and describe the optical response of the medium.
The dielectric function, (), is given in Equation (1):
() = 	() + 
() (1)
Here 	() and 
() are the real and imaginary parts
of the dielectric function, respectively. The complex
2() is defined as Equation (2):39

() =
2 2
0
2e
V
u r E Ek
c
k v c
k
v
k
c
k
vπ

  
 ⋅ − −∑
, ,
) (2)
where V is the unit cell volume, u is the vector defining
the polarization of the incident electric field, c and v are
the valence band and conduction band, respectively, r is
the momentum operator, and both Ek
c and Ek
v are eigen-
states. This expression is similar to Fermi’s golden rule
for time-dependent perturbations. The 
() can be
thought of as detailing the real transitions between the
unoccupied and occupied electronic states. The ima-
ginary and real parts of the dielectric function are linked
by a Kramers-Kronig transform because the dielectric
function describes a causal response. Furthermore, the
Kramers-Kronig transform is also used to obtain the
	(), which is:18
 
  
 



1 2 2
0
1
2
( )
' ( ' )
( ' ) ( )
'= +
−
∞
∫π P d (3)
Both imaginary parts and the real of the dielectric
function can also be used to calculate important optical
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
738 Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Calculated total (DOS) and partial densities (PDOS) of states for MgxZn1-xAl2O4 (x=0, x=0.25, x=0.5, x=0.75, x=1)
functions, such as the refractive index n() and the
reflectivity R():18
n( )
( ) ( )





   


	






= +
+⎡
⎣
⎢
⎤
⎦
⎥
2
(4)
R( )





=
−
+
⎡
⎣⎢
⎤
⎦⎥
1
1
2
(5)
The peak of the imaginary part of the dielectric
function corresponds to electronic transitions from the
valence to conduction bands, depending on the electronic
transition energy of the conduction and valence band,
i.e., the band-gap energies. The instrumental smearing of
0.3 eV is used to simulate the broadening effects. The
dielectric functions for various x are displayed in Figure
6. It can be seen that the first critical point starts at (3.49,
3.61, 3.90, 4.22 and 4.63) eV from x=0 to x=1. These are
known as the edge of the optical absorption. Further-
more, the different widths of the non-zero part of 
()
are related with different widths of the adsorption
spectra, as shown in Figure 6, suggesting that absorption
region for MgxZn1-xAl2O4(x=0) is wider than that of
MgxZn1-xAl2O4}(x=1). Other peak locates at (7.38, 10.87,
14.63 and 15.54) eV for x=0, (7.49, 7.37, 11.65, 12.83
and 14.22) eV for x=0.25, (6.83, 7.84, 10.31, 12.81 and
15.36) eV for x=0.5, (5.81, 10.16, 12.85 and 15.85) eV
for x=0.75 and (6.25, 7.69, 10.01, 12.91 and 16.16) eV
for x=1, in comparison with other theoretical values
listed in Table 3. These peaks of origins and features on
the basis of decomposing each spectrum to its individual
pair contribution from each pair of valence and conduc-
tion bands are obtained by the analysis of 
().18 The
calculated real part of 	() in infinite wavelength (i.e.,
the limit of zero energy) for x=0 is 3.15 and for x=1 is
2.93, which are in agreement with previous theoretical
values such as 2.6691 for ZnAl2O418 and 3.112 for
MgAl2O4.25 In addition, the 	() in infinite wavelength
is 3.09 for x=0.25, 3.03 for x=0.5, and 2.98 for x=0.75,
signaling a negative trend with the increase of the doping
amount x. These can also indicate that the real part 	()
of the limit of zero energy has a square fit relationship
with the refractive index n localized zero energy. The
	() reaches a maximum value of 4.22 at about 6.15 eV
for x=0, 4.28 at 8.88 eV for x=0.25, 4.30 at 8.82 eV for
x=0.5, 4.29 at 8.84 eV for x=0.75 and 4.39 at 8.73 for
x=1.
Table 3: The imaginary part (2()) of dielectric function for
MgxZn1-xAl2O4 (x=0, x=1).
ZnAl2O4 MgAl2O4
Present Calc.18 Calc.40 Present Calc.23 Calc.25
Threshol
d Peak
position
3.49 4.31 ~4.25 4.63 7.4
7.38 7.82 ~8.01 6.25 6.27
10.87 10.98 7.69 7.44
14.63 14.08 10.01 9.42
15.54 12.91 10.19
The refractive index and the extinction coefficient are
shown in Figure 7. One can see that the static refractive
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743 739
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7: Refractive index of Mg doped ions at various doping con-
centration
Figure 6: Dielectric function of Mg doped ions at different doping
concentration
index n(0) at the zero frequency limits is 1.77, 1.76,
1.73, 1.72 and 1.71 from x=0 to x=1, which are con-
sistent with values derived from the real part 	() of the
dielectric function mentioned above. They reach a maxi-
mum value of 2.12 at 10.05 eV for x=0, 2.14 at 9.01 eV
for x=0.25, 2.13 at 8.97 eV for x=0.5, 2.14 at 8.95 ev for
x=0.75 and 2.16 at 8.84 eV for x=1, compared to other
theoretical values listed in Table 4. We note that the
refractive index n(0) decreases with an increase of x. It
indicates that the n(0) likely is decided by the bond
length of Zn-O, Al-O, Mg-O and the volume of
MgxZn1-xAl2O4.
Table 4: Calculated refractive index for MgxZn1-xAl2O4 (x=0, x=1)
n (0)
n
Maximum Energy(eV)
ZnAl2O4
Present 1.77 2.13 10.15
Calc. 1.63 2.16 10.71
Calc. 1.74
MgAl2O4
Present 1.71 2.16 8.84
Calc.53 1.61 2.40 11.35
Calc.25 1.763
Expt.54 1.71
The different band structure is responsible for diffe-
rent structure in the optical spectra because the optical
spectra originates from the transitions from the top va-
lence band to the bottom conduction band. For instance,
a different band gap corresponds to different absorption
edge. The relationship between the absorption edge and
the band gap is described by Equation (6):
hc
d T
A
hc
Eg 
⎛
⎝
⎜ ⎞
⎠
⎟ ⎛
⎝
⎜ ⎞
⎠
⎟ ⎛
⎝
⎜ ⎞
⎠
⎟ = −⎛⎝
⎜ ⎞
⎠
⎟
2 2
21 1ln (6)
where  is the wavelength, T is the transmissivity, A is
the transition coefficients of the direct band gap, d is the
thickness and Eg is the value of the band gap. As shown
in Figure 8a, it is found that MgxZn1-xAl2O4(x=0) shows
an absorption shoulder around 360 nm, which is in
agreement with the experimental value. For instance, the
optical absorbance spectrum of ZnAl2O4 was detected in
the 250–400 nm region at room temperature,17 and its
adsorption shoulder was also observed in the 290–375 nm
wavelength region for the sample calcined at 700 °C.52
MgxZn1-xAl2O4(x=1) shows an absorption shoulder
around 270 nm, meaning that the absorption shoulder
undergoes a small blue shift. This is in agreement with
the imaginary part 
() of the dielectric function, and
also suggests that perfect MgAl2O4 spinel is a
transparent crystal, and its transparency is higher than
that of ZnAl2O4 in the visible region. Meanwhile, the
absorption shoulder is 340 nm for x=0.25, 310 nm for
x=0.50, 290 nm for x=0.75 and 270 nm for x=1.
Moreover, as shown in Figure 8b, where it seems that
the absorption region shifts towards lower energy for
ZnAl2O4 than MgAl2O4 due to the smaller band gap of
the former than the latter, as predicted by the imaginary
part of the dielectric function as well as the data derived
from Equation (6). The absorption edge starts from
about (3.66, 3.95, 4.18, 4.50 and 4.86) eV from x=0 to
x=1. This is mainly attributed to electronic transitions
from the O-2p located at the top of the valence band to
the empty Zn-4s electron states dominating the bottom
of the conduction band for x=0. Since Zn is replaced by
Mg, electronic transitions from O-2p to the empty
Mg-3p electron states, yielding an increase in band gap,
so that absorption edge starts at higher energy (4.86
eV), due to Mg-3p electron states dominating the
bottom of the conduction band for x=1. Other peaks
locates at (7.46, 11.77, 13.80, 16.05 and 17.15) eV for
x=0, (6.75, 7.57, 11.91, 13.63 and 17.30) eV for x=0.25,
(5.61, 13.41 and 17.35) for x=0.5, (5.85, 10.49, 13.21
and 17.58) eV for x=0.75 and (6.28, 7.76, 10.21, 13.33,
15.11 and 17.66) eV for x=1, in comparison with other
theoretical values listed in Table 5.
Table 5: The adsorption spectra of perfect MgxZn1-xAl2O4 (x=0, x=1).
ZnAl2O4 MgAl2O4
Present Calc.18 Calc.40 Present Calc.14
Threshold
Peak
position
3.66 4.20 ~4.31 4.86 ~5.05
7.46 7.83 ~8.03 6.28 ~6.14
13.80 14.17 7.76 ~7.03
16.05 16.41 10.21 ~10.15
17.15 17.35 13.33 13.75
The energy-loss function has a large effect on usage
of material. For example, the sharper peak of energy
loss, explained narrower scope, the better the optical
storage efficiency. Besides, energy loss function is an
important factor for describing the energy loss of a fast
electron traversing in a material, and the prominent
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
740 Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: Absorption spectra of Mg doped ions at various doping
concentration: a) the unit of the vertical axis for absorption spectra is
wavelength (nm), b) the unit of abscissa is energy (eV)
peaks in the energy-loss function representing a charac-
teristic associated with the plasma resonance (a collec-
tive oscillation of the valence electrons).51 The energy-
loss function is displayed in Figure 9. It can be seen that
the primary peak emerges at 22.43 eV for x=0
(ZnAl2O4), which is consistent with the theoretical value
of 20.99 eV.18 Additionally, the energy-loss function for
ZnAl2O4 between 0 eV and 12 eV is in agreement with a
previous report.40 Other peaks are located at 22.19 eV for
x=0.25, 22.00 eV for x=0.5, 21.78 eV for x=0.75 and
21.60 eV for x=1.
There is close relationship between reflectivity and
band gap, i.e., the higher coefficient of reflectivity (R(0))
at zero frequency corresponds to a smaller band gap.
Additionally, the reflection takes an important role in
optical losses. As can be seen from Figure 10 the high
reflectivity localizes between 18.61 eV and 22.32 eV,
which is consistent with the energy-loss function men-
tioned above.
4 CONCLUSIONS
The effect of Mg+2 versus Zn+2 by comparing the
structural, electronic and optical properties of
MgxZn1-xAl2O4 (from x=0 to x=1) has been implemented
using the DFT+U method. The calculations show that the
volume of the unit cell changes linearly with increasing
Mg-doping amount (x). The Al-O bond length of
MgxZn1-xAl2O4 increases with x, suggesting it has a
weaker bond stiffness. On the basis of the imaginary
parts of the dielectric function, it is found that critical
points occur at about 3.49 eV for ZnAl2O4, which is
lower than that for Mg-doped ZnAl2O4, indicating that
absorption shoulder undergoes a blue shift in the UV
region as evidenced by the absorption spectra. The blue
shift is probably due to the electronic transition from
O-2p to the empty Mg-3p electron states. For real parts
of dielectric function, the limit of zero energy has a
square fit relationship with the refractive index (n(0))
localized zero energy, which is verified by the reflec-
tivity spectra. Furthermore, the absorbance spectra,
reflectivity spectra and the refractive index show that
these compounds are suitable for UV reflective coatings.
The peak of the energy-loss function shows that
MgxZn1-xAl2O4 can be a candidate for optical storage ma-
terials.
Acknowledgment
This work is supported by the project from
Chongqing Municipal Education Commission (No.
KJ1601218), Young Foundation of Yangtze Normai
University (No. 2015XJXM29) and Talent Program of
Yangtze Normai University (No. 2015KYQD03).
6 REFERENCES
1 N. Bouropoulos, I. Tsiaoussis, P. Poulopoulos, P. Roditis, S. Bas-
koutas, ZnO controllable sized quantum dots produced by polyol
method: an experimental and theoretical study, Mater. Lett., 62
(2008), 3533–3535, doi:10.1016/j.matlet.2008.03.044
2 W. S. Tzing, W. H. Tuan, The strength of duplex Al2O3-ZnAl2O4
composite, J. Mater. Sci. Lett., 15.16 (1996), 1395–1396,
doi:10.1007/BF00275286
3 S. A.T. Redfen, R. J. Harrison, H. S. C. O’Neill, D. R. R. Wood,
Thermodynamics and kinetics of cation ordering in MgAl2O4 spinel
up to 1600 °C from in situ neutron diffraction, Amer. Mineral., 84.3
(1999), 299–310, doi:10.2138/am-1999-0313
4 H. Cynn, S. K. Sharma, T. F. Cooney, M. Nicol, High-temperature
Raman investigation of order-disorder behavior in the MgAl2O4
spinel, Phys. Rev. B, 45.1 (1992), 500–502, doi:10.1103/PhysRevB.
45.500
5 QivaStar. Inc, Ridgeerest. CA, and Hugest Space and Communi-
cation CO., El. Segundo, CA, U. S. Pat. N0. 5820669, 1998
6 K. P. Surendran, P.V. Bijumon, P. Mohanan, M.T. Sebastian,
(1-x)MgAl2O4-xTiO2 dielectrics for microwave and millimeter wave
applications, Appl. Phys., A 81.4 (2005), 823–826, doi:10.1007/
s00339-005-3282-5
7 P. Fu, W. Z. Lu, W. Lei, Y. Xu, X. H. Wang, J. M. Wu, Transparent
polycrystalline MgAl2O4 ceramic fabricated by spark plasma
sintering: Microwave dielectric and optical properties, Ceram. Int.,
39.3 (2013), 2481–2487, doi:10.1016/j.ceramint.2012.09.006
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743 741
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 10: Calculated reflection of MgxZn1-xAl2O4 obtained from the
CASTEP calculation
Figure 9: Calculated energy loss function of MgxZn1-xAl2O4 obtained
from the CASTEP calculations
index n(0) at the zero frequency limits is 1.77, 1.76,
1.73, 1.72 and 1.71 from x=0 to x=1, which are con-
sistent with values derived from the real part 	() of the
dielectric function mentioned above. They reach a maxi-
mum value of 2.12 at 10.05 eV for x=0, 2.14 at 9.01 eV
for x=0.25, 2.13 at 8.97 eV for x=0.5, 2.14 at 8.95 ev for
x=0.75 and 2.16 at 8.84 eV for x=1, compared to other
theoretical values listed in Table 4. We note that the
refractive index n(0) decreases with an increase of x. It
indicates that the n(0) likely is decided by the bond
length of Zn-O, Al-O, Mg-O and the volume of
MgxZn1-xAl2O4.
Table 4: Calculated refractive index for MgxZn1-xAl2O4 (x=0, x=1)
n (0)
n
Maximum Energy(eV)
ZnAl2O4
Present 1.77 2.13 10.15
Calc. 1.63 2.16 10.71
Calc. 1.74
MgAl2O4
Present 1.71 2.16 8.84
Calc.53 1.61 2.40 11.35
Calc.25 1.763
Expt.54 1.71
The different band structure is responsible for diffe-
rent structure in the optical spectra because the optical
spectra originates from the transitions from the top va-
lence band to the bottom conduction band. For instance,
a different band gap corresponds to different absorption
edge. The relationship between the absorption edge and
the band gap is described by Equation (6):
hc
d T
A
hc
Eg 
⎛
⎝
⎜ ⎞
⎠
⎟ ⎛
⎝
⎜ ⎞
⎠
⎟ ⎛
⎝
⎜ ⎞
⎠
⎟ = −⎛⎝
⎜ ⎞
⎠
⎟
2 2
21 1ln (6)
where  is the wavelength, T is the transmissivity, A is
the transition coefficients of the direct band gap, d is the
thickness and Eg is the value of the band gap. As shown
in Figure 8a, it is found that MgxZn1-xAl2O4(x=0) shows
an absorption shoulder around 360 nm, which is in
agreement with the experimental value. For instance, the
optical absorbance spectrum of ZnAl2O4 was detected in
the 250–400 nm region at room temperature,17 and its
adsorption shoulder was also observed in the 290–375 nm
wavelength region for the sample calcined at 700 °C.52
MgxZn1-xAl2O4(x=1) shows an absorption shoulder
around 270 nm, meaning that the absorption shoulder
undergoes a small blue shift. This is in agreement with
the imaginary part 
() of the dielectric function, and
also suggests that perfect MgAl2O4 spinel is a
transparent crystal, and its transparency is higher than
that of ZnAl2O4 in the visible region. Meanwhile, the
absorption shoulder is 340 nm for x=0.25, 310 nm for
x=0.50, 290 nm for x=0.75 and 270 nm for x=1.
Moreover, as shown in Figure 8b, where it seems that
the absorption region shifts towards lower energy for
ZnAl2O4 than MgAl2O4 due to the smaller band gap of
the former than the latter, as predicted by the imaginary
part of the dielectric function as well as the data derived
from Equation (6). The absorption edge starts from
about (3.66, 3.95, 4.18, 4.50 and 4.86) eV from x=0 to
x=1. This is mainly attributed to electronic transitions
from the O-2p located at the top of the valence band to
the empty Zn-4s electron states dominating the bottom
of the conduction band for x=0. Since Zn is replaced by
Mg, electronic transitions from O-2p to the empty
Mg-3p electron states, yielding an increase in band gap,
so that absorption edge starts at higher energy (4.86
eV), due to Mg-3p electron states dominating the
bottom of the conduction band for x=1. Other peaks
locates at (7.46, 11.77, 13.80, 16.05 and 17.15) eV for
x=0, (6.75, 7.57, 11.91, 13.63 and 17.30) eV for x=0.25,
(5.61, 13.41 and 17.35) for x=0.5, (5.85, 10.49, 13.21
and 17.58) eV for x=0.75 and (6.28, 7.76, 10.21, 13.33,
15.11 and 17.66) eV for x=1, in comparison with other
theoretical values listed in Table 5.
Table 5: The adsorption spectra of perfect MgxZn1-xAl2O4 (x=0, x=1).
ZnAl2O4 MgAl2O4
Present Calc.18 Calc.40 Present Calc.14
Threshold
Peak
position
3.66 4.20 ~4.31 4.86 ~5.05
7.46 7.83 ~8.03 6.28 ~6.14
13.80 14.17 7.76 ~7.03
16.05 16.41 10.21 ~10.15
17.15 17.35 13.33 13.75
The energy-loss function has a large effect on usage
of material. For example, the sharper peak of energy
loss, explained narrower scope, the better the optical
storage efficiency. Besides, energy loss function is an
important factor for describing the energy loss of a fast
electron traversing in a material, and the prominent
C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
740 Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: Absorption spectra of Mg doped ions at various doping
concentration: a) the unit of the vertical axis for absorption spectra is
wavelength (nm), b) the unit of abscissa is energy (eV)
peaks in the energy-loss function representing a charac-
teristic associated with the plasma resonance (a collec-
tive oscillation of the valence electrons).51 The energy-
loss function is displayed in Figure 9. It can be seen that
the primary peak emerges at 22.43 eV for x=0
(ZnAl2O4), which is consistent with the theoretical value
of 20.99 eV.18 Additionally, the energy-loss function for
ZnAl2O4 between 0 eV and 12 eV is in agreement with a
previous report.40 Other peaks are located at 22.19 eV for
x=0.25, 22.00 eV for x=0.5, 21.78 eV for x=0.75 and
21.60 eV for x=1.
There is close relationship between reflectivity and
band gap, i.e., the higher coefficient of reflectivity (R(0))
at zero frequency corresponds to a smaller band gap.
Additionally, the reflection takes an important role in
optical losses. As can be seen from Figure 10 the high
reflectivity localizes between 18.61 eV and 22.32 eV,
which is consistent with the energy-loss function men-
tioned above.
4 CONCLUSIONS
The effect of Mg+2 versus Zn+2 by comparing the
structural, electronic and optical properties of
MgxZn1-xAl2O4 (from x=0 to x=1) has been implemented
using the DFT+U method. The calculations show that the
volume of the unit cell changes linearly with increasing
Mg-doping amount (x). The Al-O bond length of
MgxZn1-xAl2O4 increases with x, suggesting it has a
weaker bond stiffness. On the basis of the imaginary
parts of the dielectric function, it is found that critical
points occur at about 3.49 eV for ZnAl2O4, which is
lower than that for Mg-doped ZnAl2O4, indicating that
absorption shoulder undergoes a blue shift in the UV
region as evidenced by the absorption spectra. The blue
shift is probably due to the electronic transition from
O-2p to the empty Mg-3p electron states. For real parts
of dielectric function, the limit of zero energy has a
square fit relationship with the refractive index (n(0))
localized zero energy, which is verified by the reflec-
tivity spectra. Furthermore, the absorbance spectra,
reflectivity spectra and the refractive index show that
these compounds are suitable for UV reflective coatings.
The peak of the energy-loss function shows that
MgxZn1-xAl2O4 can be a candidate for optical storage ma-
terials.
Acknowledgment
This work is supported by the project from
Chongqing Municipal Education Commission (No.
KJ1601218), Young Foundation of Yangtze Normai
University (No. 2015XJXM29) and Talent Program of
Yangtze Normai University (No. 2015KYQD03).
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C. XIANG et al.: ELECTRONIC AND OPTICAL PROPERTIES OF THE SPINEL OXIDES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 735–743 743
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
S.-H. KIM et al.: SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING TO MICRO-PULSE
745–749
SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING
TO MICRO-PULSE ELECTROCHEMICAL MACHINING
KARAKTERISTIKE POVR[INE INVAR ZLITINE GLEDE NA
MIKROPULZNO ELEKTROKEMI^NO OBDELAVO
Seong-Hyun Kim, Seung-Geon Choi, Woong-Kirl Choi, Eun-Sang Lee
Inha University, Department of Mechanical Engineering, 253 Yonghyun-Dong, Nam-Gu, Incheon 402-751, Republic of Korea
kimseonghyun@inha.edu, leees@inha.ac.kr
Prejem rokopisa – received: 2016-07-15; sprejem za objavo – accepted for publication: 2017-01-24
doi:10.17222/mit.2016.187
Invar is a 36 % nickel iron alloy that has a low thermal expansion, compared to other metals and alloys, at temperatures ranging
from room temperature up to approximately 230 °C. Invar alloy is ductile, easily weldable and its machinability is similar to that
of austenitic stainless steel. Due to the low thermal expansion of Invar, it is used for shadow masks for display devices such as
UHDTV – organic light-emitting diode. In this study, micro-pulse electrochemical machining (MPECM), which is a non-contact
ultra-precision machining method, was developed to manufacture Invar sheets; optimum parameters of MPECM were defined
and the basic MPECM experiments were carried out on an Invar sheet. The optimum parameters were determined with pulse-on
time and duty-ratio analysis. The experimental results show that MPECM is hard to control. Therefore, using ultrashort charging
times and very high pulses, it is possible to achieve a successful anodic dissolution at a very small electrode gap. Hence, a
longer pulse-on time and a small electrode gap may provide the scope for further improvement of the machining accuracy by
controlling the localization effect. Furthermore, the machining depth and MPECM efficiency were investigated with respect to
various parameters and pulse-on time, considering different duty-ratio conditions.
Keywords: Invar alloy, electrochemical machining, surface characteristics
Invar je 36 % zlitina niklja in `eleza, ki ima v primerjavi z vsemi kovinami in zlitinami, nizek koeficient toplotnega raztezka v
obmo~ju sobne temperature pa do okoli 230 ° C. Invar zlitina je duktilna, primerna za enostavno varjenje in njena obdelovalnost
je podobna avstenitnemu nerjavnemu jeklu. Zaradi nizkega koeficienta termi~nega raztezka, se zlitina Invar uporablja v maskah
za sen~enje displejev v OLED-UHD TV (angl.: Organic Light Emitting Diode – Ultra High Definition TeleVision) napravah. V
tej {tudiji je bila razvita metoda mikropulzne elektrokemi~ne obdelave kovin (angl. MPECM), ki je brezkontaktna ultraprecizna
tehnologija mehanske obdelave. Izvedeni so bili osnovni poskusi na zlitini Invar. Dolo~eni so bili optimalni parametri postopka.
Eksperimentalni rezultati ka`ejo, da je obdelovalnost z MPECM te`ko nadzorovati. Izjemno kratki ~asi, z zelo visokimi
napetostnimi impulzi, omogo~ajo nastanek pogojev za uspe{no anodno raztapljanje kovine pri zelo majhni re`i elektrode. Zato
se pri zelo visokem {tevilov impulzov in manj{em razmiku elektrod, lahko zagotovi nadaljnje izbolj{anje to~nosti obdelave z
MPECM. Nadalje je bila raziskana obdelovalna globina in obdelovalnost z MPECM pri razli~nih procesnih parametrih
(napetost-~as-storilnost).
Klju~ne besede: Invar zlitina, elektrokemi~na obdelava, povr{inske karakteristike
1 INTRODUCTION
Invar is an alloy of 64 % iron and 36 % nickel. Its
coefficient of expansion is very low, namely 1.8 × 10–6
cm/°C, at temperatures ranging from a cryogenic tem-
perature of –196 °C to a Currie temperature of 260 °C.
This value is 10 times lower than that of stainless steel
and 100 times lower than that of iron, which makes Invar
appropriate for applications in various fields of study.
Invar is used in ultra-high-definition televisions, display
shadow masks, organic light-emitting diodes, controllers
for quality, colors, and for the definition of images in
mobile devices and other instruments that require precise
measurements.1–3 Invar is also used in electron-gun
electrodes (devices that eject an electron beam into the
hole in a shadow mask), lead frames (to protect semicon-
ductor chips from external damage) and other industrial
devices.
Invar is currently manufactured with mechanical pro-
cesses or electroforming; other methodologies have also
been suggested. To maintain the specific characteristics
of Invar alloys, advanced machining methods are being
studied, such as laser machining, supersonic-wave ma-
chining, micro-end mills/drilling, electropolishing and
electrochemical machining.4–6 Moreover, the effects of
hydrodynamics and temperature on the electrode po-
sition of Fe–Ni (Invar) alloys has been investigated with
DC, pulse, and pulse-reverse electrodeposition tech-
niques.7,8 MPECM has seen a resurgence of industrial
interest within the last decade, due to its many
advantages, such as no tool wear, stress-free and smooth
surfaces, and the ability to machine complexly shaped
products made from different materials, regardless of
their hardness and high strength, high tension, or
whether they are heat-resistant materials.9 In this study,
the possibility of pulse electrochemical machining is
investigated by analyzing the machinability characteris-
tics of processed Invar prior to pulse electrochemical
machining.
Materiali in tehnologije / Materials and technology 51 (2017) 5, 745–749 745
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 62-4-023.7:669.018.47:621.039.333 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)745(2017)
S.-H. KIM et al.: SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING TO MICRO-PULSE
745–749
SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING
TO MICRO-PULSE ELECTROCHEMICAL MACHINING
KARAKTERISTIKE POVR[INE INVAR ZLITINE GLEDE NA
MIKROPULZNO ELEKTROKEMI^NO OBDELAVO
Seong-Hyun Kim, Seung-Geon Choi, Woong-Kirl Choi, Eun-Sang Lee
Inha University, Department of Mechanical Engineering, 253 Yonghyun-Dong, Nam-Gu, Incheon 402-751, Republic of Korea
kimseonghyun@inha.edu, leees@inha.ac.kr
Prejem rokopisa – received: 2016-07-15; sprejem za objavo – accepted for publication: 2017-01-24
doi:10.17222/mit.2016.187
Invar is a 36 % nickel iron alloy that has a low thermal expansion, compared to other metals and alloys, at temperatures ranging
from room temperature up to approximately 230 °C. Invar alloy is ductile, easily weldable and its machinability is similar to that
of austenitic stainless steel. Due to the low thermal expansion of Invar, it is used for shadow masks for display devices such as
UHDTV – organic light-emitting diode. In this study, micro-pulse electrochemical machining (MPECM), which is a non-contact
ultra-precision machining method, was developed to manufacture Invar sheets; optimum parameters of MPECM were defined
and the basic MPECM experiments were carried out on an Invar sheet. The optimum parameters were determined with pulse-on
time and duty-ratio analysis. The experimental results show that MPECM is hard to control. Therefore, using ultrashort charging
times and very high pulses, it is possible to achieve a successful anodic dissolution at a very small electrode gap. Hence, a
longer pulse-on time and a small electrode gap may provide the scope for further improvement of the machining accuracy by
controlling the localization effect. Furthermore, the machining depth and MPECM efficiency were investigated with respect to
various parameters and pulse-on time, considering different duty-ratio conditions.
Keywords: Invar alloy, electrochemical machining, surface characteristics
Invar je 36 % zlitina niklja in `eleza, ki ima v primerjavi z vsemi kovinami in zlitinami, nizek koeficient toplotnega raztezka v
obmo~ju sobne temperature pa do okoli 230 ° C. Invar zlitina je duktilna, primerna za enostavno varjenje in njena obdelovalnost
je podobna avstenitnemu nerjavnemu jeklu. Zaradi nizkega koeficienta termi~nega raztezka, se zlitina Invar uporablja v maskah
za sen~enje displejev v OLED-UHD TV (angl.: Organic Light Emitting Diode – Ultra High Definition TeleVision) napravah. V
tej {tudiji je bila razvita metoda mikropulzne elektrokemi~ne obdelave kovin (angl. MPECM), ki je brezkontaktna ultraprecizna
tehnologija mehanske obdelave. Izvedeni so bili osnovni poskusi na zlitini Invar. Dolo~eni so bili optimalni parametri postopka.
Eksperimentalni rezultati ka`ejo, da je obdelovalnost z MPECM te`ko nadzorovati. Izjemno kratki ~asi, z zelo visokimi
napetostnimi impulzi, omogo~ajo nastanek pogojev za uspe{no anodno raztapljanje kovine pri zelo majhni re`i elektrode. Zato
se pri zelo visokem {tevilov impulzov in manj{em razmiku elektrod, lahko zagotovi nadaljnje izbolj{anje to~nosti obdelave z
MPECM. Nadalje je bila raziskana obdelovalna globina in obdelovalnost z MPECM pri razli~nih procesnih parametrih
(napetost-~as-storilnost).
Klju~ne besede: Invar zlitina, elektrokemi~na obdelava, povr{inske karakteristike
1 INTRODUCTION
Invar is an alloy of 64 % iron and 36 % nickel. Its
coefficient of expansion is very low, namely 1.8 × 10–6
cm/°C, at temperatures ranging from a cryogenic tem-
perature of –196 °C to a Currie temperature of 260 °C.
This value is 10 times lower than that of stainless steel
and 100 times lower than that of iron, which makes Invar
appropriate for applications in various fields of study.
Invar is used in ultra-high-definition televisions, display
shadow masks, organic light-emitting diodes, controllers
for quality, colors, and for the definition of images in
mobile devices and other instruments that require precise
measurements.1–3 Invar is also used in electron-gun
electrodes (devices that eject an electron beam into the
hole in a shadow mask), lead frames (to protect semicon-
ductor chips from external damage) and other industrial
devices.
Invar is currently manufactured with mechanical pro-
cesses or electroforming; other methodologies have also
been suggested. To maintain the specific characteristics
of Invar alloys, advanced machining methods are being
studied, such as laser machining, supersonic-wave ma-
chining, micro-end mills/drilling, electropolishing and
electrochemical machining.4–6 Moreover, the effects of
hydrodynamics and temperature on the electrode po-
sition of Fe–Ni (Invar) alloys has been investigated with
DC, pulse, and pulse-reverse electrodeposition tech-
niques.7,8 MPECM has seen a resurgence of industrial
interest within the last decade, due to its many
advantages, such as no tool wear, stress-free and smooth
surfaces, and the ability to machine complexly shaped
products made from different materials, regardless of
their hardness and high strength, high tension, or
whether they are heat-resistant materials.9 In this study,
the possibility of pulse electrochemical machining is
investigated by analyzing the machinability characteris-
tics of processed Invar prior to pulse electrochemical
machining.
Materiali in tehnologije / Materials and technology 51 (2017) 5, 745–749 745
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 62-4-023.7:669.018.47:621.039.333 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)745(2017)
2 EXPERIMENTAL PART
2.1 Theory of the MPECM system
Figure 1 shows a schematic diagram of the combined
electrochemical process and the MPECM set-up with a
previously developed microprobe. A pulsed voltage from
a high-voltage pulse generator (Hewlett-Packard
HP8116A) is applied to the gap between the anode and
cathode, and the electrolyte solution is supplied through
a hole in the cathode or an external electrolyte supplier.
This homogenizes the electrochemical reaction and
removes the gas and heat generated. The anode is the
Invar alloy with a thickness of 30 μm and the cathode is
stainless steel with a 1-mm diameter.
The fully programmable 50 MHz pulse generator has
a variable pulse width (10 ns min) and the maximum
output amplitude of the maximum voltage of 32 V p-p
(in an open circuit). The high-speed data-acquisition sys-
tem was provided by a TDS 2002C digital oscilloscope
(Tektronix) with a bandwidth of 200 MHz and a
sampling rate of 500 MS/s. Typical experimental condi-
tions are shown in Table 1.
Table 1: Experimental conditions
Parameter Value
Pulse generator
spec.
HP8116A (Hewlett-Packard), 50MHz,
32 V p-p, various (10 ns) pulse widths
Workpiece (anode) Invar (Fe 64 %, Ni 36 %) 30-μmthickness
Electrode (cathode)
Stainless steel 304 (ø 1 mm)
Non-insulation, Insulation
Digital oscilloscope Tektronix TDS 2002C
MPECM time 40 min (pulse-on time: 500 ns-2 μs,Duty factor: 33–66 %)
Applied voltage (V) 16
Electrode gap (μm) 50
Electrolyte Sodium nitrite + sodium tartrate + DIwater
Measurement
Optical microscopic analysis, Scanning
electron microscopy (SEM),
Non-contact 3D profile
2.2 Electrode type
The concentration of the current at the edges and
irregular parts of an electrode can lead to localized
sedimentation on the anode surface. A specially designed
electrode is recommended for electrochemical micro-
machining to control the dimensions of the Invar surface.
Figure 2 compares the surfaces that were electrochemi-
cally micro-machined using two different electrodes.
Electrode 1, with the traditional design, had a concen-
trated current at the edge, which is marked as a 2.
Electrode 2 has a non-conducting part that is marked as a
3. The current distribution for electrode 2 is uniform on
the surface, which is marked as a 4. Localized sedimen-
tation was identified on the anode workpiece surface in
the shape profile created by electrode 1, but none was
identified for electrode 2. This results from non-conduct-
ing part 3 on the side of the electrode. The cathode
design should be taken into account to precisely control
the shape and prevent sedimentation.
When the electrodes are far apart, the voltage flows
in the regions other than the processed area, even when
no machining is intended. To reduce the unwanted
voltage flowing in these regions, an electrode that is
coated on the top can be used as the tool. To minimize
this voltage and observe the behavior of the machining
current, the sides of the tool need to be completely
insulated.
3 RESULTS AND DISCUSSION
3.1 Machining characteristics versus the pulse-on time
using a non-insulated electrode
The machining characteristics were studied using a
voltage of 16 V, a gap of 50 μm a pulse-off time of 1 μs
and a machining time of 40 min. An insulated electrode
was used, and the pulse-on time was changed from
500 ns to 2 μs. Figure 3 shows that the machining radius
was larger as the pulse-on time increased. The processed
radius was 559.04 μm at the pulse-on time of 500 ns,
620.76 μm at 600 ns, 675.21 μm at 800 ns, 689.73 μm at
1 μs, 693.36 μm at 1.5 μs, and 715.14 μm at 2 μs. When
the pulse-on time was 1 μs or higher, the machining
radius drastically increased. Figure 4 shows the results
after the machining with respect to the pulse-on time.
S.-H. KIM et al.: SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING TO MICRO-PULSE
746 Materiali in tehnologije / Materials and technology 51 (2017) 5, 745–749
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Comparison between two different electrodes
Figure 1: Micro-pulse electrochemical machining system
When the pulse-on time was 500–800 ns, the machinabi-
lity declined due to a low current density, but at 1–2 μs,
the machinability improved due to a high current density.
The precision and the resulting image became better and
better.
3.2 Machining characteristics versus the pulse-on time
using an insulated electrode
The machining characteristics were also examined
using an insulated electrode, with the other conditions
kept the same as in the previous experiment. Figure 5
shows that the machining radius was consistent with the
increase in the pulse-on time. The processed radius was
642.54 μm at 500 ns, 656.81 μm at 600 ns, 644.76 μm at
800 ns, 638.72 μm at 1 μs, 647.77 μm at 1.5 μs, and
656.81 μm at 2 μs.
Figure 6 shows the results after the machining. The
machining may occur on the sides of the electrode with-
out the insulation, but the insulated electrode prevents
this and makes the machining consistent with a very fine
outcome.
3.3 Surface analysis versus the pulse-on time using
insulated and uninsulated electrodes
Figure 7a shows a SEM image of the Invar film
surface after the machining with the uninsulated elec-
trode. The processed area is clearly displayed. Figure 7b
shows the non-contact 3D profile measured. The tapering
remained on the side of the Invar film surface after the
experiment. Figure 7c shows the results for the pulse-on
time of 2 μs, which allowed the best machinability with
both types of electrode. Figure 8 shows the equivalent
results obtained with the insulated electrode. The voltage
in Figure 8c is slightly increased compared to Figure
7c.
With the uninsulated electrode, the processed radius
is increased, with almost no influence on the depth. This
is attributable to the parts other than the intended ma-
chining area with the repeated charging and discharging
in the dual layer of electricity because the voltage pulse
is exhausted on the other sides of the uninsulated elec-
trode. However, with the insulated electrode, the voltage
is concentrated only in the machining area, so the radius
is consistently maintained and the machining depth
increases. Figures 7c and 8c show that the voltage
permitted during the actual machining is lower with the
uninsulated electrode than with the insulated one.
S.-H. KIM et al.: SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING TO MICRO-PULSE
Materiali in tehnologije / Materials and technology 51 (2017) 5, 745–749 747
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Crater-shape patterns with pulse-on time / duty factor and
insulated electrode: a) 500 ns / 33 %, b) 600 ns / 37 %, c) 800 ns / 44
%, d) 1 μs / 50 %, e) 1.5 μs / 60 %, f) 2 μs / 66 %
Figure 5: Increase of hole-diameter size versus pulse-on time
Figure 3: Increase of hole-diameter size versus pulse-on time
Figure 4: Crater-shape patterns with pulse-on time / duty factor and
non-insulated electrode: a) 500 ns / 33 %, b) 600 ns / 37 %, c) 800 ns /
40 %, d) 1 μs / 50 %, e) 1.5 μs / 60 %, f) 2 μs / 66 %
2 EXPERIMENTAL PART
2.1 Theory of the MPECM system
Figure 1 shows a schematic diagram of the combined
electrochemical process and the MPECM set-up with a
previously developed microprobe. A pulsed voltage from
a high-voltage pulse generator (Hewlett-Packard
HP8116A) is applied to the gap between the anode and
cathode, and the electrolyte solution is supplied through
a hole in the cathode or an external electrolyte supplier.
This homogenizes the electrochemical reaction and
removes the gas and heat generated. The anode is the
Invar alloy with a thickness of 30 μm and the cathode is
stainless steel with a 1-mm diameter.
The fully programmable 50 MHz pulse generator has
a variable pulse width (10 ns min) and the maximum
output amplitude of the maximum voltage of 32 V p-p
(in an open circuit). The high-speed data-acquisition sys-
tem was provided by a TDS 2002C digital oscilloscope
(Tektronix) with a bandwidth of 200 MHz and a
sampling rate of 500 MS/s. Typical experimental condi-
tions are shown in Table 1.
Table 1: Experimental conditions
Parameter Value
Pulse generator
spec.
HP8116A (Hewlett-Packard), 50MHz,
32 V p-p, various (10 ns) pulse widths
Workpiece (anode) Invar (Fe 64 %, Ni 36 %) 30-μmthickness
Electrode (cathode)
Stainless steel 304 (ø 1 mm)
Non-insulation, Insulation
Digital oscilloscope Tektronix TDS 2002C
MPECM time 40 min (pulse-on time: 500 ns-2 μs,Duty factor: 33–66 %)
Applied voltage (V) 16
Electrode gap (μm) 50
Electrolyte Sodium nitrite + sodium tartrate + DIwater
Measurement
Optical microscopic analysis, Scanning
electron microscopy (SEM),
Non-contact 3D profile
2.2 Electrode type
The concentration of the current at the edges and
irregular parts of an electrode can lead to localized
sedimentation on the anode surface. A specially designed
electrode is recommended for electrochemical micro-
machining to control the dimensions of the Invar surface.
Figure 2 compares the surfaces that were electrochemi-
cally micro-machined using two different electrodes.
Electrode 1, with the traditional design, had a concen-
trated current at the edge, which is marked as a 2.
Electrode 2 has a non-conducting part that is marked as a
3. The current distribution for electrode 2 is uniform on
the surface, which is marked as a 4. Localized sedimen-
tation was identified on the anode workpiece surface in
the shape profile created by electrode 1, but none was
identified for electrode 2. This results from non-conduct-
ing part 3 on the side of the electrode. The cathode
design should be taken into account to precisely control
the shape and prevent sedimentation.
When the electrodes are far apart, the voltage flows
in the regions other than the processed area, even when
no machining is intended. To reduce the unwanted
voltage flowing in these regions, an electrode that is
coated on the top can be used as the tool. To minimize
this voltage and observe the behavior of the machining
current, the sides of the tool need to be completely
insulated.
3 RESULTS AND DISCUSSION
3.1 Machining characteristics versus the pulse-on time
using a non-insulated electrode
The machining characteristics were studied using a
voltage of 16 V, a gap of 50 μm a pulse-off time of 1 μs
and a machining time of 40 min. An insulated electrode
was used, and the pulse-on time was changed from
500 ns to 2 μs. Figure 3 shows that the machining radius
was larger as the pulse-on time increased. The processed
radius was 559.04 μm at the pulse-on time of 500 ns,
620.76 μm at 600 ns, 675.21 μm at 800 ns, 689.73 μm at
1 μs, 693.36 μm at 1.5 μs, and 715.14 μm at 2 μs. When
the pulse-on time was 1 μs or higher, the machining
radius drastically increased. Figure 4 shows the results
after the machining with respect to the pulse-on time.
S.-H. KIM et al.: SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING TO MICRO-PULSE
746 Materiali in tehnologije / Materials and technology 51 (2017) 5, 745–749
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Comparison between two different electrodes
Figure 1: Micro-pulse electrochemical machining system
When the pulse-on time was 500–800 ns, the machinabi-
lity declined due to a low current density, but at 1–2 μs,
the machinability improved due to a high current density.
The precision and the resulting image became better and
better.
3.2 Machining characteristics versus the pulse-on time
using an insulated electrode
The machining characteristics were also examined
using an insulated electrode, with the other conditions
kept the same as in the previous experiment. Figure 5
shows that the machining radius was consistent with the
increase in the pulse-on time. The processed radius was
642.54 μm at 500 ns, 656.81 μm at 600 ns, 644.76 μm at
800 ns, 638.72 μm at 1 μs, 647.77 μm at 1.5 μs, and
656.81 μm at 2 μs.
Figure 6 shows the results after the machining. The
machining may occur on the sides of the electrode with-
out the insulation, but the insulated electrode prevents
this and makes the machining consistent with a very fine
outcome.
3.3 Surface analysis versus the pulse-on time using
insulated and uninsulated electrodes
Figure 7a shows a SEM image of the Invar film
surface after the machining with the uninsulated elec-
trode. The processed area is clearly displayed. Figure 7b
shows the non-contact 3D profile measured. The tapering
remained on the side of the Invar film surface after the
experiment. Figure 7c shows the results for the pulse-on
time of 2 μs, which allowed the best machinability with
both types of electrode. Figure 8 shows the equivalent
results obtained with the insulated electrode. The voltage
in Figure 8c is slightly increased compared to Figure
7c.
With the uninsulated electrode, the processed radius
is increased, with almost no influence on the depth. This
is attributable to the parts other than the intended ma-
chining area with the repeated charging and discharging
in the dual layer of electricity because the voltage pulse
is exhausted on the other sides of the uninsulated elec-
trode. However, with the insulated electrode, the voltage
is concentrated only in the machining area, so the radius
is consistently maintained and the machining depth
increases. Figures 7c and 8c show that the voltage
permitted during the actual machining is lower with the
uninsulated electrode than with the insulated one.
S.-H. KIM et al.: SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING TO MICRO-PULSE
Materiali in tehnologije / Materials and technology 51 (2017) 5, 745–749 747
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Crater-shape patterns with pulse-on time / duty factor and
insulated electrode: a) 500 ns / 33 %, b) 600 ns / 37 %, c) 800 ns / 44
%, d) 1 μs / 50 %, e) 1.5 μs / 60 %, f) 2 μs / 66 %
Figure 5: Increase of hole-diameter size versus pulse-on time
Figure 3: Increase of hole-diameter size versus pulse-on time
Figure 4: Crater-shape patterns with pulse-on time / duty factor and
non-insulated electrode: a) 500 ns / 33 %, b) 600 ns / 37 %, c) 800 ns /
40 %, d) 1 μs / 50 %, e) 1.5 μs / 60 %, f) 2 μs / 66 %
4 CONCLUSIONS
This study demonstrated that MPECM has suitable
process characteristics for machining Invar-alloy sur-
faces. The machined shape and surface qualities, such as
the surface condition and the machining defect, were
observed. The important factors influencing the hole-
shape quality of the Invar alloy are the pulse-on time, the
duty factor, the shape of the cathode and the electrolytic
process parameters. During the application of MPECM
to a fine sheet metal, the electrochemical machining
characteristics of the surface of the alloy were revealed
and analyzed with respect to different electrode shapes,
pulse-on times and applied voltages.
When the uninsulated electrode was used, the
machining speed differed depending on the area of the
tool electrode that was submerged in the liquid. With a
greater machining depth, the machining speed was re-
duced and there was almost no machining after reaching
a certain limit. This is attributable to the current being
exhausted because of the repeated charging and dis-
charging in the dual layer of electricity by the pulse, for
any part other than the intended machining part.
The insulated electrode prevents the permitted volt-
age from being exhausted and improves the machin-
ability with the increasing machining depth. In addition,
since the voltage flows only into the machining area, a
taper does not occur since the sides of the electrode are
well insulated. Accordingly, if insulated electrodes are
used, it is possible to increase the precision of the form
and the processing depth.
This study indicates that the electrode shape and
suitable MPECM conditions affect the machinability,
and that there is a close interrelation between the electric
field on the Invar-alloy surface and the polishing rate. A
comparison of the insulated and uninsulated electrodes
was made, and the effect of the pulse-on time on the
electrochemical machinability was determined. MPECM
can be recommended to practitioners in different fields
to improve the conditions for making holes in the Invar
alloy.
The electrolyte of sodium nitrite + sodium tartrate +
DI water was used and holes with a ø1 mm diameter
were made into the 30-μm-thick Invar alloy. The future
work, therefore, can be made when more engineering
data has been gathered. In addition, more research needs
to be done on the related factors, such as the surface
roughness and the shape accuracy, using controlled
micro-pulses.
Acknowledgment
This work was supported by the Inha University
Research Foundation, Korea.
5 REFERENCES
1 D. Grimmentt, M. Schwartz, K. Node, A comparison of DC and
pulsed Fe–Ni alloy deposits, Journal of The Electrochemical Society,
140 (1993) 4, 973–978, doi:10.1149/1.2056238
2 A. Stupnik, M. Leisch, Study on the surface topology of
vacuum-fired stainless steel by scanning tunnelling microscopy,
Vacuum, 81 (2007) 6, 748–751, doi:10.1016/j.vacuum.2005.11.061
3 L. S. Andrade, S. C. Xavier, R. C. Rocha-Filho, N. Bocchi, S. R.
Biaggio, Electropolishing of AISI-304 stainless steel using an
oxidizing solution originally used for electrochemical coloration,
Electrochimica Acta, 50 (2005) 13, 2623–2627, doi:10.1016/
j.electacta.2004.11.007
4 Y. N. Hu, H. Zhou, L. P. Liao, H. B. Deng, Surface quality analysis
of the electropolishing of cemented carbide, Journal of Materials
Processing Technology, 139 (2003) 1–3, 253–256, doi:10.1016/
S0924-0136(03)00230-9
S.-H. KIM et al.: SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING TO MICRO-PULSE
748 Materiali in tehnologije / Materials and technology 51 (2017) 5, 745–749
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7: SEM images and non-contact 3D profiles of crater-shaped
patterns at different pulse-on times using the uninsulated electrode
Figure 8: SEM images and non-contact 3D profiles of crater-shaped
patterns at different pulse-on times using the insulated electrode
5 D. Grimmentt, M. Schwartz, K. Node, A Comparison of DC and
Pulsed Fe-Ni Alloy Deposits, Journal of The Electrochemical
Society, 140 (1993) 4, 973–978, doi:10.1149/1.2056238
6 S. Fujimoto, K. Tsujino, T. Shibata, Growth and properties of Cr-rich
thick and porous oxide films on Type 304 stainless steel formed by
square wave potential pulse polarisation, Electrochimica Acta, 47
(2001) 47, 543–551, doi:10.1016/S0013-4686(01)00782-4
7 T. Koyano, M. Kunieda, Ultra-Short Pulse ECM Using Electrostatic
Induction Feeding Method, Procedia CIRP, 6 (2013), 390–394,
doi:10.1016/j.procir.2013.03.066
8 M. Boxhammer, S. Altmannshofer, Model Predictive Control in
Pulsed Electrochemical Machining, Journal of Process Control, 24
(2014) 1, 296–303, doi:10.1016/j.jprocont.2013.11.003
9 E. Lee, T. Shin, B. Kim, S. Baek, Investigation of short pulse elec-
trochemical machining for groove process on Ni-Ti shape memory
alloy, International Journal of Precision Engineering and Manu-
facturing, 11 (2010) 1, 113–118, doi:10.1007/s12541-010-0014-3
S.-H. KIM et al.: SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING TO MICRO-PULSE
Materiali in tehnologije / Materials and technology 51 (2017) 5, 745–749 749
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
4 CONCLUSIONS
This study demonstrated that MPECM has suitable
process characteristics for machining Invar-alloy sur-
faces. The machined shape and surface qualities, such as
the surface condition and the machining defect, were
observed. The important factors influencing the hole-
shape quality of the Invar alloy are the pulse-on time, the
duty factor, the shape of the cathode and the electrolytic
process parameters. During the application of MPECM
to a fine sheet metal, the electrochemical machining
characteristics of the surface of the alloy were revealed
and analyzed with respect to different electrode shapes,
pulse-on times and applied voltages.
When the uninsulated electrode was used, the
machining speed differed depending on the area of the
tool electrode that was submerged in the liquid. With a
greater machining depth, the machining speed was re-
duced and there was almost no machining after reaching
a certain limit. This is attributable to the current being
exhausted because of the repeated charging and dis-
charging in the dual layer of electricity by the pulse, for
any part other than the intended machining part.
The insulated electrode prevents the permitted volt-
age from being exhausted and improves the machin-
ability with the increasing machining depth. In addition,
since the voltage flows only into the machining area, a
taper does not occur since the sides of the electrode are
well insulated. Accordingly, if insulated electrodes are
used, it is possible to increase the precision of the form
and the processing depth.
This study indicates that the electrode shape and
suitable MPECM conditions affect the machinability,
and that there is a close interrelation between the electric
field on the Invar-alloy surface and the polishing rate. A
comparison of the insulated and uninsulated electrodes
was made, and the effect of the pulse-on time on the
electrochemical machinability was determined. MPECM
can be recommended to practitioners in different fields
to improve the conditions for making holes in the Invar
alloy.
The electrolyte of sodium nitrite + sodium tartrate +
DI water was used and holes with a ø1 mm diameter
were made into the 30-μm-thick Invar alloy. The future
work, therefore, can be made when more engineering
data has been gathered. In addition, more research needs
to be done on the related factors, such as the surface
roughness and the shape accuracy, using controlled
micro-pulses.
Acknowledgment
This work was supported by the Inha University
Research Foundation, Korea.
5 REFERENCES
1 D. Grimmentt, M. Schwartz, K. Node, A comparison of DC and
pulsed Fe–Ni alloy deposits, Journal of The Electrochemical Society,
140 (1993) 4, 973–978, doi:10.1149/1.2056238
2 A. Stupnik, M. Leisch, Study on the surface topology of
vacuum-fired stainless steel by scanning tunnelling microscopy,
Vacuum, 81 (2007) 6, 748–751, doi:10.1016/j.vacuum.2005.11.061
3 L. S. Andrade, S. C. Xavier, R. C. Rocha-Filho, N. Bocchi, S. R.
Biaggio, Electropolishing of AISI-304 stainless steel using an
oxidizing solution originally used for electrochemical coloration,
Electrochimica Acta, 50 (2005) 13, 2623–2627, doi:10.1016/
j.electacta.2004.11.007
4 Y. N. Hu, H. Zhou, L. P. Liao, H. B. Deng, Surface quality analysis
of the electropolishing of cemented carbide, Journal of Materials
Processing Technology, 139 (2003) 1–3, 253–256, doi:10.1016/
S0924-0136(03)00230-9
S.-H. KIM et al.: SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING TO MICRO-PULSE
748 Materiali in tehnologije / Materials and technology 51 (2017) 5, 745–749
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7: SEM images and non-contact 3D profiles of crater-shaped
patterns at different pulse-on times using the uninsulated electrode
Figure 8: SEM images and non-contact 3D profiles of crater-shaped
patterns at different pulse-on times using the insulated electrode
5 D. Grimmentt, M. Schwartz, K. Node, A Comparison of DC and
Pulsed Fe-Ni Alloy Deposits, Journal of The Electrochemical
Society, 140 (1993) 4, 973–978, doi:10.1149/1.2056238
6 S. Fujimoto, K. Tsujino, T. Shibata, Growth and properties of Cr-rich
thick and porous oxide films on Type 304 stainless steel formed by
square wave potential pulse polarisation, Electrochimica Acta, 47
(2001) 47, 543–551, doi:10.1016/S0013-4686(01)00782-4
7 T. Koyano, M. Kunieda, Ultra-Short Pulse ECM Using Electrostatic
Induction Feeding Method, Procedia CIRP, 6 (2013), 390–394,
doi:10.1016/j.procir.2013.03.066
8 M. Boxhammer, S. Altmannshofer, Model Predictive Control in
Pulsed Electrochemical Machining, Journal of Process Control, 24
(2014) 1, 296–303, doi:10.1016/j.jprocont.2013.11.003
9 E. Lee, T. Shin, B. Kim, S. Baek, Investigation of short pulse elec-
trochemical machining for groove process on Ni-Ti shape memory
alloy, International Journal of Precision Engineering and Manu-
facturing, 11 (2010) 1, 113–118, doi:10.1007/s12541-010-0014-3
S.-H. KIM et al.: SURFACE CHARACTERISTICS OF INVAR ALLOY ACCORDING TO MICRO-PULSE
Materiali in tehnologije / Materials and technology 51 (2017) 5, 745–749 749
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
751–758
DURABILITY OF MATERIALS BASED ON A
POLYMER-SILICATE MATRIX AND A LIGHTWEIGHT
AGGREGATE EXPOSED TO AGGRESSIVE INFLUENCES
COMBINED WITH HIGH TEMPERATURES
VZDR@LJIVOST MATERIALOV NA OSNOVI IZ
POLIMER-SILIKATNIH MATRIC IN LAHKEGA DODATKA,
IZPOSTAVLJENIH AGRESIVNIM VPLIVOM V KOMBINACIJI Z
VISOKIMI TEMPERATURAMI
Tomá{ Melichar, Jiøí Byd`ovský, Ámos Dufka
Brno University of Technology, Faculty of Civil Engineering, Veveøí 331/95, 602 00 Brno, Czech Republic
melichar.t@fce.vutbr.cz
Prejem rokopisa – received: 2016-07-15; sprejem za objavo – accepted for publication: 2017-03-16
doi:10.17222/mit.2016.190
The paper presents the results and findings of the research focused on an examination of progressive degradation of newly
developed composite materials. These materials consisted of a polymer-silicate matrix and a mix of fillers with a considerable
proportion of a porous aggregate. The matrix also contained a larger amount of alternative raw materials, in particular high-
temperature fly ash and blast furnace slag. Prepared mixes were tested after different periods of exposure to an aggressive
environment – for 45 d and 90 d (50 and 100 cycles). First, after 45 d and 90 d, reference materials (i.e., those not exposed to
aggressive influences) were tested. In parallel, tests were carried out with materials exposed to a solution of chloride ions and
subjected to cyclic freeze and thaw. The last mode of testing in this research included materials exposed to cyclic freeze and
thaw, chloride-ion solution and the consequent thermal load of up to 1000 °C. The level of degradation was evaluated by means
of physico-mechanical, physico-chemical and microstructural test methods. It was found that the exposure to cyclic freeze and
thaw, water and chloride ions has no considerable influence on the reduction of the thermal resistance of the developed mix
designs of polymer-silicate-based composites.
Keywords: durability, porous aggregate, polymer-silicate matrix, long-term exposure, chloride ions, water, frost, fire, micro-
structure
^lanek predstavlja rezultate in izsledke raziskave, osredoto~ene na preverjanje napredujo~ega slab{anja na novo razvitih kompo-
zitnih materialov. Ti materiali so sestavljeni iz polimer-silikatnih matric in me{anice polnil z znatnim dele`em poroznih
agregatov. Matrica je vsebovala tudi ve~jo koli~ino alternativnih surovih materialov, visokotemperaturni dimni{ki pepel in
`lindro. Pripravljene me{anice so bile testirane po razli~nih ~asih izpostavljanja agresivnemu okolju – 45 d in 90 d (v 50 in 100
ciklih). Najprej so bili, po 45 d in 90 d, testirani referen~ni materiali, tisti, ki niso bili izpostavljeni agresivnim vplivom.
Vzporedno pa so bila izvedena testiranja materialov, ki so bili izpostavljeni raztopini kloridnih ionov ter nato cikli~no
zamrznjeni in nato staljeni. Zadnji na~in testiranja v tej raziskavi je bilo izpostavljanje materialov cikli~nemu zamrzovanju in
taljenju, kloridni raztopini ionov in ter toplotni obremenitvi do 1000 °C. Stopnja degradacije je bila ocenjena glede na
fizikalno-mehanske, fizikalno-kemi~ne in mikrostrukturne testne metode. Ugotovljeno je bilo, da izpostavljenost cikli~nemu
zamrzovanje in taljenju, vodi in kloridnim ionom, nima ve~jega vpliva na zmanj{anje termi~ne odpornosti na razvite me{anice iz
polimer-silikatnih kompozitov.
Klju~ne besede: vzdr`ljivost, porozne me{anice, polimer-silikatna osnova, dalj{e izpostavljanje, kloridni ioni, voda, led, ogenj,
mikrostruktura
1 INTRODUCTION
Technical papers, stating results of the research
focused on the problems of the resistance of silicate
(cement) matrix based composites, focus mainly on an
assessment of the influences of individual adverse fac-
tors and aggressive environments. Scientists do not focus
on a synergic action of more than two types of adverse
influences, one of which is extremely high temperature.
Several factors are crucial for a gradual degradation of
polymer-silicate composites exposed to chloride ions,
frost and high temperature. The first of them is mecha-
nical deterioration of the composite caused by freeze-
thaw cycles together with the water containing chloride
ions. A water solution with chlorides penetrates through
the surface and subsequently deeper into the structure of
a composite material with a capillary-porous system.
Depending on the depth of the ingression of the men-
tioned solution, affected areas are damaged during the
temperatures changes above and below zero. The depth
of penetration depends on the components used, in parti-
cular, the material base of the matrix, the used additions,
their activities and the structure of the aggregate.
This problem is solved, in considerable detail, by J.
G. Jang et al.1 Apart from the influence of individual
factors, the authors also described a combination of
aggressive CO2 and Cl– ions, and it was proved that when
the two act in synergy, the penetration of Cl ions is
Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758 751
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:620.17:620.192.47 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)751(2017)
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
751–758
DURABILITY OF MATERIALS BASED ON A
POLYMER-SILICATE MATRIX AND A LIGHTWEIGHT
AGGREGATE EXPOSED TO AGGRESSIVE INFLUENCES
COMBINED WITH HIGH TEMPERATURES
VZDR@LJIVOST MATERIALOV NA OSNOVI IZ
POLIMER-SILIKATNIH MATRIC IN LAHKEGA DODATKA,
IZPOSTAVLJENIH AGRESIVNIM VPLIVOM V KOMBINACIJI Z
VISOKIMI TEMPERATURAMI
Tomá{ Melichar, Jiøí Byd`ovský, Ámos Dufka
Brno University of Technology, Faculty of Civil Engineering, Veveøí 331/95, 602 00 Brno, Czech Republic
melichar.t@fce.vutbr.cz
Prejem rokopisa – received: 2016-07-15; sprejem za objavo – accepted for publication: 2017-03-16
doi:10.17222/mit.2016.190
The paper presents the results and findings of the research focused on an examination of progressive degradation of newly
developed composite materials. These materials consisted of a polymer-silicate matrix and a mix of fillers with a considerable
proportion of a porous aggregate. The matrix also contained a larger amount of alternative raw materials, in particular high-
temperature fly ash and blast furnace slag. Prepared mixes were tested after different periods of exposure to an aggressive
environment – for 45 d and 90 d (50 and 100 cycles). First, after 45 d and 90 d, reference materials (i.e., those not exposed to
aggressive influences) were tested. In parallel, tests were carried out with materials exposed to a solution of chloride ions and
subjected to cyclic freeze and thaw. The last mode of testing in this research included materials exposed to cyclic freeze and
thaw, chloride-ion solution and the consequent thermal load of up to 1000 °C. The level of degradation was evaluated by means
of physico-mechanical, physico-chemical and microstructural test methods. It was found that the exposure to cyclic freeze and
thaw, water and chloride ions has no considerable influence on the reduction of the thermal resistance of the developed mix
designs of polymer-silicate-based composites.
Keywords: durability, porous aggregate, polymer-silicate matrix, long-term exposure, chloride ions, water, frost, fire, micro-
structure
^lanek predstavlja rezultate in izsledke raziskave, osredoto~ene na preverjanje napredujo~ega slab{anja na novo razvitih kompo-
zitnih materialov. Ti materiali so sestavljeni iz polimer-silikatnih matric in me{anice polnil z znatnim dele`em poroznih
agregatov. Matrica je vsebovala tudi ve~jo koli~ino alternativnih surovih materialov, visokotemperaturni dimni{ki pepel in
`lindro. Pripravljene me{anice so bile testirane po razli~nih ~asih izpostavljanja agresivnemu okolju – 45 d in 90 d (v 50 in 100
ciklih). Najprej so bili, po 45 d in 90 d, testirani referen~ni materiali, tisti, ki niso bili izpostavljeni agresivnim vplivom.
Vzporedno pa so bila izvedena testiranja materialov, ki so bili izpostavljeni raztopini kloridnih ionov ter nato cikli~no
zamrznjeni in nato staljeni. Zadnji na~in testiranja v tej raziskavi je bilo izpostavljanje materialov cikli~nemu zamrzovanju in
taljenju, kloridni raztopini ionov in ter toplotni obremenitvi do 1000 °C. Stopnja degradacije je bila ocenjena glede na
fizikalno-mehanske, fizikalno-kemi~ne in mikrostrukturne testne metode. Ugotovljeno je bilo, da izpostavljenost cikli~nemu
zamrzovanje in taljenju, vodi in kloridnim ionom, nima ve~jega vpliva na zmanj{anje termi~ne odpornosti na razvite me{anice iz
polimer-silikatnih kompozitov.
Klju~ne besede: vzdr`ljivost, porozne me{anice, polimer-silikatna osnova, dalj{e izpostavljanje, kloridni ioni, voda, led, ogenj,
mikrostruktura
1 INTRODUCTION
Technical papers, stating results of the research
focused on the problems of the resistance of silicate
(cement) matrix based composites, focus mainly on an
assessment of the influences of individual adverse fac-
tors and aggressive environments. Scientists do not focus
on a synergic action of more than two types of adverse
influences, one of which is extremely high temperature.
Several factors are crucial for a gradual degradation of
polymer-silicate composites exposed to chloride ions,
frost and high temperature. The first of them is mecha-
nical deterioration of the composite caused by freeze-
thaw cycles together with the water containing chloride
ions. A water solution with chlorides penetrates through
the surface and subsequently deeper into the structure of
a composite material with a capillary-porous system.
Depending on the depth of the ingression of the men-
tioned solution, affected areas are damaged during the
temperatures changes above and below zero. The depth
of penetration depends on the components used, in parti-
cular, the material base of the matrix, the used additions,
their activities and the structure of the aggregate.
This problem is solved, in considerable detail, by J.
G. Jang et al.1 Apart from the influence of individual
factors, the authors also described a combination of
aggressive CO2 and Cl– ions, and it was proved that when
the two act in synergy, the penetration of Cl ions is
Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758 751
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:620.17:620.192.47 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)751(2017)
increased. This implies that the testing of a synergic
action of negative influences affecting a given material is
quite important because it is impossible to give an exact
prediction only on the basis of an evaluation of indi-
vidual aggressive substances. M. Maes et al.2 confirmed
that a replacement of Portland cement with blast-furnace
slag in amounts of 50 % and 70 % considerably in-
creases the resistance to the actions of Cl– and SO4–2. It is
also interesting that increasing the contents of sulfates in
a solution increases the attack of chlorides. On the other
hand, the concentration of chlorides has no influence on
the sulfate attack on the mortar with blast-furnace slag.
Slag generally has good resistance to sulfate attack,
which was proved by M. Maes et al.2, inter alia, with
negligible changes in the volume and expansion. This
finding is quite important from the viewpoint of a posi-
tive influence of blast-furnace slag on the resistance of a
cement matrix to extreme temperatures.
The presence of Cl– ions in silicates is also important
for a possible formation of Friedel’s salt – 3CaO·Al2O3·
·CaCl2·10H2O. This chemical compound can develop
from free Cl– ions, which bind to C3A or C4AF. However,
the reaction of the mentioned components is consider-
ably influenced by several factors. As A. Delagrave et
al.3 and D. Izquierdo et al.4, Q. Yuan5 and L. Tang6 state,
concentration of chloride ions, type of cement,
replacement of cement, temperature or water-cement
ratio are important. According to F. P. Glasser et al.7, a
penetration of chloride ions with the consequent
formation of Friedel’s salt can, inter alia, cause a gradual
filling of the porous structure of a given composite,
which eventually slows down the transportation of Cl–
ions. The filling of pores is critical particularly for the
thermal resistance where a porous structure is an
advantage. The thermal decomposition of Friedel’s salt is
also important, i.e., gases developed during the
dissociation, components remaining as residua and their
capability for a further reaction with water consequently
lead to an increase in the volume.
Durability of mortars with blast-furnace slag as the
filler is described by Santamaría-Vicario et al.8 The
research focused on the influence of a combination of
several adverse environments – frost with water, alter-
nating between humid and dry environment, simulation
of a marine environment or crystallization of salts and
gaseous SO2. Based on the results and findings, the
possibility of using slag as a replacement of the aggre-
gate was confirmed. The development and research of
mortars with blast-furnace slag resistant to fire up to a
temperature of 900 °C is presented by S. Aydin9. The
tests involving blast-furnace slag as a replacement of the
binder included mortars containing up to 80 % of the
slag. Pumice was used as the filler. Mortars with 40 % of
the slag showed the best test results. The residual com-
pressive strength of these mortars exposed to 900 °C was
44 %, after they were gradually cooled down. After rapid
cooling with water, the residual compressive strength of
the mortar was around 40 %. The replacement of a fine
aggregate with industrial by-products is a part of the
research described by I. Yüksel et al.10 The amount of
the substitution was from 10–50 %, with the subsequent
thermal load of up to 800 °C at the age of 90 d. A posi-
tive influence of the replacement of the commonly used
dense aggregate with byproducts of the energy and
foundry industry was unambiguously proved. The
bottom ash showed the best results. One of the few
publications, in which the authors provide the results of a
research focused on the synergic action of high
temperature and adverse environment (in particular
CO2), was published by G. Yuan and Q. Li11. The authors
focus on the problems of concrete with various
water-cement ratios exposed to temperatures of up to 700
°C. One of the subsequent testing methods was the
exposure to an environment with a higher concentration
of CO2. The aim of the research presented by G. Yuan
and Q. Li11 was the verification of the influence of a
possible surface treatment of the tested concrete after its
exposure to heat. The results show that as the exposure
temperature grows, the resistance to carbonization
decreases.
2 METHODOLOGY OF EXPERIMENTAL WORK
As the aim of the research was to design a material
resistant to adverse environmental influences and
possible fire at the same time, the selected matrix
contained an increased amount of blast-furnace slag and
high-temperature fly ash. Fly ash as a substitution binder
increasing the thermal resistance was researched by A.
Nadeem et al.12 Microsilica was used as a supporting
addition and stabilizing component; its positive effect in
cement-based composites on the resistance to high
temperatures was proven, for example, by T. Harun and
C. Ahmet.13 Two types of aggregate were used. The
coarser fraction of 1–2 mm was amphibolite, which is
mined as a primary raw material. The fine-grained
residue from washing the aggregate, produced as a
by-product of adjusting the granulometric composition
of the mined amphibolite, was used. The remaining part
– a fine fraction of 0–1 mm – was fly-ash agloporite; the
dominant part of agloporite is an energy-industry
by-product – fly ash. As agloporite is created, through a
self-burning process, from alternative raw-material
resources, it can be characterized as an advantageous and
cheap raw material, which is relatively environmentally
friendly. The suitability of agloporite as a filler for
materials with an assumed thermal resistance can be
presumed, for example, on the basis of the results stated
by V. ^erný and [. Keprdová.14 Mix designs of tested
composite materials are listed in Table 1. Agloporite was
saturated with water before its application in the mix.
This fact was taken into account when correcting the
water-cement ratio, which was adjusted to achieve good
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
752 Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
workability of the fresh mix (a slump of around 160
mm).
Out of each mix design, four sets of test specimens of
(40 × 40 × 160) mm were made. The first and the second
sets were treated in standardized conditions; then, after
28 d of maturing, the specimens were exposed to an
adverse environment. In the adverse environment, the
temperature was cyclically changing (above and below
zero), combined with a saturation of the solution with
chloride ions (in accordance with ^SN EN17). After 50
cycles (and after 45 d elapsed from the standardized
maturation), half of the specimens were subjected to
tests of strength characteristics. The other half of the test
specimens were exposed to temperatures of (22, 400,
600 and 1000) °C. The rate of the temperature increase
was around 10 °C min–1 with an isothermal dwell time of
90 min and slow cooling down. A similar procedure was
also applied for the third and fourth set, where the test
conditions were different regarding the number of
cycles: 100 cycles were made in an adverse environment
so that the specimens could be tested after 90 d. It is
important to emphasize that the testing was carried out
after 45 d or 90 d; however, the cycles were counted after
the standardized age of 28 d. The reason for this was the
exposure of the test specimens to the adverse envi-
ronment after the maturing process. The end of the
cyclical exposure at the ages of 45 d or 90 d simulated
the early age of a real structure.
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758 753
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Bending tensile strengths of mixtures MBS and MFA after
45 d and 90 d (R – temperature exposure; Cl – combination of
chloride ions, frost and exposure to extreme temperatures)
Figure 2: Percentage change of compressive strength due to
chloride-ion solution and frost attack – weight change of mixtures
MBS and MFA after 45 d and 90 d
Figure 3: Change of temperature resistance (compressive strength)
due to chloride-ion solution and frost attack – weight change of mix-
tures MBS and MFA after 45 d and 90 d
Figure 1: Compressive strengths of mixtures MBS and MFA after 45
d and 90 d (R – temperature exposure; Cl – combination of chloride
ions, frost and exposure to extreme temperatures)
Table 1: Composition of mixtures MBS and MFA
Component Unit Mixture
MBS MFA
Cement kg m–3 443 443
Blast furnace slag kg m–3 251 –
Fly ash kg m–3 – 238
Vinyl acetate copolymer % (wC+FBS, FA) 3 3
Microsillica % (wC+FBS, FA) 7 7
Amphibolite (by-product)
<0,063 mm kg m
–3 33 33
Porous aggregate 0–1 mm kg m–3 641 641
Amphibolite 1–2 mm kg m–3 792 792
Polypropylene fibres kg m–-3 1.1 1.1
Water kg m–3 193 201
increased. This implies that the testing of a synergic
action of negative influences affecting a given material is
quite important because it is impossible to give an exact
prediction only on the basis of an evaluation of indi-
vidual aggressive substances. M. Maes et al.2 confirmed
that a replacement of Portland cement with blast-furnace
slag in amounts of 50 % and 70 % considerably in-
creases the resistance to the actions of Cl– and SO4–2. It is
also interesting that increasing the contents of sulfates in
a solution increases the attack of chlorides. On the other
hand, the concentration of chlorides has no influence on
the sulfate attack on the mortar with blast-furnace slag.
Slag generally has good resistance to sulfate attack,
which was proved by M. Maes et al.2, inter alia, with
negligible changes in the volume and expansion. This
finding is quite important from the viewpoint of a posi-
tive influence of blast-furnace slag on the resistance of a
cement matrix to extreme temperatures.
The presence of Cl– ions in silicates is also important
for a possible formation of Friedel’s salt – 3CaO·Al2O3·
·CaCl2·10H2O. This chemical compound can develop
from free Cl– ions, which bind to C3A or C4AF. However,
the reaction of the mentioned components is consider-
ably influenced by several factors. As A. Delagrave et
al.3 and D. Izquierdo et al.4, Q. Yuan5 and L. Tang6 state,
concentration of chloride ions, type of cement,
replacement of cement, temperature or water-cement
ratio are important. According to F. P. Glasser et al.7, a
penetration of chloride ions with the consequent
formation of Friedel’s salt can, inter alia, cause a gradual
filling of the porous structure of a given composite,
which eventually slows down the transportation of Cl–
ions. The filling of pores is critical particularly for the
thermal resistance where a porous structure is an
advantage. The thermal decomposition of Friedel’s salt is
also important, i.e., gases developed during the
dissociation, components remaining as residua and their
capability for a further reaction with water consequently
lead to an increase in the volume.
Durability of mortars with blast-furnace slag as the
filler is described by Santamaría-Vicario et al.8 The
research focused on the influence of a combination of
several adverse environments – frost with water, alter-
nating between humid and dry environment, simulation
of a marine environment or crystallization of salts and
gaseous SO2. Based on the results and findings, the
possibility of using slag as a replacement of the aggre-
gate was confirmed. The development and research of
mortars with blast-furnace slag resistant to fire up to a
temperature of 900 °C is presented by S. Aydin9. The
tests involving blast-furnace slag as a replacement of the
binder included mortars containing up to 80 % of the
slag. Pumice was used as the filler. Mortars with 40 % of
the slag showed the best test results. The residual com-
pressive strength of these mortars exposed to 900 °C was
44 %, after they were gradually cooled down. After rapid
cooling with water, the residual compressive strength of
the mortar was around 40 %. The replacement of a fine
aggregate with industrial by-products is a part of the
research described by I. Yüksel et al.10 The amount of
the substitution was from 10–50 %, with the subsequent
thermal load of up to 800 °C at the age of 90 d. A posi-
tive influence of the replacement of the commonly used
dense aggregate with byproducts of the energy and
foundry industry was unambiguously proved. The
bottom ash showed the best results. One of the few
publications, in which the authors provide the results of a
research focused on the synergic action of high
temperature and adverse environment (in particular
CO2), was published by G. Yuan and Q. Li11. The authors
focus on the problems of concrete with various
water-cement ratios exposed to temperatures of up to 700
°C. One of the subsequent testing methods was the
exposure to an environment with a higher concentration
of CO2. The aim of the research presented by G. Yuan
and Q. Li11 was the verification of the influence of a
possible surface treatment of the tested concrete after its
exposure to heat. The results show that as the exposure
temperature grows, the resistance to carbonization
decreases.
2 METHODOLOGY OF EXPERIMENTAL WORK
As the aim of the research was to design a material
resistant to adverse environmental influences and
possible fire at the same time, the selected matrix
contained an increased amount of blast-furnace slag and
high-temperature fly ash. Fly ash as a substitution binder
increasing the thermal resistance was researched by A.
Nadeem et al.12 Microsilica was used as a supporting
addition and stabilizing component; its positive effect in
cement-based composites on the resistance to high
temperatures was proven, for example, by T. Harun and
C. Ahmet.13 Two types of aggregate were used. The
coarser fraction of 1–2 mm was amphibolite, which is
mined as a primary raw material. The fine-grained
residue from washing the aggregate, produced as a
by-product of adjusting the granulometric composition
of the mined amphibolite, was used. The remaining part
– a fine fraction of 0–1 mm – was fly-ash agloporite; the
dominant part of agloporite is an energy-industry
by-product – fly ash. As agloporite is created, through a
self-burning process, from alternative raw-material
resources, it can be characterized as an advantageous and
cheap raw material, which is relatively environmentally
friendly. The suitability of agloporite as a filler for
materials with an assumed thermal resistance can be
presumed, for example, on the basis of the results stated
by V. ^erný and [. Keprdová.14 Mix designs of tested
composite materials are listed in Table 1. Agloporite was
saturated with water before its application in the mix.
This fact was taken into account when correcting the
water-cement ratio, which was adjusted to achieve good
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
752 Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
workability of the fresh mix (a slump of around 160
mm).
Out of each mix design, four sets of test specimens of
(40 × 40 × 160) mm were made. The first and the second
sets were treated in standardized conditions; then, after
28 d of maturing, the specimens were exposed to an
adverse environment. In the adverse environment, the
temperature was cyclically changing (above and below
zero), combined with a saturation of the solution with
chloride ions (in accordance with ^SN EN17). After 50
cycles (and after 45 d elapsed from the standardized
maturation), half of the specimens were subjected to
tests of strength characteristics. The other half of the test
specimens were exposed to temperatures of (22, 400,
600 and 1000) °C. The rate of the temperature increase
was around 10 °C min–1 with an isothermal dwell time of
90 min and slow cooling down. A similar procedure was
also applied for the third and fourth set, where the test
conditions were different regarding the number of
cycles: 100 cycles were made in an adverse environment
so that the specimens could be tested after 90 d. It is
important to emphasize that the testing was carried out
after 45 d or 90 d; however, the cycles were counted after
the standardized age of 28 d. The reason for this was the
exposure of the test specimens to the adverse envi-
ronment after the maturing process. The end of the
cyclical exposure at the ages of 45 d or 90 d simulated
the early age of a real structure.
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758 753
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Bending tensile strengths of mixtures MBS and MFA after
45 d and 90 d (R – temperature exposure; Cl – combination of
chloride ions, frost and exposure to extreme temperatures)
Figure 2: Percentage change of compressive strength due to
chloride-ion solution and frost attack – weight change of mixtures
MBS and MFA after 45 d and 90 d
Figure 3: Change of temperature resistance (compressive strength)
due to chloride-ion solution and frost attack – weight change of mix-
tures MBS and MFA after 45 d and 90 d
Figure 1: Compressive strengths of mixtures MBS and MFA after 45
d and 90 d (R – temperature exposure; Cl – combination of chloride
ions, frost and exposure to extreme temperatures)
Table 1: Composition of mixtures MBS and MFA
Component Unit Mixture
MBS MFA
Cement kg m–3 443 443
Blast furnace slag kg m–3 251 –
Fly ash kg m–3 – 238
Vinyl acetate copolymer % (wC+FBS, FA) 3 3
Microsillica % (wC+FBS, FA) 7 7
Amphibolite (by-product)
<0,063 mm kg m
–3 33 33
Porous aggregate 0–1 mm kg m–3 641 641
Amphibolite 1–2 mm kg m–3 792 792
Polypropylene fibres kg m–-3 1.1 1.1
Water kg m–3 193 201
3 RESULTS
The diagrams below show the dependency of indivi-
dual observed parameters on the exposure to aggressive
environments and high temperatures, including their
combination (Figures 1 to 6). The first compressive
strength and tensile bending strength were examined.
The first diagram of a given set – observed parameters
(Figures 1 to 4) – shows the development of the above-
mentioned parameters in a given environment and at a
given temperature (labeling: R – with no aggressive ex-
posure, Cl – exposed to a solution of chlorides and frost).
The second diagram of a given set shows the percentage
changes of a particular parameter (Figures 2 and 5). The
last diagram of a given parameter (Figures 3 and 6)
shows the development of the temperature resistance
(which is characterized by the percentage decrease of a
given parameter). The curves presented as a solid line
(dash line) denote mix designs MBS (MFA), i.e., the
binder modified with blast-furnace slag (high-tempe-
rature fly ash).
Due to the chloride-ion solution and frost, failures are
possible even inside the composite. The formation of
cracks is the most likely failure. One of the key methods
of the assessment of the structure of a material exposed
to high temperatures is X-ray tomography (and CT). The
CT examination focused on the presence of cracks, in
particular, their number, width and orientation. It was
important to determine where the cracks were found: the
matrix, the aggregate or the transition zone. The location
of cracks is also important: close to the surface or in the
inner structure. Figures 7 to 10 show the test specimens
exposed to adverse conditions. To identify and evaluate
phase changes, an X-ray diffraction analysis was used
(hereinafter referred to as XRD). The observed para-
meters were phase changes and the formation of new
chemical compounds as products of the reaction of
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
754 Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: CT picture – slice of MBS-45/Cl (45 d, Cl – exposure to
chloride solution and frost environment, 50 cycles), exposure to
1000 °C
Figure 6: Change of temperature resistance (bending tensile strength)
due to chloride-ion solution and frost attack – weight change of
mixtures MBS and MFA after 45 d and 90 d
Figure 7: CT picture – slice of MBS-45/R (45 d, R – reference sample
without chloride solution and frost environment), exposure to 1000 °C
Figure 5: Percentage change of bending tensile strength due to chlo-
ride-ion solution and frost attack – weight change of mixtures MBS
and MFA after 45 d and 90 d
chlorides with the matrix of a composite (Table 2). To
make the picture complete, electron microscopy was
used (hereinafter referred to as SEM). SEM focused on
the formation of micro-cracks and the products related to
the crystallization of chloride ions or their residua
(Figures 11 and 12).
Table 2: XRD – mineralogical composition of mixtures MBS and
MFA – selected samples
Mixture-age/
environment,
temperature
Identified components
MBS-90/R,
22
portlandite, calcite, -quartz, ettringite, mullite,
chlorite, actinolite, anorthite, hematite,
akermanite, merwinite, gehlenite, CSH,
amorphous phase
MBS-90/R,
1000
-quartz, mullite, chlorite, actinolite, anorthite,
hematite, akermanite, merwinite, wolastonite,
amorphous phase
MBS-90/Cl,
22
portlandite, calcite,  -quartz, ettringite,
mullite, chlorite, hematite, akermanite,
actinolite, wolastonite, gehlenite, biotite CSH,
amorphous phase
MBS-90/Cl,
1000
-quartz, mullite, actinolite, anorthite, hematite,
gehlenite, amorphous phase wolastonite
MFA-90/R,
22
portlandite, calcite, -quartz, ettringite, mullite,
chlorite, actinolite, anorthite, hematite, biotite
CSH, amorphous phase
MFA-90/R,
1000
-quartz, mullite, chlorite, actinolite,
amorphous phase, anorthite, hematite,
wolastonite
MFA-90/Cl,
22
portlandite, ettringite, mullite, calcite, -quartz,
anorthite, biotite, hematite CSH, amorphous
phase
MFA-90/Cl,
1000
-quartz, mullite, amorphous phase, anorthite,
hematite, wolastonite
4 DISCUSSION
The diagrams (Figures 1 to 6) show the development
of individual parameters and differences between the
tested mix designs at a given age and exposed to a given
environment.
Reference materials MBS-90 and MFA-90 showed
compressive strengths of 48 N mm–2 and 45 N mm–2. As
the diagram shows (Figure 1), the compressive strength
decreases more considerably after an exposure to an
aggressive environment. This fact is supported by the test
results of the materials stored in laboratory conditions (in
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758 755
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 11: Microstructure of MBS-45/Cl (45 d, Cl – exposure to
chloride solution and frost environment, 50 cycles), exposure to
1000 °C, 5000× magnification
Figure 9: CT picture – slice of MBS-90/Cl (90 d, Cl – exposure to
chloride solution and frost environment, 100 cycles), exposure to
1000 °C
Figure 10: CT picture – slice of MFA-90/Cl (90 d, Cl – exposure to
chloride solution and frost environment, 100 cycles), exposure to
1000 °C
3 RESULTS
The diagrams below show the dependency of indivi-
dual observed parameters on the exposure to aggressive
environments and high temperatures, including their
combination (Figures 1 to 6). The first compressive
strength and tensile bending strength were examined.
The first diagram of a given set – observed parameters
(Figures 1 to 4) – shows the development of the above-
mentioned parameters in a given environment and at a
given temperature (labeling: R – with no aggressive ex-
posure, Cl – exposed to a solution of chlorides and frost).
The second diagram of a given set shows the percentage
changes of a particular parameter (Figures 2 and 5). The
last diagram of a given parameter (Figures 3 and 6)
shows the development of the temperature resistance
(which is characterized by the percentage decrease of a
given parameter). The curves presented as a solid line
(dash line) denote mix designs MBS (MFA), i.e., the
binder modified with blast-furnace slag (high-tempe-
rature fly ash).
Due to the chloride-ion solution and frost, failures are
possible even inside the composite. The formation of
cracks is the most likely failure. One of the key methods
of the assessment of the structure of a material exposed
to high temperatures is X-ray tomography (and CT). The
CT examination focused on the presence of cracks, in
particular, their number, width and orientation. It was
important to determine where the cracks were found: the
matrix, the aggregate or the transition zone. The location
of cracks is also important: close to the surface or in the
inner structure. Figures 7 to 10 show the test specimens
exposed to adverse conditions. To identify and evaluate
phase changes, an X-ray diffraction analysis was used
(hereinafter referred to as XRD). The observed para-
meters were phase changes and the formation of new
chemical compounds as products of the reaction of
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
754 Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: CT picture – slice of MBS-45/Cl (45 d, Cl – exposure to
chloride solution and frost environment, 50 cycles), exposure to
1000 °C
Figure 6: Change of temperature resistance (bending tensile strength)
due to chloride-ion solution and frost attack – weight change of
mixtures MBS and MFA after 45 d and 90 d
Figure 7: CT picture – slice of MBS-45/R (45 d, R – reference sample
without chloride solution and frost environment), exposure to 1000 °C
Figure 5: Percentage change of bending tensile strength due to chlo-
ride-ion solution and frost attack – weight change of mixtures MBS
and MFA after 45 d and 90 d
chlorides with the matrix of a composite (Table 2). To
make the picture complete, electron microscopy was
used (hereinafter referred to as SEM). SEM focused on
the formation of micro-cracks and the products related to
the crystallization of chloride ions or their residua
(Figures 11 and 12).
Table 2: XRD – mineralogical composition of mixtures MBS and
MFA – selected samples
Mixture-age/
environment,
temperature
Identified components
MBS-90/R,
22
portlandite, calcite, -quartz, ettringite, mullite,
chlorite, actinolite, anorthite, hematite,
akermanite, merwinite, gehlenite, CSH,
amorphous phase
MBS-90/R,
1000
-quartz, mullite, chlorite, actinolite, anorthite,
hematite, akermanite, merwinite, wolastonite,
amorphous phase
MBS-90/Cl,
22
portlandite, calcite,  -quartz, ettringite,
mullite, chlorite, hematite, akermanite,
actinolite, wolastonite, gehlenite, biotite CSH,
amorphous phase
MBS-90/Cl,
1000
-quartz, mullite, actinolite, anorthite, hematite,
gehlenite, amorphous phase wolastonite
MFA-90/R,
22
portlandite, calcite, -quartz, ettringite, mullite,
chlorite, actinolite, anorthite, hematite, biotite
CSH, amorphous phase
MFA-90/R,
1000
-quartz, mullite, chlorite, actinolite,
amorphous phase, anorthite, hematite,
wolastonite
MFA-90/Cl,
22
portlandite, ettringite, mullite, calcite, -quartz,
anorthite, biotite, hematite CSH, amorphous
phase
MFA-90/Cl,
1000
-quartz, mullite, amorphous phase, anorthite,
hematite, wolastonite
4 DISCUSSION
The diagrams (Figures 1 to 6) show the development
of individual parameters and differences between the
tested mix designs at a given age and exposed to a given
environment.
Reference materials MBS-90 and MFA-90 showed
compressive strengths of 48 N mm–2 and 45 N mm–2. As
the diagram shows (Figure 1), the compressive strength
decreases more considerably after an exposure to an
aggressive environment. This fact is supported by the test
results of the materials stored in laboratory conditions (in
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758 755
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 11: Microstructure of MBS-45/Cl (45 d, Cl – exposure to
chloride solution and frost environment, 50 cycles), exposure to
1000 °C, 5000× magnification
Figure 9: CT picture – slice of MBS-90/Cl (90 d, Cl – exposure to
chloride solution and frost environment, 100 cycles), exposure to
1000 °C
Figure 10: CT picture – slice of MFA-90/Cl (90 d, Cl – exposure to
chloride solution and frost environment, 100 cycles), exposure to
1000 °C
diagram "22"). A combination of frost and chloride ions
caused a decrease in the strength of 27.5 % (MBS-90Cl)
and 29.6 % (MFA-90Cl). These values can be assessed
as adequate for the given type of material (i.e., the
material with a porous aggregate) after a given number
of cycles. The thermal resistance (a relative decrease in
the compressive strength) of the materials in individual
environments after being exposed to 1000 °C was
approximately 64–69 % (72–78 %) for the reference
samples (the samples exposed to the aggressive envi-
ronment). It is interesting that slag (i.e., MBS) showed
better results than the other reference material; however,
after the exposure to the aggressive environment and
high temperature, fly ash was better (i.e., MFA). For
illustration, MBS-90/Cl (MFA-90/Cl) showed a decrease
in the compressive strength of 75.1 % (74.1 %; Figure
2). A slight increase in the thermal resistance developed
over time in spite of the higher number of cycles in the
aggressive environment. The only exception was
MFA-90/Cl. An evaluation of the relative change in the
thermal resistance caused by the aggressive environment
shows an irregular development of the curves (Figure 3).
This is particularly clear for temperatures of 600 °C and
above. It is evident that the exposure to an aggressive
environment and 1000 °C caused a decrease in the
observed parameter of MBS (MFA) of approximately
11 % (2.5–7.2 %).
Further, the bending tensile strength was evaluated.
The differences are less obvious compared to the
development of the compressive strength (Figures 1 and
4). The increase in the bending strength in time was also
less considerable than the increase in the compressive
strength. The bending strengths of the test specimens not
exposed to aggressive environments and a thermal load
were from 7.3 N mm–2 to 7.6 N mm–2 (6.9 N mm–2 to
7.2 N mm–2) for mix designs MBS (MFA). After the ex-
posure of the samples to the aggressive environment with
chloride ions, the decrease was 12.6 % for MBS-90 and
13.9 % for MFA-90.
The decrease in the tensile bending strength caused
by the exposure to aggressive environments is then much
lower than that of the compressive strength. The negative
impact on the thermal resistance was also smaller. The
decrease in the tensile bending strength was 65.8 %
(62.3 %) for MBS-90/R (MFA-90/R) after the exposure
to 1000 °C (Figure 5). These facts imply that the substi-
tution of the primary binder with fly ash had a more
positive effect from the viewpoint of a combined ex-
posure involving an aggressive environment and a ther-
mal load. On the other hand, comparing the changes in
the thermal resistance in time and after the exposure to
adverse environments shows (Figure 6) that blast-fur-
nace slag had better results (MBS). The reduction in the
thermal resistance after the exposure to 1000 °C was
5.3 % for MBS-90/Cl-R and 6.1 % for MFA/Cl-R. The
reduction in the thermal resistance of MFA after the
exposure to 600 °C is also interesting for both tested
ages. It can be assumed that used high-temperature fly
ash contains components that react with chloride
solutions, possibly only during the thermal load. It could
also be a synergy reaction of both mentioned influences.
The synergy effect of an aggressive environment and
extreme temperatures causes a considerable degradation
of the structures of the materials based on a silicate
matrix and a porous aggregate. The given case is con-
firmed with the results of physico-mechanical tests of
materials MBS and MFA. The CT examination proved
the above character of the crack changes after the expo-
sure to the aggressive environment (chlorides and frost)
and the subsequent extreme thermal load. On the refe-
rence materials, only a small amount of randomly
oriented cracks was identified (Figure 7). The cracks
were rather small, i.e., with a smaller width. These
cracks mostly did not reach the transition zone between
the matrix and dense aggregate. On the contrary, the
number of cracks of the materials exposed to the aggres-
sive environment was larger (Figures 8 to 10). An
increase in the width and number of the cracks is evident
as the number of aggressive cycles and age of the sam-
ples grow. However, the increase was not considerable.
The orientation of the cracks was easily observable; the
cracks pointed towards the center of a test specimen
(Figures 8 and 9). The width and number of the cracks
and more cracks in the transit zone between the dense
aggregate and the matrix were evident compared to the
reference materials.
The last representative sample was FMA-60/Cl (Fig-
ure 10). It is interesting that its cracks are more close to
the surface, with an orientation similar to that of MBS.
Chloride ions combined with frost and the subsequent
thermal shock damaged more considerably the surface
areas of MFA, whereas MBS was damaged more evenly
throughout the whole structure.
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
756 Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 12: Microstructure of MBS-90/Cl (90 days, Cl – exposure to
chloride solution and frost environment, 100 cycles), exposure to
1000 °C, 5000× magnification
The XRD analysis was used for the identification of
crystal phases and their changes. XRD did not identify
Friedel’s salt in any of the tested specimens. The crucial
factor of the formation of this chemical compound is the
temperature, which is, inter alia, mentioned by Dela-
grave et al.3, Izquierdo et al.4, Yuan5 and Tang6. It is then
evident that the damage of the structure caused by the
combination of frost and the solution with chloride ions
is purely mechanical. The salt in the solution – NaCl – is
transported into the porous structure of the composite by
equalizing the concentration gradient. Sudden changes of
the temperatures above and below zero then cause the
crystallization of the salt and the change of the phase of
water, which necessarily causes a decrease in the me-
chanical properties of the given material. This corres-
ponds with the findings determined with CT. As regards
the changes in the structures of the analyzed materials, a
decomposition of chemical compounds (in particular
mineralogical phases) typical for silicates and the used
aggregate, due to the increasing temperature, was ob-
served. The development shows peaks with a decreasing
intensity and a gradual ceasing of portlandite, calcite and
CSH phases. At temperatures over 600 °C, mew phases
were observed (like -C2S).
SEM together with an element analysis showed the
presence of Cl– ions in some areas. Microcracks and the
degree of decomposition of the hydration products of the
matrix or the changes in the structure of the used aggre-
gate were marked, as shown in Figures 11 and 12.
Almost no failures were observable in the transit zone
between the porous aggregate (agloporite) and the
matrix. As regards the age of the tested materials/number
of cycles, no considerable differences in the microstruc-
ture were observed. A 5000× magnification clearly
shows a degraded polymer-silicate matrix. SEM images
clearly show spherical particles of fly ash also in MBS.
This was caused by an imperfect manufacturing process
of agloporite, during which not all granules (for the
production of the aggregate) were perfectly compacted
and fired. However, as this fly ash had gone through the
thermic process twice, it is almost an inert filler from the
viewpoint of high temperature. Only few microscopic
cracks were observed in the matrix.
5 CONCLUSIONS
The influence of a combination of an aggressive
environment (chloride solution and freeze-thaw cycles)
and a thermal load (of up to 1000 °C) was tested and
analyzed in detail on newly designed materials as it is
not a thoroughly examined area. Fly-ash agloporite is a
promising aggregate and its influence on polymer-sili-
cate materials is still not quite clear. The advantage of
agloporite is its thermal resistance and very good
interaction with a polymer-silicate matrix. However, an
exposure to cyclic changing of temperatures above and
below zero, intensified by a salt solution, makes the use
of the porous aggregate difficult. Blast-furnace slag
(MBS) and fly ash (MFA) were used for a modification
of the matrix composition. Physico-chemical and micro-
structural analyses proved that the action of the
above-mentioned environment disturbs the structures of
MBS and MFA in a mechanical way, causing an expan-
sion due to the change of the phase of water or crystalli-
zation of salt rather than the formation of new chemical
compounds. The extreme temperature then only inten-
sifies these phenomena by degrading the structures, in
particular that of the matrix. It was proved that CT is
crucial for the explanation of the changes/failures of the
inner structure. This analytical method is quite suitable
for examining the three-dimensional structure of a given
material. It is particularly applicable for revealing va-
rious cracks, their orientation, localization or number.
The tested materials have a large potential for future
research. An examination of the behavior of the mate-
rials over a longer time (e.g., 180 d and 360 d) would
also be very interesting.
Acknowledgements
This paper has been worked out with the financial
support of The Czech Science Foundation (GA^R),
project GA15-07657S, "Study of kinetics of processes
occurring in the composite system at extreme tem-
peratures and exposed to an aggressive environment", as
well as under project No. LO1408 "AdMaS UP –
Advanced Materials, Structures and Technologies",
supported by the Ministry of Education, Youth and
Sports under the National Sustainability Programme I.
6 REFERENCES
1 J. G. Jang, H. J. Kim, H. K. Kim, H. K. Lee, Resistance of coal
bottom ash mortar against the coupled deterioration of carbonation
and chloride penetration, Materials & Design, 93 (2016), 160–167,
doi:10.1016/j.matdes.2015.12.074
2 M. Maes, N. De Belie, Resistance of concrete and mortar against
combined attack of chloride and sodium sulphate, Cement and Con-
crete Composites, 53 (2014), 59–72, doi:10.1016/j.cemconcomp.
2014.06.013
3 A. Delagrave, J. Marchand, J. P. Ollivier, S. Julien, K. Hazrati,
Chloride binding capacity of various hydrated cement paste systems,
Advanced Cement Based Materials, 6 (1997), 28–35, doi:
10.1016/S1065-7355(97)90003-1
4 D. Izquierdo, C. Alonso, C. Andrade, M. Castellote, Potentiostatic
determination of chloride threshold values for rebar depassivation –
experimental and statistical study, Electrochim. Acta, 49 (2004)
17–18, 2731–2739, doi:10.1016/j.electacta.2004.01.034
5 Q. Yuan, Fundamental studies on test methods for the transport of
chloride ions in cementitious materials, Doctoral Thesis, Ghent,
Ghent University, 2009
6 L. Tang, Chloride transport in concrete-measurement and prediction,
Doctoral Thesis, Göteborg, Chalmers University of Technology,
1996
7 F. P. Glasser, J. Marchand, E. Samson, Durability of concrete –
degradation phenomena involving detrimental chemical reactions,
Cem. Concr. Res, 38 (2007) 2, 226–246, doi:10.1016/j.cemconres.
2007.09.015
8 I. Santamaría-Vicario, A. Rodríguez, C. Junco, S. Gutiérrez-Gonzá-
lez, V. Calderón, Durability behavior of steelmaking slag masonry
mortars, Materials & Design, 97 (2016), 307–315, doi:10.1016/
j.matdes.2016.02.080
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758 757
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
diagram "22"). A combination of frost and chloride ions
caused a decrease in the strength of 27.5 % (MBS-90Cl)
and 29.6 % (MFA-90Cl). These values can be assessed
as adequate for the given type of material (i.e., the
material with a porous aggregate) after a given number
of cycles. The thermal resistance (a relative decrease in
the compressive strength) of the materials in individual
environments after being exposed to 1000 °C was
approximately 64–69 % (72–78 %) for the reference
samples (the samples exposed to the aggressive envi-
ronment). It is interesting that slag (i.e., MBS) showed
better results than the other reference material; however,
after the exposure to the aggressive environment and
high temperature, fly ash was better (i.e., MFA). For
illustration, MBS-90/Cl (MFA-90/Cl) showed a decrease
in the compressive strength of 75.1 % (74.1 %; Figure
2). A slight increase in the thermal resistance developed
over time in spite of the higher number of cycles in the
aggressive environment. The only exception was
MFA-90/Cl. An evaluation of the relative change in the
thermal resistance caused by the aggressive environment
shows an irregular development of the curves (Figure 3).
This is particularly clear for temperatures of 600 °C and
above. It is evident that the exposure to an aggressive
environment and 1000 °C caused a decrease in the
observed parameter of MBS (MFA) of approximately
11 % (2.5–7.2 %).
Further, the bending tensile strength was evaluated.
The differences are less obvious compared to the
development of the compressive strength (Figures 1 and
4). The increase in the bending strength in time was also
less considerable than the increase in the compressive
strength. The bending strengths of the test specimens not
exposed to aggressive environments and a thermal load
were from 7.3 N mm–2 to 7.6 N mm–2 (6.9 N mm–2 to
7.2 N mm–2) for mix designs MBS (MFA). After the ex-
posure of the samples to the aggressive environment with
chloride ions, the decrease was 12.6 % for MBS-90 and
13.9 % for MFA-90.
The decrease in the tensile bending strength caused
by the exposure to aggressive environments is then much
lower than that of the compressive strength. The negative
impact on the thermal resistance was also smaller. The
decrease in the tensile bending strength was 65.8 %
(62.3 %) for MBS-90/R (MFA-90/R) after the exposure
to 1000 °C (Figure 5). These facts imply that the substi-
tution of the primary binder with fly ash had a more
positive effect from the viewpoint of a combined ex-
posure involving an aggressive environment and a ther-
mal load. On the other hand, comparing the changes in
the thermal resistance in time and after the exposure to
adverse environments shows (Figure 6) that blast-fur-
nace slag had better results (MBS). The reduction in the
thermal resistance after the exposure to 1000 °C was
5.3 % for MBS-90/Cl-R and 6.1 % for MFA/Cl-R. The
reduction in the thermal resistance of MFA after the
exposure to 600 °C is also interesting for both tested
ages. It can be assumed that used high-temperature fly
ash contains components that react with chloride
solutions, possibly only during the thermal load. It could
also be a synergy reaction of both mentioned influences.
The synergy effect of an aggressive environment and
extreme temperatures causes a considerable degradation
of the structures of the materials based on a silicate
matrix and a porous aggregate. The given case is con-
firmed with the results of physico-mechanical tests of
materials MBS and MFA. The CT examination proved
the above character of the crack changes after the expo-
sure to the aggressive environment (chlorides and frost)
and the subsequent extreme thermal load. On the refe-
rence materials, only a small amount of randomly
oriented cracks was identified (Figure 7). The cracks
were rather small, i.e., with a smaller width. These
cracks mostly did not reach the transition zone between
the matrix and dense aggregate. On the contrary, the
number of cracks of the materials exposed to the aggres-
sive environment was larger (Figures 8 to 10). An
increase in the width and number of the cracks is evident
as the number of aggressive cycles and age of the sam-
ples grow. However, the increase was not considerable.
The orientation of the cracks was easily observable; the
cracks pointed towards the center of a test specimen
(Figures 8 and 9). The width and number of the cracks
and more cracks in the transit zone between the dense
aggregate and the matrix were evident compared to the
reference materials.
The last representative sample was FMA-60/Cl (Fig-
ure 10). It is interesting that its cracks are more close to
the surface, with an orientation similar to that of MBS.
Chloride ions combined with frost and the subsequent
thermal shock damaged more considerably the surface
areas of MFA, whereas MBS was damaged more evenly
throughout the whole structure.
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
756 Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 12: Microstructure of MBS-90/Cl (90 days, Cl – exposure to
chloride solution and frost environment, 100 cycles), exposure to
1000 °C, 5000× magnification
The XRD analysis was used for the identification of
crystal phases and their changes. XRD did not identify
Friedel’s salt in any of the tested specimens. The crucial
factor of the formation of this chemical compound is the
temperature, which is, inter alia, mentioned by Dela-
grave et al.3, Izquierdo et al.4, Yuan5 and Tang6. It is then
evident that the damage of the structure caused by the
combination of frost and the solution with chloride ions
is purely mechanical. The salt in the solution – NaCl – is
transported into the porous structure of the composite by
equalizing the concentration gradient. Sudden changes of
the temperatures above and below zero then cause the
crystallization of the salt and the change of the phase of
water, which necessarily causes a decrease in the me-
chanical properties of the given material. This corres-
ponds with the findings determined with CT. As regards
the changes in the structures of the analyzed materials, a
decomposition of chemical compounds (in particular
mineralogical phases) typical for silicates and the used
aggregate, due to the increasing temperature, was ob-
served. The development shows peaks with a decreasing
intensity and a gradual ceasing of portlandite, calcite and
CSH phases. At temperatures over 600 °C, mew phases
were observed (like -C2S).
SEM together with an element analysis showed the
presence of Cl– ions in some areas. Microcracks and the
degree of decomposition of the hydration products of the
matrix or the changes in the structure of the used aggre-
gate were marked, as shown in Figures 11 and 12.
Almost no failures were observable in the transit zone
between the porous aggregate (agloporite) and the
matrix. As regards the age of the tested materials/number
of cycles, no considerable differences in the microstruc-
ture were observed. A 5000× magnification clearly
shows a degraded polymer-silicate matrix. SEM images
clearly show spherical particles of fly ash also in MBS.
This was caused by an imperfect manufacturing process
of agloporite, during which not all granules (for the
production of the aggregate) were perfectly compacted
and fired. However, as this fly ash had gone through the
thermic process twice, it is almost an inert filler from the
viewpoint of high temperature. Only few microscopic
cracks were observed in the matrix.
5 CONCLUSIONS
The influence of a combination of an aggressive
environment (chloride solution and freeze-thaw cycles)
and a thermal load (of up to 1000 °C) was tested and
analyzed in detail on newly designed materials as it is
not a thoroughly examined area. Fly-ash agloporite is a
promising aggregate and its influence on polymer-sili-
cate materials is still not quite clear. The advantage of
agloporite is its thermal resistance and very good
interaction with a polymer-silicate matrix. However, an
exposure to cyclic changing of temperatures above and
below zero, intensified by a salt solution, makes the use
of the porous aggregate difficult. Blast-furnace slag
(MBS) and fly ash (MFA) were used for a modification
of the matrix composition. Physico-chemical and micro-
structural analyses proved that the action of the
above-mentioned environment disturbs the structures of
MBS and MFA in a mechanical way, causing an expan-
sion due to the change of the phase of water or crystalli-
zation of salt rather than the formation of new chemical
compounds. The extreme temperature then only inten-
sifies these phenomena by degrading the structures, in
particular that of the matrix. It was proved that CT is
crucial for the explanation of the changes/failures of the
inner structure. This analytical method is quite suitable
for examining the three-dimensional structure of a given
material. It is particularly applicable for revealing va-
rious cracks, their orientation, localization or number.
The tested materials have a large potential for future
research. An examination of the behavior of the mate-
rials over a longer time (e.g., 180 d and 360 d) would
also be very interesting.
Acknowledgements
This paper has been worked out with the financial
support of The Czech Science Foundation (GA^R),
project GA15-07657S, "Study of kinetics of processes
occurring in the composite system at extreme tem-
peratures and exposed to an aggressive environment", as
well as under project No. LO1408 "AdMaS UP –
Advanced Materials, Structures and Technologies",
supported by the Ministry of Education, Youth and
Sports under the National Sustainability Programme I.
6 REFERENCES
1 J. G. Jang, H. J. Kim, H. K. Kim, H. K. Lee, Resistance of coal
bottom ash mortar against the coupled deterioration of carbonation
and chloride penetration, Materials & Design, 93 (2016), 160–167,
doi:10.1016/j.matdes.2015.12.074
2 M. Maes, N. De Belie, Resistance of concrete and mortar against
combined attack of chloride and sodium sulphate, Cement and Con-
crete Composites, 53 (2014), 59–72, doi:10.1016/j.cemconcomp.
2014.06.013
3 A. Delagrave, J. Marchand, J. P. Ollivier, S. Julien, K. Hazrati,
Chloride binding capacity of various hydrated cement paste systems,
Advanced Cement Based Materials, 6 (1997), 28–35, doi:
10.1016/S1065-7355(97)90003-1
4 D. Izquierdo, C. Alonso, C. Andrade, M. Castellote, Potentiostatic
determination of chloride threshold values for rebar depassivation –
experimental and statistical study, Electrochim. Acta, 49 (2004)
17–18, 2731–2739, doi:10.1016/j.electacta.2004.01.034
5 Q. Yuan, Fundamental studies on test methods for the transport of
chloride ions in cementitious materials, Doctoral Thesis, Ghent,
Ghent University, 2009
6 L. Tang, Chloride transport in concrete-measurement and prediction,
Doctoral Thesis, Göteborg, Chalmers University of Technology,
1996
7 F. P. Glasser, J. Marchand, E. Samson, Durability of concrete –
degradation phenomena involving detrimental chemical reactions,
Cem. Concr. Res, 38 (2007) 2, 226–246, doi:10.1016/j.cemconres.
2007.09.015
8 I. Santamaría-Vicario, A. Rodríguez, C. Junco, S. Gutiérrez-Gonzá-
lez, V. Calderón, Durability behavior of steelmaking slag masonry
mortars, Materials & Design, 97 (2016), 307–315, doi:10.1016/
j.matdes.2016.02.080
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758 757
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
9 S. Aydin, Development of a high-temperature-resistant mortar by
using slag and pumice, Fire Safety Journal, 43 (2008), 610–617,
doi:10.1016/j.firesaf.2008.02.001
10 Ý. Yüksel, R. Siddique, Ö. Özkan, Influence of high temperature on
the properties of concretes made with industrial by-products as fine
aggregate replacement, Construction and Building Materials, 25
(2011), 967–972, doi:10.1016/j.conbuildmat.2010.06.085
11 G. Yuan, Q. Li, The use of surface coating in enhancing the mecha-
nical properties and durability of concrete exposed to elevated
temperature, Construction and Building Materials, 95 (2015),
375–383, doi:10.1016/j.conbuildmat.2015.07.120
12 A. Nadeem, S. A. Memon, T. Yiu Lo, Mechanical performance,
durability, qualitative and quantitative analysis of microstructure of
fly ash and Metakaolin mortar at elevated temperatures, Construction
and Building Materials, 38 (2013), 338–347, doi:10.1016/
j.conbuildmat.2012.08.042
13 T. Harun, C. Ahmet, An experimental investigation of bond and
compressive strength of concrete with mineral admixtures at high
temperature, Arab. J. Sci. Eng., 33 (2008) 2B, 443–449
14 V. ^erný, [. Keprdová, Usability of fly ashes from Czech Republic
for sintered artificial aggregate, Advanced Materials Research,
(2014), 805–808
15 ^SN EN 1504-3, Products and systems for the protection and repair
of concrete structures – Definitions, requirements, quality control
and evaluation of conformity – Part 3: Structural and non-structural
repair, ^NI, 2006
16 ^SN EN 1363-1, Fire resistance tests – Part 1: General requirements,
^NI, 2013
17 ^SN 73 1326, including Z1, Resistance of cement concrete surface
to water and defrosting chemicals, ^NI, 2003
18 ^SN EN 12190, Products and systems for the protection and repair
of concrete structures – Test methods – Determination of com-
pressive strength of repair mortar, ^NI, 1999
19 ^SN EN 196-1, Methods of testing cement – Part 1: Determination
of strength, ^NI, 2005
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
758 Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
759–768
INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE
GRAIN-GROWTH BEHAVIOUR OF NEW Fe-Al BASED ALLOYS
WITH FINE Al2O3 PRECIPITATES
VPLIV TERMOMEHANSKE OBDELAVE FeAl ZLITIN S FINIMI
Al2O3 IZLO^KI NA RAST ZRN
Bohuslav Ma{ek1, Omid Khalaj1, Hana Jirková1, Jiøí Svoboda2,
Dagmar Bublíková1
1University of West Bohemia, Regional Technology Institute, Univerzitní 22, 306 14 Pilsen, Czech Republic
2Institute of Physics of Materials, Academy of Sciences of the Czech Republic, @i`kova 22, 616 62 Brno, Czech Republic
svobj@ipm.cz
Prejem rokopisa – received: 2016-07-20; sprejem za objavo – accepted for publication: 2017-04-20
doi:10.17222/mit.2016.232
To obtain the superior-high-temperature creep strength, a transformation of a fine-grained structure to large grains due to
abnormal grain growth or recrystallization is an important process in oxide-dispersion-strengthened (ODS) alloys. The
processing of steel is enabled with powder metallurgy, which utilizes powders consisting of a Fe-Al metal matrix with a large O
content, prepared with mechanical alloying, and their hot consolidation due to rolling. The thermomechanical characteristics of
new ODS alloys with a Fe-Al matrix are investigated in terms of the changes in the grain-size distribution. The recrystallization
and grain growth were quantified after heating up to 1200 °C, which is the typical consolidation temperature for standard
nanostructured ferritic steels. The results show that new ODS alloys are significantly affected by the thermomechanical
treatment leading to microstructural changes. Recrystallization is mostly affected by decreasing the deformation and increasing
the holding time, which leads to a growth of the grain size.
Keywords: grain growth, ODS alloys, steel, Fe-Al, Al2O3
Da bi pridobili najvi{jo temperaturo mo~i lezenja, je transformacija drobnozrnate strukture v ve~ja zrna ali celo v abnormalno
velika oziroma rekristalizacija pomemben proces pri zlitinah, oja~anih z oksidno disperzijo (angl. ODS). Obdelava jekla je
omogo~ena z metalurgijo v prahu, ki pomaga prahom, ki vsebujejo metalno matrico Fe-Al z veliko vsebnostjo kisika,
pripravljeno z mehanskim legiranjem in njihovo vro~o konsolidacijo pri valjanju. Termomehanske karakteristike novih
ODS-zlitin s Fe-Al matrico so bile preiskovane, ker je pri{lo do sprememb v velikosti porazdelitve zrn. Rekristalizacija in rast
zrn sta bili ovrednoteni po segretju do 1200 °C, ki je tipi~na temperatura konsolidacije za nanostrukturirana feritna jekla.
Rezultati ka`ejo, da na ODS-zlitine zelo vpliva termomehanska obdelava, ki pripelje do sprememb v mikrostrukturi. Na
rekristalizacijo najve~krat vpliva deformacija in pove~anje ~asa zadr`anosti, ki povzro~i rast velikost zrn.
Klju~ne besede: rast zrn, ODS-zlitine, jeklo, Fe-Al, Al2O3
1 INTRODUCTION
Historically, ODS alloys have been employed to
improve high-temperature mechanical properties. In
1910, the first utilization of an oxide dispersion was
reported by W. D. Coolidge. Using the classical powder
metallurgy, a tungsten-based alloy reinforced with tho-
rium oxides was developed to impede high-temperature
grain growth and therefore increase the life span of a
tungsten-filament lamp.1 After this first development,
several other applications with various metallic matrices
such as aluminium or nickel were implemented over the
decades. A notable improvement in the field was made
by J. S. Benjamin at the International Nickel Company
(INCO) laboratory: he proposed a new process based on
high-energy-milling powder metallurgy – later called
mechanical alloying.2 This process was introduced to
obtain a fine and homogeneous oxide dispersion within a
nickel matrix. Its aim was to produce high-temperature-
resistant materials for gas-turbine applications.3
Currently, mechanical alloying is still considered to
be the most effective process for obtaining fine and
homogeneously distributed particles. The volume frac-
tion of dispersed spherical oxides (usually Y2O3) is
typically below 1 % and the oxides typically have a
mean size of 5–30 nm. The mechanically alloyed powder
is then consolidated at high temperatures and pressures
to produce the bulk material in the form of a bar or tube
stock. Subsequently, different thermomechanical treat-
ments are applied to optimize its microstructure and
mechanical properties.
Oxides are much more resistant to coarsening in the
coarse-grained ferrite than the ’-precipitates in super-
alloys, and the limiting temperature for a long-term
operation is between 1000 °C and 1100 °C when the
mean size of oxides is 20–30 nm. This clearly indicates
that oxides are extremely stable and the microstructural
stability of ODS alloys is much higher than that of
nickel-based superalloys. However, a loss of mechanical
properties due to the coarsening of oxides was also
Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768 759
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:621.78:669.15 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)759(2017)
9 S. Aydin, Development of a high-temperature-resistant mortar by
using slag and pumice, Fire Safety Journal, 43 (2008), 610–617,
doi:10.1016/j.firesaf.2008.02.001
10 Ý. Yüksel, R. Siddique, Ö. Özkan, Influence of high temperature on
the properties of concretes made with industrial by-products as fine
aggregate replacement, Construction and Building Materials, 25
(2011), 967–972, doi:10.1016/j.conbuildmat.2010.06.085
11 G. Yuan, Q. Li, The use of surface coating in enhancing the mecha-
nical properties and durability of concrete exposed to elevated
temperature, Construction and Building Materials, 95 (2015),
375–383, doi:10.1016/j.conbuildmat.2015.07.120
12 A. Nadeem, S. A. Memon, T. Yiu Lo, Mechanical performance,
durability, qualitative and quantitative analysis of microstructure of
fly ash and Metakaolin mortar at elevated temperatures, Construction
and Building Materials, 38 (2013), 338–347, doi:10.1016/
j.conbuildmat.2012.08.042
13 T. Harun, C. Ahmet, An experimental investigation of bond and
compressive strength of concrete with mineral admixtures at high
temperature, Arab. J. Sci. Eng., 33 (2008) 2B, 443–449
14 V. ^erný, [. Keprdová, Usability of fly ashes from Czech Republic
for sintered artificial aggregate, Advanced Materials Research,
(2014), 805–808
15 ^SN EN 1504-3, Products and systems for the protection and repair
of concrete structures – Definitions, requirements, quality control
and evaluation of conformity – Part 3: Structural and non-structural
repair, ^NI, 2006
16 ^SN EN 1363-1, Fire resistance tests – Part 1: General requirements,
^NI, 2013
17 ^SN 73 1326, including Z1, Resistance of cement concrete surface
to water and defrosting chemicals, ^NI, 2003
18 ^SN EN 12190, Products and systems for the protection and repair
of concrete structures – Test methods – Determination of com-
pressive strength of repair mortar, ^NI, 1999
19 ^SN EN 196-1, Methods of testing cement – Part 1: Determination
of strength, ^NI, 2005
T. MELICHAR et al.: DURABILITY OF MATERIALS BASED ON A POLYMER-SILICATE MATRIX ...
758 Materiali in tehnologije / Materials and technology 51 (2017) 5, 751–758
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
759–768
INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE
GRAIN-GROWTH BEHAVIOUR OF NEW Fe-Al BASED ALLOYS
WITH FINE Al2O3 PRECIPITATES
VPLIV TERMOMEHANSKE OBDELAVE FeAl ZLITIN S FINIMI
Al2O3 IZLO^KI NA RAST ZRN
Bohuslav Ma{ek1, Omid Khalaj1, Hana Jirková1, Jiøí Svoboda2,
Dagmar Bublíková1
1University of West Bohemia, Regional Technology Institute, Univerzitní 22, 306 14 Pilsen, Czech Republic
2Institute of Physics of Materials, Academy of Sciences of the Czech Republic, @i`kova 22, 616 62 Brno, Czech Republic
svobj@ipm.cz
Prejem rokopisa – received: 2016-07-20; sprejem za objavo – accepted for publication: 2017-04-20
doi:10.17222/mit.2016.232
To obtain the superior-high-temperature creep strength, a transformation of a fine-grained structure to large grains due to
abnormal grain growth or recrystallization is an important process in oxide-dispersion-strengthened (ODS) alloys. The
processing of steel is enabled with powder metallurgy, which utilizes powders consisting of a Fe-Al metal matrix with a large O
content, prepared with mechanical alloying, and their hot consolidation due to rolling. The thermomechanical characteristics of
new ODS alloys with a Fe-Al matrix are investigated in terms of the changes in the grain-size distribution. The recrystallization
and grain growth were quantified after heating up to 1200 °C, which is the typical consolidation temperature for standard
nanostructured ferritic steels. The results show that new ODS alloys are significantly affected by the thermomechanical
treatment leading to microstructural changes. Recrystallization is mostly affected by decreasing the deformation and increasing
the holding time, which leads to a growth of the grain size.
Keywords: grain growth, ODS alloys, steel, Fe-Al, Al2O3
Da bi pridobili najvi{jo temperaturo mo~i lezenja, je transformacija drobnozrnate strukture v ve~ja zrna ali celo v abnormalno
velika oziroma rekristalizacija pomemben proces pri zlitinah, oja~anih z oksidno disperzijo (angl. ODS). Obdelava jekla je
omogo~ena z metalurgijo v prahu, ki pomaga prahom, ki vsebujejo metalno matrico Fe-Al z veliko vsebnostjo kisika,
pripravljeno z mehanskim legiranjem in njihovo vro~o konsolidacijo pri valjanju. Termomehanske karakteristike novih
ODS-zlitin s Fe-Al matrico so bile preiskovane, ker je pri{lo do sprememb v velikosti porazdelitve zrn. Rekristalizacija in rast
zrn sta bili ovrednoteni po segretju do 1200 °C, ki je tipi~na temperatura konsolidacije za nanostrukturirana feritna jekla.
Rezultati ka`ejo, da na ODS-zlitine zelo vpliva termomehanska obdelava, ki pripelje do sprememb v mikrostrukturi. Na
rekristalizacijo najve~krat vpliva deformacija in pove~anje ~asa zadr`anosti, ki povzro~i rast velikost zrn.
Klju~ne besede: rast zrn, ODS-zlitine, jeklo, Fe-Al, Al2O3
1 INTRODUCTION
Historically, ODS alloys have been employed to
improve high-temperature mechanical properties. In
1910, the first utilization of an oxide dispersion was
reported by W. D. Coolidge. Using the classical powder
metallurgy, a tungsten-based alloy reinforced with tho-
rium oxides was developed to impede high-temperature
grain growth and therefore increase the life span of a
tungsten-filament lamp.1 After this first development,
several other applications with various metallic matrices
such as aluminium or nickel were implemented over the
decades. A notable improvement in the field was made
by J. S. Benjamin at the International Nickel Company
(INCO) laboratory: he proposed a new process based on
high-energy-milling powder metallurgy – later called
mechanical alloying.2 This process was introduced to
obtain a fine and homogeneous oxide dispersion within a
nickel matrix. Its aim was to produce high-temperature-
resistant materials for gas-turbine applications.3
Currently, mechanical alloying is still considered to
be the most effective process for obtaining fine and
homogeneously distributed particles. The volume frac-
tion of dispersed spherical oxides (usually Y2O3) is
typically below 1 % and the oxides typically have a
mean size of 5–30 nm. The mechanically alloyed powder
is then consolidated at high temperatures and pressures
to produce the bulk material in the form of a bar or tube
stock. Subsequently, different thermomechanical treat-
ments are applied to optimize its microstructure and
mechanical properties.
Oxides are much more resistant to coarsening in the
coarse-grained ferrite than the ’-precipitates in super-
alloys, and the limiting temperature for a long-term
operation is between 1000 °C and 1100 °C when the
mean size of oxides is 20–30 nm. This clearly indicates
that oxides are extremely stable and the microstructural
stability of ODS alloys is much higher than that of
nickel-based superalloys. However, a loss of mechanical
properties due to the coarsening of oxides was also
Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768 759
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:621.78:669.15 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)759(2017)
observed in fine-grained ODS steels at temperatures of
about 800 °C for the systems with an extremely fine
oxide dispersion (about 5 nm).4 Thus, the size of oxides
as well as the grain size are also important for the
stability of the microstructure – coarser oxides are much
more stable than fine ones, which is in line with the
cubic law for coarsening kinetics.
In the consolidation step, the processing temperatures
are critical in order to retain the nanocrystalline structure
generated during the mechanical alloying and to impede
the particle coarsening and grain growth.5–9 The Ni- and
F-based ODS alloys rely on the formation of slowly
growing and strongly adherent chromium and aluminium
scales for their high temperature oxidation/corrosion
resistance. In the present study, the ODS alloys consist of
a ferritic Fe-Al matrix strengthened with about 4 %
volume fraction of Al2O3 particles.10–15 In order to obtain
a more detailed insight into these new groups of
materials, this paper will concentrate on both the general
microstructure and mechanical properties of ODS steels,
the phenomenon of recrystallization and the related
microstructural evolutions.
2 EXPERIMENTAL PART
2.1 Material preparation
The classical processing route used to produce Fe-Al
based alloys with fine Al2O3 precipitates is highly
dependent on powder metallurgy (Figure 1). The first
step of the elaboration is the mechanical alloying of
powders consisting of Fe-11 w/% Al matrix (90 w/% Fe
and 10 w/% Al) and 1 w/% of O2 in gas, which is
absorbed, in a low-energy ball mill developed by the
authors, enabling evacuation and filling with oxygen. It
has two steel containers (each holding 24 L) and each
container is filled with 80 steel balls with a diameter of
40 mm. The revolution speed varies between 20–75
min–1. However, the speed for our research was kept
constant (75 min–1) for all the material preparations.
This step allows the powder to be forced into a solid
solution. The milled powder is then deposited into a steel
container with a diameter of 70 mm, evacuated
(degassed) and sealed by welding. The steel container is
then heated up to a temperature of 750 °C and rolled in a
hot rolling mill to a thickness of 20 mm in the first
rolling step, then heated up to a temperature of 900 °C
and rolled to a thickness of 8 mm in the second step. A
6-mm-thick sheet of the ODS alloy covered on both
sides with a 1-mm-thick scale from the rolling container
was produced in this way. Afterwards, the specimens
were cut with a water jet (Figure 5).
Eight types of material were used in this research as
described in Table 1. They are all based on a Fe-10w/%
Al ferritic matrix with different particle sizes and 4 %
volume fraction of Al2O3.
Al2O3 powder was added to prepare the composite;
fine oxides in the ODS alloys were obtained with the
internal oxidation during mechanical alloying and
precipitated during hot consolidation. It should be noted
that in the case of Material 1, the Al2O3 particles were
added, but in all the other materials, oxides were pro-
duced due to internal oxidation as shown in Figure 1.
Different sizes of the oxides in ODS steels are due to
different heat treatments after hot rolling. SEM obser-
vations indicated several inhomogeneities due to the
material sticking to the walls of the milling container
during mechanical alloying.
Table 1: Material parameters
Material
No.
Material
type
Milling
time
(hours)
Ferritic
matrix
Vol.%
of
Al2O3
Typical
particle
size
(nm)
1 ODS Alloy 100 Fe10wt%Al 10 300
2 ODS Alloy 100 Fe10wt%Al 6 50-200
3 ODS Alloy 150 Fe10wt%Al 6 50-150
4 ODS Alloy 200 Fe10wt%Al 6 30-150
5 ODS Alloy 245 Fe10wt%Al 7 20-50
6* ODS Alloy 245 Fe10wt%Al 7 20-50
7* ODS Alloy 245 Fe10wt%Al 7 20-50
8* ODS Alloy 245 Fe10wt%Al 7 20-50
* Different rolling conditions
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
760 Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Material-preparation process These inhomogeneities can also influence the mecha-
nical properties of the material, but the mechanical
alloying process is steadily optimized with respect to the
homogeneity of the materials.
2.2 Specimen preparation
Two specimen types (Figure 2) were selected from
several examples, which exhibited the most homo-
geneous temperature fields for specimen type 1 and
simplicity of tests for specimen type 2. Specimen type 1
was chosen from six types of specimen shapes with
different active parts in the middle. All the samples were
carefully monitored using a thermal camera to see how
the temperature field is distributed within the active part
when heated up to 1200 °C. So, regarding the results,
specimen type 1 (5a) has the best homogeneity regarding
the temperature distribution. On the other hand, speci-
men type 2 is designed to have four different deforma-
tions (5, 8, 20 and 50) % at the same time. The shortest
part (9 mm) is designed to have no deformation during
the pressing and the rest have appropriate values of
deformation. This specimen is held in a hydraulic
forging press with two positioning side plates (10 mm
thick) (Figure 3) in order not to apply more deformation
than required. The prepared containers were annealed at
1000 °C over 16 h. After normal cooling at room tem-
perature, all the specimens were cut by a water-jet ma-
chine in the longitudinal direction (Figure 4) and then all
the specimens were removed from the steel containers.
After grinding, the thickness of the specimens was
approximately 6 mm.
2.3 Testing equipment
In order to speed up the tests, a hydraulic forging
press (Figure 5) was used to apply four different defor-
mations at the same time on the type-2 specimens. Also,
in order to investigate the thermomechanical treatment of
the specimens, a servo-hydraulic MTS thermomecha-
nical simulator (Figure 6) was used, which allows
various temperature-deformation paths to be run in order
to find the conditions leading to, e.g., the most effective
grain coarsening due to recrystallization. The thermo-
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768 761
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Positions of specimens on a rolled semi-product: a) type 1,
b) type 2
Figure 2: Specimen dimensions in mm: a) type 1, b) type 2
Figure 3: Positions of specimens between holding plates Figure 5: Hydraulic forging press
observed in fine-grained ODS steels at temperatures of
about 800 °C for the systems with an extremely fine
oxide dispersion (about 5 nm).4 Thus, the size of oxides
as well as the grain size are also important for the
stability of the microstructure – coarser oxides are much
more stable than fine ones, which is in line with the
cubic law for coarsening kinetics.
In the consolidation step, the processing temperatures
are critical in order to retain the nanocrystalline structure
generated during the mechanical alloying and to impede
the particle coarsening and grain growth.5–9 The Ni- and
F-based ODS alloys rely on the formation of slowly
growing and strongly adherent chromium and aluminium
scales for their high temperature oxidation/corrosion
resistance. In the present study, the ODS alloys consist of
a ferritic Fe-Al matrix strengthened with about 4 %
volume fraction of Al2O3 particles.10–15 In order to obtain
a more detailed insight into these new groups of
materials, this paper will concentrate on both the general
microstructure and mechanical properties of ODS steels,
the phenomenon of recrystallization and the related
microstructural evolutions.
2 EXPERIMENTAL PART
2.1 Material preparation
The classical processing route used to produce Fe-Al
based alloys with fine Al2O3 precipitates is highly
dependent on powder metallurgy (Figure 1). The first
step of the elaboration is the mechanical alloying of
powders consisting of Fe-11 w/% Al matrix (90 w/% Fe
and 10 w/% Al) and 1 w/% of O2 in gas, which is
absorbed, in a low-energy ball mill developed by the
authors, enabling evacuation and filling with oxygen. It
has two steel containers (each holding 24 L) and each
container is filled with 80 steel balls with a diameter of
40 mm. The revolution speed varies between 20–75
min–1. However, the speed for our research was kept
constant (75 min–1) for all the material preparations.
This step allows the powder to be forced into a solid
solution. The milled powder is then deposited into a steel
container with a diameter of 70 mm, evacuated
(degassed) and sealed by welding. The steel container is
then heated up to a temperature of 750 °C and rolled in a
hot rolling mill to a thickness of 20 mm in the first
rolling step, then heated up to a temperature of 900 °C
and rolled to a thickness of 8 mm in the second step. A
6-mm-thick sheet of the ODS alloy covered on both
sides with a 1-mm-thick scale from the rolling container
was produced in this way. Afterwards, the specimens
were cut with a water jet (Figure 5).
Eight types of material were used in this research as
described in Table 1. They are all based on a Fe-10w/%
Al ferritic matrix with different particle sizes and 4 %
volume fraction of Al2O3.
Al2O3 powder was added to prepare the composite;
fine oxides in the ODS alloys were obtained with the
internal oxidation during mechanical alloying and
precipitated during hot consolidation. It should be noted
that in the case of Material 1, the Al2O3 particles were
added, but in all the other materials, oxides were pro-
duced due to internal oxidation as shown in Figure 1.
Different sizes of the oxides in ODS steels are due to
different heat treatments after hot rolling. SEM obser-
vations indicated several inhomogeneities due to the
material sticking to the walls of the milling container
during mechanical alloying.
Table 1: Material parameters
Material
No.
Material
type
Milling
time
(hours)
Ferritic
matrix
Vol.%
of
Al2O3
Typical
particle
size
(nm)
1 ODS Alloy 100 Fe10wt%Al 10 300
2 ODS Alloy 100 Fe10wt%Al 6 50-200
3 ODS Alloy 150 Fe10wt%Al 6 50-150
4 ODS Alloy 200 Fe10wt%Al 6 30-150
5 ODS Alloy 245 Fe10wt%Al 7 20-50
6* ODS Alloy 245 Fe10wt%Al 7 20-50
7* ODS Alloy 245 Fe10wt%Al 7 20-50
8* ODS Alloy 245 Fe10wt%Al 7 20-50
* Different rolling conditions
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
760 Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Material-preparation process These inhomogeneities can also influence the mecha-
nical properties of the material, but the mechanical
alloying process is steadily optimized with respect to the
homogeneity of the materials.
2.2 Specimen preparation
Two specimen types (Figure 2) were selected from
several examples, which exhibited the most homo-
geneous temperature fields for specimen type 1 and
simplicity of tests for specimen type 2. Specimen type 1
was chosen from six types of specimen shapes with
different active parts in the middle. All the samples were
carefully monitored using a thermal camera to see how
the temperature field is distributed within the active part
when heated up to 1200 °C. So, regarding the results,
specimen type 1 (5a) has the best homogeneity regarding
the temperature distribution. On the other hand, speci-
men type 2 is designed to have four different deforma-
tions (5, 8, 20 and 50) % at the same time. The shortest
part (9 mm) is designed to have no deformation during
the pressing and the rest have appropriate values of
deformation. This specimen is held in a hydraulic
forging press with two positioning side plates (10 mm
thick) (Figure 3) in order not to apply more deformation
than required. The prepared containers were annealed at
1000 °C over 16 h. After normal cooling at room tem-
perature, all the specimens were cut by a water-jet ma-
chine in the longitudinal direction (Figure 4) and then all
the specimens were removed from the steel containers.
After grinding, the thickness of the specimens was
approximately 6 mm.
2.3 Testing equipment
In order to speed up the tests, a hydraulic forging
press (Figure 5) was used to apply four different defor-
mations at the same time on the type-2 specimens. Also,
in order to investigate the thermomechanical treatment of
the specimens, a servo-hydraulic MTS thermomecha-
nical simulator (Figure 6) was used, which allows
various temperature-deformation paths to be run in order
to find the conditions leading to, e.g., the most effective
grain coarsening due to recrystallization. The thermo-
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768 761
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Positions of specimens on a rolled semi-product: a) type 1,
b) type 2
Figure 2: Specimen dimensions in mm: a) type 1, b) type 2
Figure 3: Positions of specimens between holding plates Figure 5: Hydraulic forging press
mechanical simulator also allows a combination of
tensile and compressive deformation, thus accumulating
a high plastic deformation (and a high dislocation den-
sity) in a specimen.
2.4 Testing programme
The testing programme involved two different
groups. The tests are summarized in Table 2.
Tests in group A were carried out to investigate the
thermomechanical behaviour of different materials (1–8)
at different temperatures, looking at a single tensile and
compression loading with a constant strain rate of 1 s–1
(Figure 7). In order to give a clearer comparison of the
results, only the results at room temperature (RT), (800,
1000 and 1200) °C are presented.
Tests in group B were performed to investigate the
effects of the holding time at 1000 °C and the defor-
mation percentage on materials 7 and 8 with the type-2
specimen shape (Figure 2b). The specimen was de-
signed in order to apply four different deformations (5, 8,
20 and 50 %) at the same time (Figure 8), and sub-
sequently one specimen from each material was
quenched in water for immediate cooling and the rest
were kept in the furnace at three different holding times
(10, 20, 40) h, after which they were cooled slowly to
room temperature.
Treatment No. 1 was designed to apply three
different single deformations on each sample at different
temperatures. The specimen is first heated up to the
desired temperature in 5 min and is then held for another
five minutes. Immediately after these 10 min, the first
deformation (5 % compression) is applied to the speci-
men and then the temperature is held for another five
minutes. The second deformation (3 % tension) is
applied immediately after 5 min and again the tempe-
rature is held for five minutes and after that the last
deformation (50 % tension) is applied and the specimen
is left to cool down to room temperature.
Treatment No. 2 was designed to apply four different
deformations (5, 8, 20 and 50) % at the same time on a
specimen. First, the specimen is heated to the desired
temperature in 5 minutes and immediately after that, the
deformation is applied. The temperature is kept for
different holding times (0, 10, 20, 40) h and then the
specimen is left to cool down to room temperature.
3 RESULTS AND DISCUSSION
3.1 Test group A
Figure 9 shows the stress-strain curves for all the
materials at different temperatures, with 5 % com-
pression, corresponding to Treatment No. 1. Material 2
exhibits better strength at 30 °C (almost 500 MPa) and
800 °C (almost 300 MPa); at 1200 °C, material 1 still
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
762 Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768
Figure 7: Treatment No. 1
Figure 6: Treatment with a thermomechanical simulator
Table 2: Parameters of the testing programme
Test group MaterialNo.
Treatment
No.
Holding
time (h)
Maximum temperature
(°C)
Number of
tests Purpose of tests
A 1, 2, 3, 4, 5,6, 7, 8 1 –
1200, 1100, 1000, 900,
800, RT 48
Investigation of
thermomechanical behaviour
B 7,8 2
0, 10,
20, 40
1000 40 Investigation of grain-size growth
Figure 8: Treatment No. 2
shows better strength (almost 60 MPa), but not very
different from material 2 (almost 50 Mpa). It seems that
the difference in the compression strength between these
two materials is almost 37 % at RT; however, this
difference increased to 137 % at 800 °C. On the other
hand, all the materials have almost the same elastic
modulus and none of them fails at the 5 % compression.
However, material 4 shows a strange behaviour at
800 °C. The microstructure analysis shows that this
occurred because of an inhomogeneous formation of the
particles in the specimen.
The hot-working behaviour of alloys is generally
reflected by flow curves, which are a direct consequence
of microstructural changes: the nucleation and growth of
new grains, dynamic recrystallization (DRX), the gene-
ration of dislocations, work hardening (WH), the
rearrangement of dislocations and their dynamic
recovery (DRV). In deformed materials, DRX seems to
be the prominent softening mechanism at high tempe-
ratures. DRX occurs during the straining of metals at
high temperatures, characterized by a nucleation of low-
dislocation-density grains and their posterior growth,
producing a homogeneous grain structure when dynamic
equilibrium is reached. Material 4 showed a strange
curve shape at 800 °C. The test was repeated several
times and similar behaviour was observed each time. It
was concluded that this happens because of the
inhomogeneity of the microstructure of this material.
Figure 10 shows the stress-strain curves at different
temperatures corresponding to the 3 % tension of Treat-
ment No. 1. As can be seen in Figure 10, material 2
shows a higher strength at 30 °C (almost 650 MPa) and
800 °C (almost 270 MPa), but at 1200 °C, material 1
again shows a better strength (almost 65 MPa). This
value was 35 % lower when the temperature was in-
creased. In Figure 10, the strain at the maximum tension
is almost 2 % at RT (Figure 10a), then it increases to
almost 3 % at 800 °C (Figure 10b); however, it de-
creases to almost 1 % at 1200 °C (Figure 10c). All the
materials have almost the same elastic modulus and none
of them failed at the 3 % deformation. The yield stress
and the shape of the flow curves are sensitive to tempe-
rature. Comparing all these curves, it is found that
decreasing the deformation temperature increases the
yield-stress level. In other words, it prevents softening
due to dynamic recrystallization (DRX) and dynamic
recovery (DRV) and allows the deformed metals to
exhibit work hardening (WH). For every curve, after a
rapid increase in the stress to the peak value, the flow
stress decreases monotonically towards a steady-state
regime with a varying softening rate, which typically
indicates the onset of DRX, Figure 10c.
Figure 11 shows the stress-strain curves at different
temperatures for the 50 % tension of Treatment No. 1.
All eight materials failed at RT, but only four materials
failed below the 50 % tension at higher temperatures.
Material 2 failed at the 34 % strain, materials 4 and 5
failed at around 45 % strain and material 8 failed at
around 50 % at 800 °C. At 1200 °C, only material 1
failed at 41 % and material 2 failed at the 45 % strain.
From these curves, it can also be seen that the stress evo-
lution with strain exhibits three distinct stages.
Work hardening (WH) predominates in the first stage
and causes dislocations to polygonise into stable sub-
grains. The flow stress exhibits a rapid increase with the
increasing strain up to the critical value. Then DRX
occurs due to a large difference in the dislocation density
within subgrains or grains. When the critical driving
force of DRX is attained, new grains are nucleated along
the grain boundaries, deformation bands and disloca-
tions, resulting in the formation of equiaxed DRX grains.
In the second stage, the flow stress exhibits a smaller
and smaller increase until the peak value, or an inflection
of the work-hardening rate is reached. This shows that
thermal softening becomes more and more important due
to DRX and dynamic recovery (DRV) and it exceeds
WH.
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Figure 10: Stress-strain curves (3 % tension) for: a) RT, b) 800 °C, c)
1000 °C, d) 1200 °C
Figure 9: Stress-strain curves (5 % compression) for: a) RT, b)
800 °C, c) 1000 °C, d) 1200 °C
mechanical simulator also allows a combination of
tensile and compressive deformation, thus accumulating
a high plastic deformation (and a high dislocation den-
sity) in a specimen.
2.4 Testing programme
The testing programme involved two different
groups. The tests are summarized in Table 2.
Tests in group A were carried out to investigate the
thermomechanical behaviour of different materials (1–8)
at different temperatures, looking at a single tensile and
compression loading with a constant strain rate of 1 s–1
(Figure 7). In order to give a clearer comparison of the
results, only the results at room temperature (RT), (800,
1000 and 1200) °C are presented.
Tests in group B were performed to investigate the
effects of the holding time at 1000 °C and the defor-
mation percentage on materials 7 and 8 with the type-2
specimen shape (Figure 2b). The specimen was de-
signed in order to apply four different deformations (5, 8,
20 and 50 %) at the same time (Figure 8), and sub-
sequently one specimen from each material was
quenched in water for immediate cooling and the rest
were kept in the furnace at three different holding times
(10, 20, 40) h, after which they were cooled slowly to
room temperature.
Treatment No. 1 was designed to apply three
different single deformations on each sample at different
temperatures. The specimen is first heated up to the
desired temperature in 5 min and is then held for another
five minutes. Immediately after these 10 min, the first
deformation (5 % compression) is applied to the speci-
men and then the temperature is held for another five
minutes. The second deformation (3 % tension) is
applied immediately after 5 min and again the tempe-
rature is held for five minutes and after that the last
deformation (50 % tension) is applied and the specimen
is left to cool down to room temperature.
Treatment No. 2 was designed to apply four different
deformations (5, 8, 20 and 50) % at the same time on a
specimen. First, the specimen is heated to the desired
temperature in 5 minutes and immediately after that, the
deformation is applied. The temperature is kept for
different holding times (0, 10, 20, 40) h and then the
specimen is left to cool down to room temperature.
3 RESULTS AND DISCUSSION
3.1 Test group A
Figure 9 shows the stress-strain curves for all the
materials at different temperatures, with 5 % com-
pression, corresponding to Treatment No. 1. Material 2
exhibits better strength at 30 °C (almost 500 MPa) and
800 °C (almost 300 MPa); at 1200 °C, material 1 still
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762 Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768
Figure 7: Treatment No. 1
Figure 6: Treatment with a thermomechanical simulator
Table 2: Parameters of the testing programme
Test group MaterialNo.
Treatment
No.
Holding
time (h)
Maximum temperature
(°C)
Number of
tests Purpose of tests
A 1, 2, 3, 4, 5,6, 7, 8 1 –
1200, 1100, 1000, 900,
800, RT 48
Investigation of
thermomechanical behaviour
B 7,8 2
0, 10,
20, 40
1000 40 Investigation of grain-size growth
Figure 8: Treatment No. 2
shows better strength (almost 60 MPa), but not very
different from material 2 (almost 50 Mpa). It seems that
the difference in the compression strength between these
two materials is almost 37 % at RT; however, this
difference increased to 137 % at 800 °C. On the other
hand, all the materials have almost the same elastic
modulus and none of them fails at the 5 % compression.
However, material 4 shows a strange behaviour at
800 °C. The microstructure analysis shows that this
occurred because of an inhomogeneous formation of the
particles in the specimen.
The hot-working behaviour of alloys is generally
reflected by flow curves, which are a direct consequence
of microstructural changes: the nucleation and growth of
new grains, dynamic recrystallization (DRX), the gene-
ration of dislocations, work hardening (WH), the
rearrangement of dislocations and their dynamic
recovery (DRV). In deformed materials, DRX seems to
be the prominent softening mechanism at high tempe-
ratures. DRX occurs during the straining of metals at
high temperatures, characterized by a nucleation of low-
dislocation-density grains and their posterior growth,
producing a homogeneous grain structure when dynamic
equilibrium is reached. Material 4 showed a strange
curve shape at 800 °C. The test was repeated several
times and similar behaviour was observed each time. It
was concluded that this happens because of the
inhomogeneity of the microstructure of this material.
Figure 10 shows the stress-strain curves at different
temperatures corresponding to the 3 % tension of Treat-
ment No. 1. As can be seen in Figure 10, material 2
shows a higher strength at 30 °C (almost 650 MPa) and
800 °C (almost 270 MPa), but at 1200 °C, material 1
again shows a better strength (almost 65 MPa). This
value was 35 % lower when the temperature was in-
creased. In Figure 10, the strain at the maximum tension
is almost 2 % at RT (Figure 10a), then it increases to
almost 3 % at 800 °C (Figure 10b); however, it de-
creases to almost 1 % at 1200 °C (Figure 10c). All the
materials have almost the same elastic modulus and none
of them failed at the 3 % deformation. The yield stress
and the shape of the flow curves are sensitive to tempe-
rature. Comparing all these curves, it is found that
decreasing the deformation temperature increases the
yield-stress level. In other words, it prevents softening
due to dynamic recrystallization (DRX) and dynamic
recovery (DRV) and allows the deformed metals to
exhibit work hardening (WH). For every curve, after a
rapid increase in the stress to the peak value, the flow
stress decreases monotonically towards a steady-state
regime with a varying softening rate, which typically
indicates the onset of DRX, Figure 10c.
Figure 11 shows the stress-strain curves at different
temperatures for the 50 % tension of Treatment No. 1.
All eight materials failed at RT, but only four materials
failed below the 50 % tension at higher temperatures.
Material 2 failed at the 34 % strain, materials 4 and 5
failed at around 45 % strain and material 8 failed at
around 50 % at 800 °C. At 1200 °C, only material 1
failed at 41 % and material 2 failed at the 45 % strain.
From these curves, it can also be seen that the stress evo-
lution with strain exhibits three distinct stages.
Work hardening (WH) predominates in the first stage
and causes dislocations to polygonise into stable sub-
grains. The flow stress exhibits a rapid increase with the
increasing strain up to the critical value. Then DRX
occurs due to a large difference in the dislocation density
within subgrains or grains. When the critical driving
force of DRX is attained, new grains are nucleated along
the grain boundaries, deformation bands and disloca-
tions, resulting in the formation of equiaxed DRX grains.
In the second stage, the flow stress exhibits a smaller
and smaller increase until the peak value, or an inflection
of the work-hardening rate is reached. This shows that
thermal softening becomes more and more important due
to DRX and dynamic recovery (DRV) and it exceeds
WH.
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Figure 10: Stress-strain curves (3 % tension) for: a) RT, b) 800 °C, c)
1000 °C, d) 1200 °C
Figure 9: Stress-strain curves (5 % compression) for: a) RT, b)
800 °C, c) 1000 °C, d) 1200 °C
In the third stage, three types of curves can be re-
cognized:
• A gradual decrease to a steady state with DRX soft-
ening (Figure 11c).
• A continuous increase with significant work-harden-
ing (Figure 11a).
• A continuous decrease with significant DRX soften-
ing (Figure 11b).
3.2 Test group B
Figures 12 to 19 show the microstructures of
materials 7 and 8 without deformation and with the 50 %
deformation at different annealing times of (0, 10, 20 and
40) h at 1000 °C. It can be seen that by increasing the
annealing time, the number of low-angle grain boun-
daries (<15°) decreased. This indicates that recrystalli-
zation does not occur at the zones with shorter annealing
times (Figure 12a). A few banding zones with a high
density of low-angle grain boundaries are observed in
Figures 14a and 16a showing that recrystallization
occurs in most zones. Figure 18a shows that the banding
zones with a high density of low-angle grain boundaries
cannot be observed, indicating that almost completely
recrystallized ferrite grains can be obtained after
annealing at 1000 °C for 40 h.
Similar results were achieved for the rest of the
specimens; however, there is a major difference between
the microstructures. Material 7 is partially recrystallized,
but material 8 retained fine grains. Both materials were
prepared using the same procedure, but with different
rolling pressures during the hot-rolling process. During
the rolling process, a specimen passes through different
stages. In the first stage of rolling, the material is under a
considerable mechanical stress, resulting from an in-
ternally balanced elastic strain. This elastic strain is
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Figure 11: Stress-strain curves (50 % tension) for: a) RT, b) 800 °C,
c) 1000 °C, d) 1200 °C
Figure 12: Microstructure for material 7 without annealing a) without
deformation, b) with 50 % deformation
Figure 13: Microstructure of material 8 without annealing: a) without
deformation, b) with 50 % deformation
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Figure 16: Microstructure for material 7 with 20 h annealing: a) with-
out deformation, b) with 50 % deformation
Figure 14: Microstructure for material 7 after 10 h annealing: a) with-
out deformation, b) with 50 % deformation
Figure 17: Microstructure for material 8 h with 20 h annealing:
a) without deformation, b) with 50 % deformation
Figure 15: Microstructure for material 8 after 10 h annealing: a) with-
out deformation, b) with 50 % deformation
In the third stage, three types of curves can be re-
cognized:
• A gradual decrease to a steady state with DRX soft-
ening (Figure 11c).
• A continuous increase with significant work-harden-
ing (Figure 11a).
• A continuous decrease with significant DRX soften-
ing (Figure 11b).
3.2 Test group B
Figures 12 to 19 show the microstructures of
materials 7 and 8 without deformation and with the 50 %
deformation at different annealing times of (0, 10, 20 and
40) h at 1000 °C. It can be seen that by increasing the
annealing time, the number of low-angle grain boun-
daries (<15°) decreased. This indicates that recrystalli-
zation does not occur at the zones with shorter annealing
times (Figure 12a). A few banding zones with a high
density of low-angle grain boundaries are observed in
Figures 14a and 16a showing that recrystallization
occurs in most zones. Figure 18a shows that the banding
zones with a high density of low-angle grain boundaries
cannot be observed, indicating that almost completely
recrystallized ferrite grains can be obtained after
annealing at 1000 °C for 40 h.
Similar results were achieved for the rest of the
specimens; however, there is a major difference between
the microstructures. Material 7 is partially recrystallized,
but material 8 retained fine grains. Both materials were
prepared using the same procedure, but with different
rolling pressures during the hot-rolling process. During
the rolling process, a specimen passes through different
stages. In the first stage of rolling, the material is under a
considerable mechanical stress, resulting from an in-
ternally balanced elastic strain. This elastic strain is
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
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Figure 11: Stress-strain curves (50 % tension) for: a) RT, b) 800 °C,
c) 1000 °C, d) 1200 °C
Figure 12: Microstructure for material 7 without annealing a) without
deformation, b) with 50 % deformation
Figure 13: Microstructure of material 8 without annealing: a) without
deformation, b) with 50 % deformation
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Figure 16: Microstructure for material 7 with 20 h annealing: a) with-
out deformation, b) with 50 % deformation
Figure 14: Microstructure for material 7 after 10 h annealing: a) with-
out deformation, b) with 50 % deformation
Figure 17: Microstructure for material 8 h with 20 h annealing:
a) without deformation, b) with 50 % deformation
Figure 15: Microstructure for material 8 after 10 h annealing: a) with-
out deformation, b) with 50 % deformation
added to the jamming of the dislocation, which occurred
during cold forming.
In Figures 12, 14, 16 and 18, there is a visible alter-
nation in the distorted shape of the cold-work crystals at
the stress-relief stage (Figures 12a, 14a, 16a and 18a).
In this stage, new crystals begin to grow in the deformed
crystals. In the next stage (Figures 12b, 14b, 16b and
18b), the small crystals that formed in the previous stage
gradually grow into bigger crystals by absorbing each
other in the cannibal fashion, thus making the structure
relatively coarse grained.
Figures 12 to 19 show that with the increasing strain,
a cellblock structure gradually develops and the sizes of
the cellblocks and the cells decrease. In other words, in
material 7 (Figures 12, 14, 16 and 18), there is a great
transformation. It can be seen that with the increasing
strain, there are changes in the spacing of the dense
dislocation walls and micro-bands (DDW-MBs) and in
the cell size. It is seen that after a 50 % deformation, the
spacing of the DDW–MBs decreased to almost 50 μm,
close to the cell size. In addition, the rate of the decrease
in spacing with the increasing strains is much larger for
geometrically necessary boundaries (GNBs) than for
incidental dislocation boundaries (IDBs). However, in
material 8 (Figures 13, 15, 17 and 19), there was no
serious transformation and the structure still shows
ferrite and pearlite. The grain size is relatively similar for
both deformations. The grain boundaries can no longer
be seen clearly, as the ferrite precipitated with the pear-
lite, creating a grey shade instead of a clear black-and-
white contrast.
Figure 20 gives an overview of the grain sizes of
materials 7 and 8 with different deformations and
annealing times. It can be seen that as material 7 is
almost recrystallized, it has a bigger grain size than
material 8. From Figure 20 it is clear that by increasing
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Figure 20: Grain size for: a) material 7, b) material 8
Figure 19: Microstructure for material 8 h with 40 h annealing: a)
without deformation, b) with 50 % deformation
Figure 18: Microstructure for material 7 with 40 h annealing: a) with-
out deformation, b) with 50 % deformation
the annealing time, the grain size increases significantly
for material 7 (Figure 20a), however the rate of increase
in grain size is lower for material 8 (Figure 20b). The
grain size reached almost 120 μm after 40 hours
annealing without any deformation in material 7, while
the grains in material 8 reached almost 10 μm under the
same conditions. On the other hand, by increasing the
applied deformation, the grain size of both materials
decreased smoothly. Material 7 reached 50 μm by apply-
ing 50 % deformation without annealing while material 8
reached almost 3 μm under the same conditions. One can
thus conclude that the stability of grain microstructure
can be significantly influenced by the processing. The
analysis of the reasons for this will be the topic of future
work by the team.
4 CONCLUSIONS
If you mean that why we didn’t present the micro-
structe of material 1-6, because all the materials except 7
has fine microstructure similar to material 8, that is why
we decide to present only the microstructure results from
material 7 and 8 and compare them in this view point.
This paper outlines the influence of thermomecha-
nical treatment on the grain growth of new Fe-Al based
alloys with fine Al2O3 precipitates. Eight materials
differing in the amount and size of the oxides embedded
in the ferritic matrix were tested under different condi-
tions. The advantages of all the materials are their
low-cost, simplicity of preparation and significant me-
chanical properties together with micro structures, due to
the Fe-Al based ferritic matrix of the ODS alloy. The
results from material 1-6 described in the relevant
section (3.1. test group A) and the results from material
7-8 in case of microstructure described in section (3.2
test group B). As all the materials except 7 has fine
microstructure similar to material 8, only the micro-
structure results from material 7 and 8 compared in this
view point. It can be concluded that in general the oxide
dispersion significantly strengthens the material.
However, the typical form of the flow curve with DRX
softening, including a single peak followed by a steady
state flow as a plateau, is more recognizable at high
temperatures than at low temperatures. This is because at
high temperatures the DRX softening compensates for
the work hardening (WH), and both the peak stress and
the onset of steady state flow are therefore shifted to
lower strain levels. The characteristics of softening flow
behaviour coupled with DRX were investigated for eight
materials and can be summarized as follows:
Decreasing deformation temperature causes the flow
stress level to increase. In other words, it prevents the
occurrence of softening due to DRX and dynamic
recovery (DRV) and causes the deformed metals to
exhibit work hardening (WH).
For every curve, after a rapid increase in the stress to
a peak value, the flow stress decreases monotonically
towards a steady state regime (a steady state flow as a
plateau due to DRX softening is more recognizable at
higher temperatures). A varying softening rate typically
indicates the onset of DRX, and the stress evolution with
strain exhibits three distinct stages.
At higher temperatures, a higher DRX softening
compensates the WH, and both the peak stress and the
onset of steady state flow are therefore shifted to lower
strain levels.
The elastic part of the total strain amplitude is always
higher than the plastic part in all specimens tested, even
for the highest total strain amplitudes of 15 %. This is
further confirmation of the strong strengthening effect of
oxide particles.
Material 7 is more crystallised than material 8. Thus,
it has a larger grain size compared to the fine grains of
material 8.
Acknowledgements
This paper includes results from projects 14-24252S
Preparation and Optimization of Creep Resistant
Submicron-Structured Composite with Fe-Al Matrix and
Al2O3 Particles subsidised by the Czech Science Founda-
tion.
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B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768 767
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
added to the jamming of the dislocation, which occurred
during cold forming.
In Figures 12, 14, 16 and 18, there is a visible alter-
nation in the distorted shape of the cold-work crystals at
the stress-relief stage (Figures 12a, 14a, 16a and 18a).
In this stage, new crystals begin to grow in the deformed
crystals. In the next stage (Figures 12b, 14b, 16b and
18b), the small crystals that formed in the previous stage
gradually grow into bigger crystals by absorbing each
other in the cannibal fashion, thus making the structure
relatively coarse grained.
Figures 12 to 19 show that with the increasing strain,
a cellblock structure gradually develops and the sizes of
the cellblocks and the cells decrease. In other words, in
material 7 (Figures 12, 14, 16 and 18), there is a great
transformation. It can be seen that with the increasing
strain, there are changes in the spacing of the dense
dislocation walls and micro-bands (DDW-MBs) and in
the cell size. It is seen that after a 50 % deformation, the
spacing of the DDW–MBs decreased to almost 50 μm,
close to the cell size. In addition, the rate of the decrease
in spacing with the increasing strains is much larger for
geometrically necessary boundaries (GNBs) than for
incidental dislocation boundaries (IDBs). However, in
material 8 (Figures 13, 15, 17 and 19), there was no
serious transformation and the structure still shows
ferrite and pearlite. The grain size is relatively similar for
both deformations. The grain boundaries can no longer
be seen clearly, as the ferrite precipitated with the pear-
lite, creating a grey shade instead of a clear black-and-
white contrast.
Figure 20 gives an overview of the grain sizes of
materials 7 and 8 with different deformations and
annealing times. It can be seen that as material 7 is
almost recrystallized, it has a bigger grain size than
material 8. From Figure 20 it is clear that by increasing
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
766 Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 20: Grain size for: a) material 7, b) material 8
Figure 19: Microstructure for material 8 h with 40 h annealing: a)
without deformation, b) with 50 % deformation
Figure 18: Microstructure for material 7 with 40 h annealing: a) with-
out deformation, b) with 50 % deformation
the annealing time, the grain size increases significantly
for material 7 (Figure 20a), however the rate of increase
in grain size is lower for material 8 (Figure 20b). The
grain size reached almost 120 μm after 40 hours
annealing without any deformation in material 7, while
the grains in material 8 reached almost 10 μm under the
same conditions. On the other hand, by increasing the
applied deformation, the grain size of both materials
decreased smoothly. Material 7 reached 50 μm by apply-
ing 50 % deformation without annealing while material 8
reached almost 3 μm under the same conditions. One can
thus conclude that the stability of grain microstructure
can be significantly influenced by the processing. The
analysis of the reasons for this will be the topic of future
work by the team.
4 CONCLUSIONS
If you mean that why we didn’t present the micro-
structe of material 1-6, because all the materials except 7
has fine microstructure similar to material 8, that is why
we decide to present only the microstructure results from
material 7 and 8 and compare them in this view point.
This paper outlines the influence of thermomecha-
nical treatment on the grain growth of new Fe-Al based
alloys with fine Al2O3 precipitates. Eight materials
differing in the amount and size of the oxides embedded
in the ferritic matrix were tested under different condi-
tions. The advantages of all the materials are their
low-cost, simplicity of preparation and significant me-
chanical properties together with micro structures, due to
the Fe-Al based ferritic matrix of the ODS alloy. The
results from material 1-6 described in the relevant
section (3.1. test group A) and the results from material
7-8 in case of microstructure described in section (3.2
test group B). As all the materials except 7 has fine
microstructure similar to material 8, only the micro-
structure results from material 7 and 8 compared in this
view point. It can be concluded that in general the oxide
dispersion significantly strengthens the material.
However, the typical form of the flow curve with DRX
softening, including a single peak followed by a steady
state flow as a plateau, is more recognizable at high
temperatures than at low temperatures. This is because at
high temperatures the DRX softening compensates for
the work hardening (WH), and both the peak stress and
the onset of steady state flow are therefore shifted to
lower strain levels. The characteristics of softening flow
behaviour coupled with DRX were investigated for eight
materials and can be summarized as follows:
Decreasing deformation temperature causes the flow
stress level to increase. In other words, it prevents the
occurrence of softening due to DRX and dynamic
recovery (DRV) and causes the deformed metals to
exhibit work hardening (WH).
For every curve, after a rapid increase in the stress to
a peak value, the flow stress decreases monotonically
towards a steady state regime (a steady state flow as a
plateau due to DRX softening is more recognizable at
higher temperatures). A varying softening rate typically
indicates the onset of DRX, and the stress evolution with
strain exhibits three distinct stages.
At higher temperatures, a higher DRX softening
compensates the WH, and both the peak stress and the
onset of steady state flow are therefore shifted to lower
strain levels.
The elastic part of the total strain amplitude is always
higher than the plastic part in all specimens tested, even
for the highest total strain amplitudes of 15 %. This is
further confirmation of the strong strengthening effect of
oxide particles.
Material 7 is more crystallised than material 8. Thus,
it has a larger grain size compared to the fine grains of
material 8.
Acknowledgements
This paper includes results from projects 14-24252S
Preparation and Optimization of Creep Resistant
Submicron-Structured Composite with Fe-Al Matrix and
Al2O3 Particles subsidised by the Czech Science Founda-
tion.
5 REFERENCES
1 M. Mohan, R. Subramanian, Z. Alam, P. C. Angelo, Evaluation of
the Mechanical Properties OF A Hot Isostatically Pressed Yttria-
Dispersed Nickel-Based Superalloy, Mater. Tehnol., 48 (2014) 6,
899–904
2 W. Quadakkers, Oxidation of ODS alloys, Journal de Physique IV, 03
(1993) C9, 177–186, doi:10.1051/jp4:1993916
3 F. Pedraza, Low Energy-High Flux Nitridation of Metal Alloys:
Mechanisms, Microstructures and High Temperatures Oxidation
Behaviour, Mater. Tehnol., 42 (2008) 4, 157–169
4 O. Khalaj, B. Ma{ek, H. Jirkova, A. Ronesova, J. Svoboda, Investi-
gation on New Creep and Oxidation Resistant Materials, Mater.
Tehnol., 49 (2015) 4, 173–179, doi:10.17222/mit.2014.210
5 F. D. Fischer, J. Svoboda, P. Fratzl, A thermodynamic approach to
grain growth and coarsening, Journal of Philosophical Magazine, 83
(2003) 9, 1075–1093, doi:10.1080/0141861031000068966
6 M. J. Alinger, G. R. Odette, D. T. Hoelzer, On the role of alloy com-
position and processing parameters in nanocluster formation and
dispersion strengthening in nanostuctured ferritic alloys, Acta
Material, 57 (2009) 2, 392–406, doi:10.1016/j.actamat.2008.09.025
7 P. Unifantowicz, Z. Oksiuta, P. Olier, Y. de Carlan, N. Baluc,
Microstructure and mechanical properties of an ODS RAF steel
fabricated by hot extrusion or hot isostatic pressing, Fusion Engi-
neering and Design, 86 (2011), 2413–2416, doi:10.1016/j.fusengdes.
2011.01.022
8 M. A. Auger, V. de Castro, T. Leguey, A. Muñoz, R. Pareja, Micro-
structure and mechanical behavior of ODS and non-ODS Fe-14Cr
model alloys produced by spark plasma sintering, Journal of Nuclear
Materials, 436 (2013) 5, 68–75, doi:10.1016/j.jnucmat.2013.01.331
9 M. Kos, J. Fer~ec, M. Brun~ko, R. Rudolf, I. An`el, Pressing of
Partially Oxide-Dispersion-Strenghtened Copper using the ECAP
Process, Mater. Tehnol., 48 (2014) 3, 379–384, UDK
621.777.2:669.35’71
10 B. Ma{ek, O. Khalaj, Z. Nový, T. Kubina, H. Jirkova, J. Svoboda, C.
[tádler, Behaviour of New ODS Alloys under Single and Multiple
Deformation, Mater. Tehnol., 50 (2016) 6, 891–898, doi:10.17222/
mit.2015.156
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Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768 767
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
11 Marmy, P., Kruml, T., Low cycle fatigue of Eurofer 97, Journal of
Nuclear Materials, 377 (2008) 1, 52–58, doi:10.1016/j.jnucmat.
2008.02.054
12 M. Mi{ovi}, N. Tadi}, M. Ja}imovi}, M. Janji}, Deformations and
Velocities during the Cold Rolling of Aluminium Alloys, Mater.
Tehnol., 50 (2016) 1, 59-67, doi:10.17222/mit.2014.250
13 A. Grajcar, Microstructure Evolution of Advanced High-Strength
Trip-Aided Bainitic Steel, Mater. Tehnol., 49 (2015) 5, 715–720,
doi:10.17222/mit.2014.154
14 B. [u{tar{i~, I. Paulin, M. Godec, S. Glode`, M. [ori, J. Flasker, A.
Koro{ec, S. Kores, G. Abramovi~, DSC/TG of Al-based Alloyed
Powders for P/M Applications, Mater. Tehnol., 48 (2014) 4, 439–450
15 F. Tehovnik, J. Burja, B. Podgornik, M. Godec, F. Vode, Micro-
structural evolution of Inconel 625 during hot rolling, Mater. Tehnol.,
49 (2015) 5, 899–904, doi:10.17222/mit.2015.274
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
768 Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
K. MATUS et al.: ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH THE USE OF HIGH-RESOLUTION ...
769–773
ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH
THE USE OF HIGH-RESOLUTION ELECTRON MICROSCOPY
AND COMPUTER SIMULATION
RAZISKAVE OBORIN V ALUMINIJEVIH ZLITINAH Z
VISOKORESOLUCIJSKO ELEKTRONSKO MIKROSKOPIJO IN
RA^UNALNI[KO SIMULACIJO
Krzysztof Matus, Anna Tomiczek, Klaudiusz Go³ombek, Miros³awa Pawlyta
Silesian University of Technology, Institute of Engineering Materials and Biomaterials, 18A Konarskiego Str.,
44100 Gliwice, Poland
krzysztof.matus@polsl.pl
Prejem rokopisa – received: 2016-07-27; sprejem za objavo – accepted for publication: 2017-01-27
doi:10.17222/mit.2016.226
This paper presents the results of the tests using high-resolution transmission electron microscopy (HRTEM) with both
transmission and scanning modes, as well as energy-dispersive-spectroscopy (EDS) investigation of the AlSi9Cu alloy after
laser surface remelting. The possibility of using a computer simulation to identify precipitates in the analysed alloy was also
explored. The obtained results and computer simulations were compared. Moreover, this article presents the advantages of the
computer aid in solving the diffraction patterns and precipitates in supercell simulations.
Keywords: precipitation, crystallography, TEM, computer simulations
V prispevku so predstavljeni rezultati preiskave AlSi9Cu zlitine z laserskim pretaljevanjem povr{in z uporabo transmisijske
elektronske mikroskopije z visoko lo~ljivostjo (angl. HRTEM) na dva na~ina: s transmisijskim skeniranjem, kot tudi z
energijsko disperzijsko spektroskopijo (angl. EDS). Z uporabo ra~unalni{ke simulacije je bila preiskana mo`nost identifikacije
delcev v analizirani zlitini. Dobljeni rezultati in ra~unalni{ke simulacije so bili primerjani. Poleg tega ~lanek predstavlja
prednosti ra~unalni{ke simulacije pri {tudiji vzorcev difrakcije in oborin v posameznih stanjih celic.
Klju~ne besede: oborine, kristalografija, TEM, ra~unalni{ke simulacije
1 INTRODUCTION
Aluminium alloys are the most widely used alloys in
modern technology, mainly due to high specific strength
and low density. Other factors behind the widespread
distribution of aluminium alloys are their excellent
electrical conductivity and high corrosion resistance.1,2
The use of laser surface treatment to improve the utility
properties of aluminium alloys allows the remelting of
the surface layer of the material or its enrichment with
alloy elements. The remelting of the material surface and
rapid crystallisation allow an improvement of the mecha-
nical properties of alloys, mainly in the field of abrasion
resistance and tribological properties of the surface.
Through the generation of plenty of fine precipitates, the
strength is increased as well. The influence of alloying
element precipitates on aluminium alloys is a complex
issue and it is still one of the most promising topics in
materials science.3–7
Laser surface treatment ensures that the processed
material obtains new properties due to the rapid disper-
sion of the heat from the melted zone. This phenomenon
enables the crystallisation of very fine precipitates to
occur. Laser remelting can be used for small and large
elements. The most common alloyed layers have a thick-
ness from 0.3 3 mm to about 3 mm. A layer formed by
remelting usually has a high homogeneity as well as a
fine crystalline structure. The use of the laser technology
makes it possible to obtain surface layers with high con-
tents of alloying elements and a unique combination of
elements that conventional alloys rarely contain.8–11
Generally, the laser surface treatment is accom-
plished with two methods, which differ from one another
based on how the alloying addition is introduced to the
surface layer (which is schematically presented in
Figure 1). When the surface of a material is subjected to
a laser beam and then melted, the process is called
remelting (Figure 1a). When an alloying material is
applied to a surface and then melted with a laser beam
(most of the materials are added in the forms of tapes,
pastes or powders), the process is called alloying (Figure
1b).12,13
This article aims to identify the precipitates in an alu-
minium alloy after the laser treatment using transmission
electron microscopy and computer simulations.
2 EXPERIMENTAL PART
For the experimental procedure, cast AlSi9Cu alumi-
nium alloy was used. Its average chemical composition
is shown in Table 1. This alloy was subjected to
Materiali in tehnologije / Materials and technology 51 (2017) 5, 769–773 769
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.1:669.715:621.385.833:004.942 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)769(2017)
11 Marmy, P., Kruml, T., Low cycle fatigue of Eurofer 97, Journal of
Nuclear Materials, 377 (2008) 1, 52–58, doi:10.1016/j.jnucmat.
2008.02.054
12 M. Mi{ovi}, N. Tadi}, M. Ja}imovi}, M. Janji}, Deformations and
Velocities during the Cold Rolling of Aluminium Alloys, Mater.
Tehnol., 50 (2016) 1, 59-67, doi:10.17222/mit.2014.250
13 A. Grajcar, Microstructure Evolution of Advanced High-Strength
Trip-Aided Bainitic Steel, Mater. Tehnol., 49 (2015) 5, 715–720,
doi:10.17222/mit.2014.154
14 B. [u{tar{i~, I. Paulin, M. Godec, S. Glode`, M. [ori, J. Flasker, A.
Koro{ec, S. Kores, G. Abramovi~, DSC/TG of Al-based Alloyed
Powders for P/M Applications, Mater. Tehnol., 48 (2014) 4, 439–450
15 F. Tehovnik, J. Burja, B. Podgornik, M. Godec, F. Vode, Micro-
structural evolution of Inconel 625 during hot rolling, Mater. Tehnol.,
49 (2015) 5, 899–904, doi:10.17222/mit.2015.274
B. MA[EK et al.: INFLUENCE OF THERMOMECHANICAL TREATMENT ON THE GRAIN-GROWTH BEHAVIOUR ...
768 Materiali in tehnologije / Materials and technology 51 (2017) 5, 759–768
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
K. MATUS et al.: ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH THE USE OF HIGH-RESOLUTION ...
769–773
ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH
THE USE OF HIGH-RESOLUTION ELECTRON MICROSCOPY
AND COMPUTER SIMULATION
RAZISKAVE OBORIN V ALUMINIJEVIH ZLITINAH Z
VISOKORESOLUCIJSKO ELEKTRONSKO MIKROSKOPIJO IN
RA^UNALNI[KO SIMULACIJO
Krzysztof Matus, Anna Tomiczek, Klaudiusz Go³ombek, Miros³awa Pawlyta
Silesian University of Technology, Institute of Engineering Materials and Biomaterials, 18A Konarskiego Str.,
44100 Gliwice, Poland
krzysztof.matus@polsl.pl
Prejem rokopisa – received: 2016-07-27; sprejem za objavo – accepted for publication: 2017-01-27
doi:10.17222/mit.2016.226
This paper presents the results of the tests using high-resolution transmission electron microscopy (HRTEM) with both
transmission and scanning modes, as well as energy-dispersive-spectroscopy (EDS) investigation of the AlSi9Cu alloy after
laser surface remelting. The possibility of using a computer simulation to identify precipitates in the analysed alloy was also
explored. The obtained results and computer simulations were compared. Moreover, this article presents the advantages of the
computer aid in solving the diffraction patterns and precipitates in supercell simulations.
Keywords: precipitation, crystallography, TEM, computer simulations
V prispevku so predstavljeni rezultati preiskave AlSi9Cu zlitine z laserskim pretaljevanjem povr{in z uporabo transmisijske
elektronske mikroskopije z visoko lo~ljivostjo (angl. HRTEM) na dva na~ina: s transmisijskim skeniranjem, kot tudi z
energijsko disperzijsko spektroskopijo (angl. EDS). Z uporabo ra~unalni{ke simulacije je bila preiskana mo`nost identifikacije
delcev v analizirani zlitini. Dobljeni rezultati in ra~unalni{ke simulacije so bili primerjani. Poleg tega ~lanek predstavlja
prednosti ra~unalni{ke simulacije pri {tudiji vzorcev difrakcije in oborin v posameznih stanjih celic.
Klju~ne besede: oborine, kristalografija, TEM, ra~unalni{ke simulacije
1 INTRODUCTION
Aluminium alloys are the most widely used alloys in
modern technology, mainly due to high specific strength
and low density. Other factors behind the widespread
distribution of aluminium alloys are their excellent
electrical conductivity and high corrosion resistance.1,2
The use of laser surface treatment to improve the utility
properties of aluminium alloys allows the remelting of
the surface layer of the material or its enrichment with
alloy elements. The remelting of the material surface and
rapid crystallisation allow an improvement of the mecha-
nical properties of alloys, mainly in the field of abrasion
resistance and tribological properties of the surface.
Through the generation of plenty of fine precipitates, the
strength is increased as well. The influence of alloying
element precipitates on aluminium alloys is a complex
issue and it is still one of the most promising topics in
materials science.3–7
Laser surface treatment ensures that the processed
material obtains new properties due to the rapid disper-
sion of the heat from the melted zone. This phenomenon
enables the crystallisation of very fine precipitates to
occur. Laser remelting can be used for small and large
elements. The most common alloyed layers have a thick-
ness from 0.3 3 mm to about 3 mm. A layer formed by
remelting usually has a high homogeneity as well as a
fine crystalline structure. The use of the laser technology
makes it possible to obtain surface layers with high con-
tents of alloying elements and a unique combination of
elements that conventional alloys rarely contain.8–11
Generally, the laser surface treatment is accom-
plished with two methods, which differ from one another
based on how the alloying addition is introduced to the
surface layer (which is schematically presented in
Figure 1). When the surface of a material is subjected to
a laser beam and then melted, the process is called
remelting (Figure 1a). When an alloying material is
applied to a surface and then melted with a laser beam
(most of the materials are added in the forms of tapes,
pastes or powders), the process is called alloying (Figure
1b).12,13
This article aims to identify the precipitates in an alu-
minium alloy after the laser treatment using transmission
electron microscopy and computer simulations.
2 EXPERIMENTAL PART
For the experimental procedure, cast AlSi9Cu alumi-
nium alloy was used. Its average chemical composition
is shown in Table 1. This alloy was subjected to
Materiali in tehnologije / Materials and technology 51 (2017) 5, 769–773 769
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.1:669.715:621.385.833:004.942 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)769(2017)
remelting using high-power diode laser (HPDL) ROFIN
SINAR DL 020 with a radiation wavelength of 940±5
nm and a laser power of 2 kW with a spot size of
1.8 mm × 6.8 mm and argon as the shielding gas.
Table 1: Chemical composition of AlSi9Cu alloy, in mass fractions
(w/%)
Si Cu Fe Mn Mg Ti Al
9.0 1.0 0.7 0.4 0.3 0.3 balance
Samples for transmission-electron-microscopy
(TEM) observations were prepared as thin foils, where
the Gatan PIPS 691 precision ion polishing system was
used to prepare them. Studies were carried out using a
scanning transmission electron microscope S/TEM Titan
80-300 type (FEI, USA), operated at 300 kV and
equipped with a EDAX Li-drifted Si EDS detector
(Philips Electronic Instruments Corp. PV97-61850 ME).
Electron-microscopy observations were performed with
a probe Cs-corrected S/TEM Titan 80-300 FEI micro-
scope, equipped with an EDAX EDS detector (energy
resolution of 132 eV). HRTEM images (300 kV, Cs = 1.2
mm, 0.2 nm point resolution) were recorded using only a
condenser lens aperture, and thus without an objective
lens aperture. HAADF-STEM images were obtained
with a HAADF STEM detector (Model Fischione 3000).
A convergence angle of 0.974° was used with the
HAADF detector’s inner collection angle of 5.443°. The
probe size used was about 0.1 nm. The diffraction
patterns were collected by exploiting both the selected
area electron diffraction (SAED) and Fourier trans-
formations from HRTEM and HAADF images.
Chemical-composition examinations were carried out
using energy-dispersive spectroscopy (EDS) for the
identification of the precipitates which occurred in the
samples. To perform a simulation of unit cells, the
CrystalMaker software was used.14 The simulation of
supercells and the results of electron diffraction were
realised with the software "Química de Solidos y
Catalisis, Estructura y Química de Nanomateriales"
developed at the University of Cadiz.15,16
3 RESULTS AND DISCUSSION
Despite a low content of titanium in the analysed
material (0.3 %, Table 1), microscopic observations
revealed the presence and dense occurrence of titanium
nanoprecipitates in the shape of sticks, sized approxi-
mately 250 nm × 50 nm (Figures 2 and 3). The reason
for this is the fact that it is a high-melting-point element.
At the time of the laser alloying, the volume of the
material was treated with a laser beam and then rapidly
melted and cooled down. When lowering the
temperature, the precipitates of titanium were the first to
crystallize. After melting all the elements, including
titanium, there was an even distribution. The amount of
titanium is limited, so the size of the crystallites formed
did not exceed 250 nm. The boundary between the
matrix and the titanium precipitates is a favorable
nucleation area for the other precipitates to grow on
titanium rods. The EDS map presents the titanium
distribution and determines the localization of titanium
precipitates in the aluminium matrix (Figure 3).
Figure 4 shows an enlarged view of one of the pre-
cipitates containing Ti. The size of the precipitate is
K. MATUS et al.: ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH THE USE OF HIGH-RESOLUTION ...
770 Materiali in tehnologije / Materials and technology 51 (2017) 5, 769–773
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Diagram of laser treatment: a) remelting, b) alloying12,13
Figure 3: STEM-HAADF/EDS map which presents a magnification
of the area from Figure 2 and shows the distribution of titanium in the
aluminium matrix
Figure 2: STEM-HAADF image of titanium precipitates in an alumi-
nium-alloy matrix
approximately 150 nm × 30 nm. The contrast in the
STEM-HAADF image indicates a difference in the
chemical composition between the precipitate and the
matrix. The matrix area and three areas characterised by
differing contrasts (different chemical compositions) and
structures (different phases) are marked by A and B–D,
respectively. The results of the chemical analysis of the
selected areas are presented in Figure 5, in graphs A-D,
and Table 2.
Table 2: Chemical compositions of the areas marked with letters A–D
in Figure 5 (in amount fractions, a/%)
Al (K) Cu (K) Mg (K) Si (K) Ti (K)
point A 93.95 6.04 – – –
point B 18.92 15.04 39.83 26.19 –
point C 67.65 5.03 3.70 23.60 –
point D 75.64 6.35 – 5.75 12.24
In the centre of Figure 4, the precipitate marked with
symbol D, which contains mainly Ti, Al and Si, is
shown. The observed precipitate had an irregular, elon-
gated shape and a size of approximately 100 nm × 20 nm.
The grain boundary between the precipitate with
titanium and the matrix served as the nucleation point for
the other precipitates. Copper results were affected by a
holder and by the pole pieces, which were made of
copper. There were at least two other types of precipi-
tates, which were marked with symbols B and C. The
phase marked as B consisted of Al, Mg, Si and Cu
(Figure 5b). The intensities of the Mg, Si and Cu peaks
(K lines) were higher than that of the matrix. At the
phase boundary between the aluminium matrix and the
precipitate marked as B, two smaller precipitates,
marked as C, can be noticed. The chemical composition
of precipitate C was similar to that of precipitate B (Al,
Si, Mg and Cu), as shown in Figure 5c, but the relative
intensities of the peaks were different. Detected Mg
content in precipitate C was practically at the back-
ground level. A fast Fourier transform (FFT) of the
precipitate marked with the letter B is presented in
Figure 6a. Compared to the EDS spectrum, the
measurement of the spacings and angles in the FFT, with
its high-resolution image, and a computer simulation of
the diffraction pattern allow the identification of the
precipitate marked with letter B as the Q phase (with a
chemical composition close to Al3Cu2Mg9Si7). A
comparison of the simulated electron diffraction and the
experimental one allows determining the orientation of
the Q-phase precipitate as [001] (Figure 6b).
The unit cell of the Q phase was simulated along the
[001] direction (Figure 7). This data was necessary to
create a supercell in the Q phase (Figure 8a) with its
characteristic arrangement of atoms and marked regular
shape between the copper atoms. This was then adapted
to an enlarged HAADF image of the investigated preci-
pitate (Figure 8b). The HAADF detector enabled the
K. MATUS et al.: ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH THE USE OF HIGH-RESOLUTION ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 769–773 771
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Results of the chemical analysis of the areas marked with letters A–D in Figure 5
Figure 4: STEM-HAADF image of a complex precipitate in the
aluminium matrix
remelting using high-power diode laser (HPDL) ROFIN
SINAR DL 020 with a radiation wavelength of 940±5
nm and a laser power of 2 kW with a spot size of
1.8 mm × 6.8 mm and argon as the shielding gas.
Table 1: Chemical composition of AlSi9Cu alloy, in mass fractions
(w/%)
Si Cu Fe Mn Mg Ti Al
9.0 1.0 0.7 0.4 0.3 0.3 balance
Samples for transmission-electron-microscopy
(TEM) observations were prepared as thin foils, where
the Gatan PIPS 691 precision ion polishing system was
used to prepare them. Studies were carried out using a
scanning transmission electron microscope S/TEM Titan
80-300 type (FEI, USA), operated at 300 kV and
equipped with a EDAX Li-drifted Si EDS detector
(Philips Electronic Instruments Corp. PV97-61850 ME).
Electron-microscopy observations were performed with
a probe Cs-corrected S/TEM Titan 80-300 FEI micro-
scope, equipped with an EDAX EDS detector (energy
resolution of 132 eV). HRTEM images (300 kV, Cs = 1.2
mm, 0.2 nm point resolution) were recorded using only a
condenser lens aperture, and thus without an objective
lens aperture. HAADF-STEM images were obtained
with a HAADF STEM detector (Model Fischione 3000).
A convergence angle of 0.974° was used with the
HAADF detector’s inner collection angle of 5.443°. The
probe size used was about 0.1 nm. The diffraction
patterns were collected by exploiting both the selected
area electron diffraction (SAED) and Fourier trans-
formations from HRTEM and HAADF images.
Chemical-composition examinations were carried out
using energy-dispersive spectroscopy (EDS) for the
identification of the precipitates which occurred in the
samples. To perform a simulation of unit cells, the
CrystalMaker software was used.14 The simulation of
supercells and the results of electron diffraction were
realised with the software "Química de Solidos y
Catalisis, Estructura y Química de Nanomateriales"
developed at the University of Cadiz.15,16
3 RESULTS AND DISCUSSION
Despite a low content of titanium in the analysed
material (0.3 %, Table 1), microscopic observations
revealed the presence and dense occurrence of titanium
nanoprecipitates in the shape of sticks, sized approxi-
mately 250 nm × 50 nm (Figures 2 and 3). The reason
for this is the fact that it is a high-melting-point element.
At the time of the laser alloying, the volume of the
material was treated with a laser beam and then rapidly
melted and cooled down. When lowering the
temperature, the precipitates of titanium were the first to
crystallize. After melting all the elements, including
titanium, there was an even distribution. The amount of
titanium is limited, so the size of the crystallites formed
did not exceed 250 nm. The boundary between the
matrix and the titanium precipitates is a favorable
nucleation area for the other precipitates to grow on
titanium rods. The EDS map presents the titanium
distribution and determines the localization of titanium
precipitates in the aluminium matrix (Figure 3).
Figure 4 shows an enlarged view of one of the pre-
cipitates containing Ti. The size of the precipitate is
K. MATUS et al.: ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH THE USE OF HIGH-RESOLUTION ...
770 Materiali in tehnologije / Materials and technology 51 (2017) 5, 769–773
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Diagram of laser treatment: a) remelting, b) alloying12,13
Figure 3: STEM-HAADF/EDS map which presents a magnification
of the area from Figure 2 and shows the distribution of titanium in the
aluminium matrix
Figure 2: STEM-HAADF image of titanium precipitates in an alumi-
nium-alloy matrix
approximately 150 nm × 30 nm. The contrast in the
STEM-HAADF image indicates a difference in the
chemical composition between the precipitate and the
matrix. The matrix area and three areas characterised by
differing contrasts (different chemical compositions) and
structures (different phases) are marked by A and B–D,
respectively. The results of the chemical analysis of the
selected areas are presented in Figure 5, in graphs A-D,
and Table 2.
Table 2: Chemical compositions of the areas marked with letters A–D
in Figure 5 (in amount fractions, a/%)
Al (K) Cu (K) Mg (K) Si (K) Ti (K)
point A 93.95 6.04 – – –
point B 18.92 15.04 39.83 26.19 –
point C 67.65 5.03 3.70 23.60 –
point D 75.64 6.35 – 5.75 12.24
In the centre of Figure 4, the precipitate marked with
symbol D, which contains mainly Ti, Al and Si, is
shown. The observed precipitate had an irregular, elon-
gated shape and a size of approximately 100 nm × 20 nm.
The grain boundary between the precipitate with
titanium and the matrix served as the nucleation point for
the other precipitates. Copper results were affected by a
holder and by the pole pieces, which were made of
copper. There were at least two other types of precipi-
tates, which were marked with symbols B and C. The
phase marked as B consisted of Al, Mg, Si and Cu
(Figure 5b). The intensities of the Mg, Si and Cu peaks
(K lines) were higher than that of the matrix. At the
phase boundary between the aluminium matrix and the
precipitate marked as B, two smaller precipitates,
marked as C, can be noticed. The chemical composition
of precipitate C was similar to that of precipitate B (Al,
Si, Mg and Cu), as shown in Figure 5c, but the relative
intensities of the peaks were different. Detected Mg
content in precipitate C was practically at the back-
ground level. A fast Fourier transform (FFT) of the
precipitate marked with the letter B is presented in
Figure 6a. Compared to the EDS spectrum, the
measurement of the spacings and angles in the FFT, with
its high-resolution image, and a computer simulation of
the diffraction pattern allow the identification of the
precipitate marked with letter B as the Q phase (with a
chemical composition close to Al3Cu2Mg9Si7). A
comparison of the simulated electron diffraction and the
experimental one allows determining the orientation of
the Q-phase precipitate as [001] (Figure 6b).
The unit cell of the Q phase was simulated along the
[001] direction (Figure 7). This data was necessary to
create a supercell in the Q phase (Figure 8a) with its
characteristic arrangement of atoms and marked regular
shape between the copper atoms. This was then adapted
to an enlarged HAADF image of the investigated preci-
pitate (Figure 8b). The HAADF detector enabled the
K. MATUS et al.: ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH THE USE OF HIGH-RESOLUTION ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 769–773 771
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Results of the chemical analysis of the areas marked with letters A–D in Figure 5
Figure 4: STEM-HAADF image of a complex precipitate in the
aluminium matrix
distinction and contrast between the columns of atoms
containing elements of differentiated value Z. Images
were obtained by recording the intensity of the scattered
electrons as a function of the incident beam position on
the sample surface. Cu has the highest value of Z,
therefore the copper atoms were characterized as having
the best contrast, or the brightest points.
Using the same method, the area marked as C in
Figure 4 was identified as the U1 phase (MgAl2Si2), and
area D was determined as the TiAl3 phase.17,18
The precipitates created during the laser remelting of
aluminium alloys have very complex chemical compo-
sitions and crystallographic structures (Figures 5 and 7
and Table 2). The interface between the matrix and the
large amount of precipitation is a unique place to
increase further the creation of the precipitates, which
are successfully identified as the participations of the
complex U1 and Q phases. Particularly the precipitates
of the Q phase, which is a major equilibrium phase of the
Al-Mg-Si-Cu system, are noteworthy. The Q phase
belongs to the P-6 space group with lattice parameters:
a = b = 0.1039 nm, c = 0.409 nm, and  =  = 90°,  =
120°.17
4 CONCLUSIONS
Laser surface treatment, due to its requirement for a
large amount of energy, leads to a dissolution of the
surface layers of alloys due to a rapid dissipation of heat
into the interior of the material. This causes the remelted
material to solidify very quickly, resulting in the
formation of fine precipitates of high-melting elements,
which are unique places for the nucleation of other
precipitates. This type of treatment can increase the
hardness due to the refinement of the microstructure. The
completed research determined that the phase that forms
first during the cooling was TiAl3, on which other
precipitates with complex chemical compositions such as
the precipitates of the Q and U1 phases grew. With the
help of the computer simulations allowing the identifi-
cation of the phases of the process, the process was
shortened in comparison to the research using the data-
bases on a computer. Moreover, the use of simulation
software to obtain diffraction patterns simplified the
processes of identifying the phases and their orientations.
The simulation of unit cells and supercells allowed the
matching of the simulated phases with their high-reso-
lution images.
Acknowledgment
This publication was financed by the Ministry of
Science and Higher Education of Poland as the statutory
financial grant of the Faculty of Mechanical Engineer-
ing, SUT.
5 REFERENCES
1 C. Cayron, P. A. Buffat, Transmission electron microscopy study of
the ’ phase (Al–Mg–Si alloys) and QC phase (Al–Cu–Mg–Si
alloys): ordering mechanism and crystallographic structure, Acta
Mater., 48 (2000), 2639–2653, doi:10.1016/S1359-6454(00)00057-4
2 D. J. Chakrabarti, D. E. Laughlin, Phase relations and precipitation
in Al–Mg–Si alloys with Cu additions, Prog. Mater. Sci., 49 (2004),
389–410, doi:10.1016/S0079-6425(03)00031-8
3 A. L. Garcia-Garcia, I. Dominguez-Lopez, L. Lopez-Jimenez, J. D.
O. Barceinas-Sanchez, Comparative quantification and statistical
analysis of ’ and  precipitates in aluminum alloy AA7075-T651 by
TEM and AFM, Mater. Charact., 87 (2014), 116–124, doi:10.1016/
j.matchar.2013.11.007
K. MATUS et al.: ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH THE USE OF HIGH-RESOLUTION ...
772 Materiali in tehnologije / Materials and technology 51 (2017) 5, 769–773
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: a) Q-phase supercell with the selected pattern of copper
atoms and b) this pattern registered on the HAADF image
Figure 7: Unit cell of precipitate Q in the [001] direction
Figure 6: a) Fast Fourier transform (FFT) of a high-resolution image
and b) simulation of the diffraction pattern of the Q-phase precipitate
in the [001] direction
4 N. K. Mukhopadhyay, H. J. Chang, J. Y. Lee, D. H. Kim, Electron
microscopy of an icosahedral phase in a rapidly solidified
Al18Mg3Mn2 complex metallic alloy, Scr. Mater., 59 (2008),
1119–1122, doi:10.1016/j.scriptamat.2008.07.024
5 A. Biswas, D. J. Siegel, D. N. Seidman, Compositional evolution of
Q-phase precipitates in an aluminum alloy, Acta Mater., 75 (2014),
322–336, doi:10.1016/j.actamat.2014.05.001
6 F. Delmas, M. J. Casanove, P. Lours, A. Couret, A. Coujou, Quan-
titative TEM study of the precipitation microstructure in aluminium
alloy Al(MgSiCu) 6056 T6, Mater. Sci. Eng. A, 373 (2004),
doi:10.1016/j.msea.2003.12.068
7 W. Yang, S. Ji, M. Wang, Z. Li, Precipitation behaviour of Al–Zn–
Mg–Cu alloy and diffraction analysis from ’ precipitates in four
variants, J. Alloys Compounds, 610 (2014), 623–629, doi:10.1016/
j.jallcom.2014.05.061
8 Z. Hu, L. Wan, S. Wu, H. Wu, X. Liu, Microstructure and mecha-
nical properties of high strength die-casting Al–Mg–Si–Mn alloy,
Mater. Des., 46 (2013), 451–456, doi:10.1016/j.matdes.2012.10.020
9 J. Aguilar, M. Fehlbier, A. Ludwig, A. Bührig-Polaczek, P. R. Sahm,
Non-equilibrium globular microstructure suitable for semisolid
casting of light metal alloys by rapid slug cooling technology
(RSCT), Mater. Sci. Eng. A, 375–377 (2004), 651–655, doi:10.1016/
j.msea.2003.10.091
10 W. M. Lee, M. A. Zikry, High strain-rate modeling of the interfacial
effects of dispersed particles in high strength aluminum alloys, Int. J.
Solids Struct., 49 (2012), 3291–3300, doi:10.1016/j.ijsolstr.2012.
07.003
11 J. F. Nie, B. C. Muddle, On the form of the age-hardening response
in high strength aluminium alloys, Mater. Sci. Eng. A, 319–321
(2001), 448–451, doi:10.1016/S0921-5093(01)01054-1
12 J. C. Betts, Laser surface modification of aluminium and magnesium
alloys, Surf. Eng. Light Alloys, Woodhead Publishing, 2010,
444–474, doi:10.1533/9781845699451.2.444
13 H. C. Man, C. T. Kwok, T. M. Yue, Cavitation erosion and corrosion
behaviour of laser surface alloyed MMC of SiC and Si3N4 on Al
alloy AA6061, Surf. Coat. Technol., 132 (2000), 11–20,
doi:10.1016/S0257-8972(00)00729-5
14 http://www.crystalmaker.com
15 C. López-Cartes, J. A. Pérez-Omil, J. M. Pintado, J. J. Calvino,
Z. C. Kang, L. Eyring, Rare-earth oxides with fluorite-related struc-
tures: Their systematic investigation using HREM images, image
simulations and electron diffraction pattern simulations, Ultramicro-
scopy 80 (1999), 19–39, doi:10.1016/S0304-3991(99)00067-4
16 S. Bernal, F. J. Botana, J. J. Calvino, C. Lopez-Cartes, J. A. Perez-
Omil, J. M. Rodriguez-Izquierdo, The interpretation of HREM
images of supported metal catalysts using image simulation: profile
view images, Ultramicroscopy, 72 (1998), 135–164, doi:10.1016/
S0304-3991(98)00009-6
17 R.-K. Pan, L. Ma, First-principles study on the elastic properties of
B’ and Q phase in Al–Mg–Si (–Cu) alloys, Phys. Scr., 87 (2013),
doi:10.1088/0031-8949/87/01/015601
18 S. J. Andersen, C. D. Marioara, R. Vissers, A. Frøseth, H. W. Zand-
bergen, The structural relation between precipitates in Al–Mg–Si
alloys, the Al-matrix and diamond silicon, with emphasis on the
trigonal phase U1-MgAl2Si2, Mat. Sci. Eng. A – Struct., 444,
(2007), 157–169, doi:10.1016/j.msea.2006.08.084
K. MATUS et al.: ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH THE USE OF HIGH-RESOLUTION ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 769–773 773
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
distinction and contrast between the columns of atoms
containing elements of differentiated value Z. Images
were obtained by recording the intensity of the scattered
electrons as a function of the incident beam position on
the sample surface. Cu has the highest value of Z,
therefore the copper atoms were characterized as having
the best contrast, or the brightest points.
Using the same method, the area marked as C in
Figure 4 was identified as the U1 phase (MgAl2Si2), and
area D was determined as the TiAl3 phase.17,18
The precipitates created during the laser remelting of
aluminium alloys have very complex chemical compo-
sitions and crystallographic structures (Figures 5 and 7
and Table 2). The interface between the matrix and the
large amount of precipitation is a unique place to
increase further the creation of the precipitates, which
are successfully identified as the participations of the
complex U1 and Q phases. Particularly the precipitates
of the Q phase, which is a major equilibrium phase of the
Al-Mg-Si-Cu system, are noteworthy. The Q phase
belongs to the P-6 space group with lattice parameters:
a = b = 0.1039 nm, c = 0.409 nm, and  =  = 90°,  =
120°.17
4 CONCLUSIONS
Laser surface treatment, due to its requirement for a
large amount of energy, leads to a dissolution of the
surface layers of alloys due to a rapid dissipation of heat
into the interior of the material. This causes the remelted
material to solidify very quickly, resulting in the
formation of fine precipitates of high-melting elements,
which are unique places for the nucleation of other
precipitates. This type of treatment can increase the
hardness due to the refinement of the microstructure. The
completed research determined that the phase that forms
first during the cooling was TiAl3, on which other
precipitates with complex chemical compositions such as
the precipitates of the Q and U1 phases grew. With the
help of the computer simulations allowing the identifi-
cation of the phases of the process, the process was
shortened in comparison to the research using the data-
bases on a computer. Moreover, the use of simulation
software to obtain diffraction patterns simplified the
processes of identifying the phases and their orientations.
The simulation of unit cells and supercells allowed the
matching of the simulated phases with their high-reso-
lution images.
Acknowledgment
This publication was financed by the Ministry of
Science and Higher Education of Poland as the statutory
financial grant of the Faculty of Mechanical Engineer-
ing, SUT.
5 REFERENCES
1 C. Cayron, P. A. Buffat, Transmission electron microscopy study of
the ’ phase (Al–Mg–Si alloys) and QC phase (Al–Cu–Mg–Si
alloys): ordering mechanism and crystallographic structure, Acta
Mater., 48 (2000), 2639–2653, doi:10.1016/S1359-6454(00)00057-4
2 D. J. Chakrabarti, D. E. Laughlin, Phase relations and precipitation
in Al–Mg–Si alloys with Cu additions, Prog. Mater. Sci., 49 (2004),
389–410, doi:10.1016/S0079-6425(03)00031-8
3 A. L. Garcia-Garcia, I. Dominguez-Lopez, L. Lopez-Jimenez, J. D.
O. Barceinas-Sanchez, Comparative quantification and statistical
analysis of ’ and  precipitates in aluminum alloy AA7075-T651 by
TEM and AFM, Mater. Charact., 87 (2014), 116–124, doi:10.1016/
j.matchar.2013.11.007
K. MATUS et al.: ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH THE USE OF HIGH-RESOLUTION ...
772 Materiali in tehnologije / Materials and technology 51 (2017) 5, 769–773
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: a) Q-phase supercell with the selected pattern of copper
atoms and b) this pattern registered on the HAADF image
Figure 7: Unit cell of precipitate Q in the [001] direction
Figure 6: a) Fast Fourier transform (FFT) of a high-resolution image
and b) simulation of the diffraction pattern of the Q-phase precipitate
in the [001] direction
4 N. K. Mukhopadhyay, H. J. Chang, J. Y. Lee, D. H. Kim, Electron
microscopy of an icosahedral phase in a rapidly solidified
Al18Mg3Mn2 complex metallic alloy, Scr. Mater., 59 (2008),
1119–1122, doi:10.1016/j.scriptamat.2008.07.024
5 A. Biswas, D. J. Siegel, D. N. Seidman, Compositional evolution of
Q-phase precipitates in an aluminum alloy, Acta Mater., 75 (2014),
322–336, doi:10.1016/j.actamat.2014.05.001
6 F. Delmas, M. J. Casanove, P. Lours, A. Couret, A. Coujou, Quan-
titative TEM study of the precipitation microstructure in aluminium
alloy Al(MgSiCu) 6056 T6, Mater. Sci. Eng. A, 373 (2004),
doi:10.1016/j.msea.2003.12.068
7 W. Yang, S. Ji, M. Wang, Z. Li, Precipitation behaviour of Al–Zn–
Mg–Cu alloy and diffraction analysis from ’ precipitates in four
variants, J. Alloys Compounds, 610 (2014), 623–629, doi:10.1016/
j.jallcom.2014.05.061
8 Z. Hu, L. Wan, S. Wu, H. Wu, X. Liu, Microstructure and mecha-
nical properties of high strength die-casting Al–Mg–Si–Mn alloy,
Mater. Des., 46 (2013), 451–456, doi:10.1016/j.matdes.2012.10.020
9 J. Aguilar, M. Fehlbier, A. Ludwig, A. Bührig-Polaczek, P. R. Sahm,
Non-equilibrium globular microstructure suitable for semisolid
casting of light metal alloys by rapid slug cooling technology
(RSCT), Mater. Sci. Eng. A, 375–377 (2004), 651–655, doi:10.1016/
j.msea.2003.10.091
10 W. M. Lee, M. A. Zikry, High strain-rate modeling of the interfacial
effects of dispersed particles in high strength aluminum alloys, Int. J.
Solids Struct., 49 (2012), 3291–3300, doi:10.1016/j.ijsolstr.2012.
07.003
11 J. F. Nie, B. C. Muddle, On the form of the age-hardening response
in high strength aluminium alloys, Mater. Sci. Eng. A, 319–321
(2001), 448–451, doi:10.1016/S0921-5093(01)01054-1
12 J. C. Betts, Laser surface modification of aluminium and magnesium
alloys, Surf. Eng. Light Alloys, Woodhead Publishing, 2010,
444–474, doi:10.1533/9781845699451.2.444
13 H. C. Man, C. T. Kwok, T. M. Yue, Cavitation erosion and corrosion
behaviour of laser surface alloyed MMC of SiC and Si3N4 on Al
alloy AA6061, Surf. Coat. Technol., 132 (2000), 11–20,
doi:10.1016/S0257-8972(00)00729-5
14 http://www.crystalmaker.com
15 C. López-Cartes, J. A. Pérez-Omil, J. M. Pintado, J. J. Calvino,
Z. C. Kang, L. Eyring, Rare-earth oxides with fluorite-related struc-
tures: Their systematic investigation using HREM images, image
simulations and electron diffraction pattern simulations, Ultramicro-
scopy 80 (1999), 19–39, doi:10.1016/S0304-3991(99)00067-4
16 S. Bernal, F. J. Botana, J. J. Calvino, C. Lopez-Cartes, J. A. Perez-
Omil, J. M. Rodriguez-Izquierdo, The interpretation of HREM
images of supported metal catalysts using image simulation: profile
view images, Ultramicroscopy, 72 (1998), 135–164, doi:10.1016/
S0304-3991(98)00009-6
17 R.-K. Pan, L. Ma, First-principles study on the elastic properties of
B’ and Q phase in Al–Mg–Si (–Cu) alloys, Phys. Scr., 87 (2013),
doi:10.1088/0031-8949/87/01/015601
18 S. J. Andersen, C. D. Marioara, R. Vissers, A. Frøseth, H. W. Zand-
bergen, The structural relation between precipitates in Al–Mg–Si
alloys, the Al-matrix and diamond silicon, with emphasis on the
trigonal phase U1-MgAl2Si2, Mat. Sci. Eng. A – Struct., 444,
(2007), 157–169, doi:10.1016/j.msea.2006.08.084
K. MATUS et al.: ANALYSIS OF PRECIPITATES IN ALUMINIUM ALLOYS WITH THE USE OF HIGH-RESOLUTION ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 769–773 773
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
775–782
MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM
A REPRESENTATIVE 2D IMAGE AND/OR A 3D
RECONSTRUCTION BASED ON A STACK OF IMAGES
VREDNOTENJE MIKROSTRUKTUR Ni-SDC KERMETA Z 2D
IN/ALI 3D METODO
Gregor Kapun1,3, Marjan Marin{ek2, Franci Merzel1, Sa{o [turm3,4,
Miran Gaber{~ek1, Tina Skalar2
1National Institute of Chemistry, Laboratory for Materials Chemistry, Hajdrihova 19, Ljubljana, Slovenia
2University of Ljubljana, Faculty for Chemistry and Chemical Technology, Ve~na pot 113, Ljubljana, Slovenia
3Jo`ef Stefan International Postgraduate School, Jamova cesta 39, 1000 Ljubljana, Slovenia
4Jo`ef Stefan Institute, Jamova cesta 39, 1000 Ljubljana, Slovenia
Tina.Skalar@fkkt.uni-lj.si
Prejem rokopisa – received: 2016-08-16; sprejem za objavo – accepted for publication: 2017-01-24
doi:10.17222/mit.2016.256
This work provides an in-depth discussion on the subject of deriving microstructural parameters from a realistic 2D image or a
3D volume reconstruction based on a stack of images. As a convenient model material, a nickel/samaria-doped-ceria (Ni-SDC)
cermet synthetized by the citrate-nitrate combustion reaction is tested. Field-emission scanning electron microscopy (FE-SEM)
for 2D microstructure evaluation or focused-ion-beam/scanning-electron-microscopy (FIB-SEM) for 3D imaging is used for the
cermet characterization. Microstructural parameters, such as the volume ratio of phases, grain morphology, contiguity of phases
and triple-point boundary density, are quantitatively evaluated. These microstructural parameters reveal that a 2D microstruc-
tural evaluation provides a relatively accurate and quick analytical approach. However, it is shown that a detailed microstruc-
tural quantification of the parameters that are closely related to transport phenomena and electrochemical reactions in a porous
cermet is only possible through the three-dimensional (3D) quantitative material characterization. Based on the microstructural
evaluation, optimization of the Ni-SDC material used is briefly discussed.
Keywords: solid-oxide fuel cell, Ni-SDC anode, microstructure, FIB tomography
Ni-SDC kermet (nikelj/cerijev oksid dopiran s samarijevim oksidom) je eden izmed najpogosteje uporabljenih anodnih
materialov v gorivnih celicah s trdnim elektrolitom (angl. SOFC). Med pripravo materiala sodi tudi kvantitativno vrednotenje
mikrostrukturnih parametrov kot so velikost in porazdelitev velikosti zrn, dele`i faz, kontinuitete, gostota trojnih to~k (angl.
TPB), gostota materiala, itd. V tem prispevku `elimo primerjati dva razli~na na~ina analiziranja teh parametrov. Prva t.i. 2D
metoda temelji na obdelavi slik, posnetih z vrsti~nim elektronskim mikroskopom s poljsko emisijo (SEM). Medtem, ko 3D
metoda upo{teva 3D rekonstrukcijo vzorca iz slik s pomo~jo sistema s fokusiranim ionskim sklopom (FIB-SEM). Omenjeni
mikrostrukturni parametri razkrivajo, da je 2D pristop razmeroma hiter in natan~en. Vendar se je izkazalo, da je za natan~nej{o
kvantitativno analizo mikrostrukture, ki je tesno povezana s prenosom snovi in elektrokemijskimi reakcijami v anodnem
materialu, potreben 3D pristop. Metodi sta bili ocenjeni glede na njuno te`avnost izvedbe in ~asovno zahtevnost.
Klju~ne besede: SOFC, Ni-SDC anoda, mikrostruktura, FIB-tomografija
1 INTRODUCTION
High energy-conversion efficiency, fuel flexibility
and environmental friendliness can be seen as great
advantages of solid-oxide fuel cells (SOFCs), which may
become one of the leading energy-conversion techno-
logies converting chemical energy into electricity using a
wide spectrum of fuels.1 An SOFC system consists of
two porous electrodes, which are separated by a solid
ceramic electrolyte. Most commonly, the anode material
is based on two-phase cermets. Due to reduced operating
temperatures (600–800 °C) in modern SOFC devices,
yttrium-stabilized zirconia (YSZ) based cermets, i.e.,
Ni-YSZ, have been typically replaced with samarium-
doped ceria (SDC) cermets, i.e., Ni-SDC 2; namely, at
comparable temperatures samarium-doped ceria exhibits
ionic conductivities approximately one order of magni-
tude greater than YSZ.3
Several synthesis approaches for cermet preparation
are described in the literature. Certain approaches such
as mechanical mixing4–6 of metal oxides are associated
with high costs due to fine particle milling and may lead
to a non-uniform particle size distribution or insufficient
continuity of ceramic, metal and pore phases in the final
cermet.7 Some of these issues may be avoided by using
alternative methods such as co-precipitation,4,8 sol-gel,9
Pechini synthesis,10–12 spray pyrolysis13 or citrate/nitrate
combustion synthesis.4 All of them rely on producing
small particles, which may eventually result in a homo-
geneous phase distribution within the cermet after
thermal treatment.
The most important requirements for efficient elec-
tro-oxidation at the anode are: i) high catalytic activity
for fuel electro-oxidation, ii) chemical-, morphological-
and dimensional-stability, iii) electronic- and ionic-con-
Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782 775
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:543.427.3:621.3.035 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)775(2017)
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
775–782
MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM
A REPRESENTATIVE 2D IMAGE AND/OR A 3D
RECONSTRUCTION BASED ON A STACK OF IMAGES
VREDNOTENJE MIKROSTRUKTUR Ni-SDC KERMETA Z 2D
IN/ALI 3D METODO
Gregor Kapun1,3, Marjan Marin{ek2, Franci Merzel1, Sa{o [turm3,4,
Miran Gaber{~ek1, Tina Skalar2
1National Institute of Chemistry, Laboratory for Materials Chemistry, Hajdrihova 19, Ljubljana, Slovenia
2University of Ljubljana, Faculty for Chemistry and Chemical Technology, Ve~na pot 113, Ljubljana, Slovenia
3Jo`ef Stefan International Postgraduate School, Jamova cesta 39, 1000 Ljubljana, Slovenia
4Jo`ef Stefan Institute, Jamova cesta 39, 1000 Ljubljana, Slovenia
Tina.Skalar@fkkt.uni-lj.si
Prejem rokopisa – received: 2016-08-16; sprejem za objavo – accepted for publication: 2017-01-24
doi:10.17222/mit.2016.256
This work provides an in-depth discussion on the subject of deriving microstructural parameters from a realistic 2D image or a
3D volume reconstruction based on a stack of images. As a convenient model material, a nickel/samaria-doped-ceria (Ni-SDC)
cermet synthetized by the citrate-nitrate combustion reaction is tested. Field-emission scanning electron microscopy (FE-SEM)
for 2D microstructure evaluation or focused-ion-beam/scanning-electron-microscopy (FIB-SEM) for 3D imaging is used for the
cermet characterization. Microstructural parameters, such as the volume ratio of phases, grain morphology, contiguity of phases
and triple-point boundary density, are quantitatively evaluated. These microstructural parameters reveal that a 2D microstruc-
tural evaluation provides a relatively accurate and quick analytical approach. However, it is shown that a detailed microstruc-
tural quantification of the parameters that are closely related to transport phenomena and electrochemical reactions in a porous
cermet is only possible through the three-dimensional (3D) quantitative material characterization. Based on the microstructural
evaluation, optimization of the Ni-SDC material used is briefly discussed.
Keywords: solid-oxide fuel cell, Ni-SDC anode, microstructure, FIB tomography
Ni-SDC kermet (nikelj/cerijev oksid dopiran s samarijevim oksidom) je eden izmed najpogosteje uporabljenih anodnih
materialov v gorivnih celicah s trdnim elektrolitom (angl. SOFC). Med pripravo materiala sodi tudi kvantitativno vrednotenje
mikrostrukturnih parametrov kot so velikost in porazdelitev velikosti zrn, dele`i faz, kontinuitete, gostota trojnih to~k (angl.
TPB), gostota materiala, itd. V tem prispevku `elimo primerjati dva razli~na na~ina analiziranja teh parametrov. Prva t.i. 2D
metoda temelji na obdelavi slik, posnetih z vrsti~nim elektronskim mikroskopom s poljsko emisijo (SEM). Medtem, ko 3D
metoda upo{teva 3D rekonstrukcijo vzorca iz slik s pomo~jo sistema s fokusiranim ionskim sklopom (FIB-SEM). Omenjeni
mikrostrukturni parametri razkrivajo, da je 2D pristop razmeroma hiter in natan~en. Vendar se je izkazalo, da je za natan~nej{o
kvantitativno analizo mikrostrukture, ki je tesno povezana s prenosom snovi in elektrokemijskimi reakcijami v anodnem
materialu, potreben 3D pristop. Metodi sta bili ocenjeni glede na njuno te`avnost izvedbe in ~asovno zahtevnost.
Klju~ne besede: SOFC, Ni-SDC anoda, mikrostruktura, FIB-tomografija
1 INTRODUCTION
High energy-conversion efficiency, fuel flexibility
and environmental friendliness can be seen as great
advantages of solid-oxide fuel cells (SOFCs), which may
become one of the leading energy-conversion techno-
logies converting chemical energy into electricity using a
wide spectrum of fuels.1 An SOFC system consists of
two porous electrodes, which are separated by a solid
ceramic electrolyte. Most commonly, the anode material
is based on two-phase cermets. Due to reduced operating
temperatures (600–800 °C) in modern SOFC devices,
yttrium-stabilized zirconia (YSZ) based cermets, i.e.,
Ni-YSZ, have been typically replaced with samarium-
doped ceria (SDC) cermets, i.e., Ni-SDC 2; namely, at
comparable temperatures samarium-doped ceria exhibits
ionic conductivities approximately one order of magni-
tude greater than YSZ.3
Several synthesis approaches for cermet preparation
are described in the literature. Certain approaches such
as mechanical mixing4–6 of metal oxides are associated
with high costs due to fine particle milling and may lead
to a non-uniform particle size distribution or insufficient
continuity of ceramic, metal and pore phases in the final
cermet.7 Some of these issues may be avoided by using
alternative methods such as co-precipitation,4,8 sol-gel,9
Pechini synthesis,10–12 spray pyrolysis13 or citrate/nitrate
combustion synthesis.4 All of them rely on producing
small particles, which may eventually result in a homo-
geneous phase distribution within the cermet after
thermal treatment.
The most important requirements for efficient elec-
tro-oxidation at the anode are: i) high catalytic activity
for fuel electro-oxidation, ii) chemical-, morphological-
and dimensional-stability, iii) electronic- and ionic-con-
Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782 775
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:543.427.3:621.3.035 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)775(2017)
ductivity, iv) chemical-, thermal- and mechanical compa-
tibility to other cell components and v) open porosity.14,15
Many of those requirements are determined by the anode
microstructure. Specifically, the performance of com-
posite cermet anodes critically depends on features like
particle size distribution, contiguity and triple phase
boundary length.4,16 In attempts to find correlations
between the anode microstructure and selected measured
properties, authors usually rely on 2D cross-sectional
images obtained from optical or electron microscopy.17,18
From these images, 3D microstructural parameters are
estimated by using porous models based on geometrical
theories, such as general effective media (GEM),19 the
contiguity concept,20 the random register network mo-
del,21 the random packing spheres model22 and the
stochastic reconstruction.23 However, these models are
all based on various assumptions, which can only
indirectly describe the real 3D microstructure. Recently,
3D reconstructions of SOFC electrodes have been per-
formed by X-ray computed tomography as well as by
focused-ion-beam-coupled scanning electron microscopy
(FIB-SEM).24–26 The recently highly improved spatial
resolution of the latter method is based on the serial
sectioning by FIB and SEM imaging of a sample
followed by reconstruction of the SEM micrographs into
a single data set, which gives 3D information of the
probed volume. Direct observation of a real electrode 3D
structure is a key towards the establishment of a reliable
relationship between a porous anode microstructure and
the cell’s power-generation performance.
The aim of this work is to compare a Ni-SDC micro-
structure obtained by the conventional 2D method using
SEM with the state-of-the-art FIB-SEM-based 3D
approach. The advantages and challenges of both me-
thods are discussed in significant detail, with a special
emphasis on the accuracy and applicability of the
obtained results as well as on the time required for pre-
paration and conducting the analysis. Both approaches
track well the characteristic microstructure parameters
that of a porous cermet structure, such as volume frac-
tion of phases, the average grain sizes and grain size
distribution in various directions, grain shape factor,
porosity (open and closed), contiguity of phases (C) and
triple-point boundary density (TPB).
2 EXPERIMENTAL PART
2.1 Material preparation
Anode materials were synthesized via the citrate-
nitrate combustion synthesis by mixing aqueous solu-
tions of metal nitrates and citric acid. Cation precursors –
hydrated metal nitrates Ce(NO3)3 · 6H2O (99 %, Sigma-
Aldrich), Sm(NO3)3 . 6H2O (99.9 %, Alfa Aesar) and
Ni(NO3)2 · 6H2O (99 %, Acros Organics) and citric acid
(100%, Carlo Erba) were separately dissolved in distilled
water. The mixture was prepared by taking care that the
molar ratio between cerium and samarium cations was
80:20, and considering that the Ni volume content in the
final composite was 50 %. In order to ensure subsequent
self-sustaining combustion the molar ratio between citric
acid and nitrates was set to 0.2. The prepared mixture
was then held over a water bath at 60 °C under vacuum
(8 mbar) for at least 6 h until it transformed into a light
green brittle gel. The dried gel was grinded and compres-
sed into pellets (f = 12 mm, h = 30 mm, p = 100 MPa),
which were then immediately ignited at the top with a
hot metal tip to start a self-sustaining combustion reac-
tion that spread as a reaction zone throughout the pellet.
The synthesized NiO-SDC powder was then crushed in
an agate mortar, wet-milled (isopropanol) in a planetary
mill (Fritsch pulverisette 7 premium line) for 15 min
with 10 mm grinding balls and 500 min–1 (1st milling)
and additional 15 min with 1 mm grinding balls and
500 min–1 (2nd milling). After drying the milled powder
was calcined at 800 °C for 1 h and subsequently
uni-axially pressed into pellets (f = 6 mm, h = 2.6 mm,
p = 100 MPa). The formed pellets were sintered at
1400 °C for 1 h. To determine the final composite micro-
structure, all the sintered tablets were polished, thermally
etched and reduced in an Ar-H2 (5 % of volume frac-
tions) atmosphere at 900 °C for 2 h.
2.2 Material characterization
The rapid increase in temperature during combustion
inside the reaction zone was measured as a temperature
profile using an optical pyrometer (Impac, IPE 140,
based on sample brightness). The measuring range of the
pyrometer is from 50 °C to 1200 °C, and it has a very
quick response time (1.5 ms). The accuracy of the op-
tically measured temperature was ±2.5 °C below 400 °C
and ±0.4 % of a measured value (in °C) above 400 °C.
The particle size observations of the as-synthesized
NiO-SDC dispersion were performed on a JEOL
JEM-2100 transmission electron microscope (TEM)
operated at 200 kV and equipped with EDX. The TEM
samples were prepared by dispersing the prepared
powder in ethanol using ultrasonication followed by
deposition of the obtained suspension on holey carbon-
coated copper grids. Particle size distributions of
unmilled and milled powders after the synthesis were
obtained using a Microtrac Bluewave particle sizer and
measured as aqueous suspensions. The shrinkage curve
of the NiO-SDC during sintering up to 1450 °C with a
constant heating rate of 10 K min–1 was obtained with a
heating microscope (Leitz Wetzlar). FE-SEM (Zeiss,
Ultra Plus) was used for a conventional 2D cross-section
imaging of sintered and reduced pellets. FIB-SEM
(Zeiss, Crossbeam 540) was used for the image data
stack acquisition of sintered anode material to obtain 3D
reconstruction of the probed volume.
2.3 2D microstructure analysis
The quantitative 2D microstructure analysis involved
the determination of the fraction of phases (jNi, jSDC),
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
776 Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
intercept grain lengths in x and y directions (Feret x and
Feret y), diameter of the area-analogue circle (d which
equals D-CIRCLE), sphericity (F-CIRCLE) and micro-
structural porosity () were performed on 4–5 digital
images using Axiovision 4.8 image-analysis software.
The definitions and equations of these parameters are
listed in Figure 1. Geometrical porosity (geometrical) was
calculated from the Archimedes densities of the pellets
using measured dimensions (diameter, height and mass)
relative to the theoretical densities. In order to determine
the contiguity values (C), which represent a frequency of
contacts between grains of one or various phases, in a
three-phase mixture the Gurland principle27,28 modified
by J. H. Lee et al.29 was used. SEM micrographs of the
given sample were lined (Figure 1) vertically and
horizontally. Along with these lines contacts between
different phases were counted. The final number of
contacts (N) was normalized to 1 mm length of sample.
Finally, contiguities were calculated using Equations (1),
(2) and (3):
C
N
N N NNi Ni
Ni Ni
Ni Ni Ni pore Ni SDC'
−
−
− − −
=
+ +
2
2
(1)
C
N
N N NSDC SDC
SDC SDC
SDC SDC SDC pore Ni SDC'
−
−
− − −
=
+ +
2
2
(2)
C
N
N N N NNi SDC
Ni SDC
Ni Ni Ni pore Ni SDC SDC SDC
−
−
− − − −
=
+ + +
2
2 2 + −N SDC pore
(3)
2.4 3D microstructure analysis
In addition to the conventional 2D microstructure
analysis, a FIB-SEM serial sectioning method was used
for a quantitative 3D examination of the anode micro-
structures. The method allows direct observation of
critical microstructure parameters like phase distribu-
tions, grain connectivity and triple phase boundary
density. The prepared Ni-SDC cermet was infiltrated
with epoxy resin (EpoFix, Struers) under vacuum in
order to easily distinguish the pores and avoid mistakes
due to the depth of field during SEM imaging. Sub-
sequently, a lamella with a thickness of 150 μm was cut
out, polished down to 50 μm and glued directly to an
aluminium stub to prevent sample drift during FIB ope-
ration. The configuration of the FIB-SEM setup is
presented in Figure 2. Initially, a layer of platinum was
"in situ" deposited on top of the region of interest to
protect the surface and prevent the curtaining effect of
the cross-section during milling. The volume of interest
was separated using an optimised U-pattern pre-milling
procedure to prevent material re-deposition and shadow-
ing of the signals used for imaging and microanalysis.
The sample was then serially sectioned using an auto-
mated slicing procedure with drift correction algorithms
to obtain a series of 2D images with narrow and repro-
ducible spacing between the individual image planes.30
Experimental milling and imaging parameters were
optimized in order to obtain a high-quality 3D recon-
struction (10 nm × 10 nm × 20 nm voxel resolution) with
phase-contrast information. Individual phases were
identified from EDXS elemental maps and further seg-
mented according to their grey level. Raw images were
pre-processed with "in-house" programing using ImageJ
followed by phase segmentation based on thresholding,
edge detection and a region growth algorithm.31,32 After
phase separation, the 3D microstructure with final di-
mensions of 10 μm × 10 μm × 10 μm was reconstructed
using the Amira 5.4.5. software package.33,34 The 3D
structure was additionally separated into particles based
upon a three-dimensional watershed algorithm applied to
a distance measurement.35,36 Volume fractions, feature
size and size distributions for each individual phase were
calculated directly from the 3D reconstruction. The
specific surface area was calculated inside the Amira
5.4.5. software using a triangular approximation based
on the marching cubes algorithm. A summary of the
results is listed in Table 1.
In the final step of the 3D quantification, the seg-
mented data cube with a 3D data matrix/grid of dimen-
sions Np × Np × Np, with Np = 400 at a resolution of 20
nm/pixel was exported for the TPB length calculation
using an "in house" programing algorithm based on the
centroid method.32,37,38 Each grid element (pixel) gets
associated with one of the three different phases: the
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782 777
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Schematic view of FIB-SEM settings and the measuring
procedure
Figure 1: In the left-hand part the definitions of FERETX, FERETY,
DCIRCLE and FCIRCLE are presented. Lined SEM micrograph in the
right-hand part (horizontal lines A, B, C, D, … and vertical lines a, b,
c, d, …) with marked Ni-SDC phase contact (mark) and listed num-
ber of Ni-SDC contacts for lines "I" (left) and "g" (right)
ductivity, iv) chemical-, thermal- and mechanical compa-
tibility to other cell components and v) open porosity.14,15
Many of those requirements are determined by the anode
microstructure. Specifically, the performance of com-
posite cermet anodes critically depends on features like
particle size distribution, contiguity and triple phase
boundary length.4,16 In attempts to find correlations
between the anode microstructure and selected measured
properties, authors usually rely on 2D cross-sectional
images obtained from optical or electron microscopy.17,18
From these images, 3D microstructural parameters are
estimated by using porous models based on geometrical
theories, such as general effective media (GEM),19 the
contiguity concept,20 the random register network mo-
del,21 the random packing spheres model22 and the
stochastic reconstruction.23 However, these models are
all based on various assumptions, which can only
indirectly describe the real 3D microstructure. Recently,
3D reconstructions of SOFC electrodes have been per-
formed by X-ray computed tomography as well as by
focused-ion-beam-coupled scanning electron microscopy
(FIB-SEM).24–26 The recently highly improved spatial
resolution of the latter method is based on the serial
sectioning by FIB and SEM imaging of a sample
followed by reconstruction of the SEM micrographs into
a single data set, which gives 3D information of the
probed volume. Direct observation of a real electrode 3D
structure is a key towards the establishment of a reliable
relationship between a porous anode microstructure and
the cell’s power-generation performance.
The aim of this work is to compare a Ni-SDC micro-
structure obtained by the conventional 2D method using
SEM with the state-of-the-art FIB-SEM-based 3D
approach. The advantages and challenges of both me-
thods are discussed in significant detail, with a special
emphasis on the accuracy and applicability of the
obtained results as well as on the time required for pre-
paration and conducting the analysis. Both approaches
track well the characteristic microstructure parameters
that of a porous cermet structure, such as volume frac-
tion of phases, the average grain sizes and grain size
distribution in various directions, grain shape factor,
porosity (open and closed), contiguity of phases (C) and
triple-point boundary density (TPB).
2 EXPERIMENTAL PART
2.1 Material preparation
Anode materials were synthesized via the citrate-
nitrate combustion synthesis by mixing aqueous solu-
tions of metal nitrates and citric acid. Cation precursors –
hydrated metal nitrates Ce(NO3)3 · 6H2O (99 %, Sigma-
Aldrich), Sm(NO3)3 . 6H2O (99.9 %, Alfa Aesar) and
Ni(NO3)2 · 6H2O (99 %, Acros Organics) and citric acid
(100%, Carlo Erba) were separately dissolved in distilled
water. The mixture was prepared by taking care that the
molar ratio between cerium and samarium cations was
80:20, and considering that the Ni volume content in the
final composite was 50 %. In order to ensure subsequent
self-sustaining combustion the molar ratio between citric
acid and nitrates was set to 0.2. The prepared mixture
was then held over a water bath at 60 °C under vacuum
(8 mbar) for at least 6 h until it transformed into a light
green brittle gel. The dried gel was grinded and compres-
sed into pellets (f = 12 mm, h = 30 mm, p = 100 MPa),
which were then immediately ignited at the top with a
hot metal tip to start a self-sustaining combustion reac-
tion that spread as a reaction zone throughout the pellet.
The synthesized NiO-SDC powder was then crushed in
an agate mortar, wet-milled (isopropanol) in a planetary
mill (Fritsch pulverisette 7 premium line) for 15 min
with 10 mm grinding balls and 500 min–1 (1st milling)
and additional 15 min with 1 mm grinding balls and
500 min–1 (2nd milling). After drying the milled powder
was calcined at 800 °C for 1 h and subsequently
uni-axially pressed into pellets (f = 6 mm, h = 2.6 mm,
p = 100 MPa). The formed pellets were sintered at
1400 °C for 1 h. To determine the final composite micro-
structure, all the sintered tablets were polished, thermally
etched and reduced in an Ar-H2 (5 % of volume frac-
tions) atmosphere at 900 °C for 2 h.
2.2 Material characterization
The rapid increase in temperature during combustion
inside the reaction zone was measured as a temperature
profile using an optical pyrometer (Impac, IPE 140,
based on sample brightness). The measuring range of the
pyrometer is from 50 °C to 1200 °C, and it has a very
quick response time (1.5 ms). The accuracy of the op-
tically measured temperature was ±2.5 °C below 400 °C
and ±0.4 % of a measured value (in °C) above 400 °C.
The particle size observations of the as-synthesized
NiO-SDC dispersion were performed on a JEOL
JEM-2100 transmission electron microscope (TEM)
operated at 200 kV and equipped with EDX. The TEM
samples were prepared by dispersing the prepared
powder in ethanol using ultrasonication followed by
deposition of the obtained suspension on holey carbon-
coated copper grids. Particle size distributions of
unmilled and milled powders after the synthesis were
obtained using a Microtrac Bluewave particle sizer and
measured as aqueous suspensions. The shrinkage curve
of the NiO-SDC during sintering up to 1450 °C with a
constant heating rate of 10 K min–1 was obtained with a
heating microscope (Leitz Wetzlar). FE-SEM (Zeiss,
Ultra Plus) was used for a conventional 2D cross-section
imaging of sintered and reduced pellets. FIB-SEM
(Zeiss, Crossbeam 540) was used for the image data
stack acquisition of sintered anode material to obtain 3D
reconstruction of the probed volume.
2.3 2D microstructure analysis
The quantitative 2D microstructure analysis involved
the determination of the fraction of phases (jNi, jSDC),
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
776 Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
intercept grain lengths in x and y directions (Feret x and
Feret y), diameter of the area-analogue circle (d which
equals D-CIRCLE), sphericity (F-CIRCLE) and micro-
structural porosity () were performed on 4–5 digital
images using Axiovision 4.8 image-analysis software.
The definitions and equations of these parameters are
listed in Figure 1. Geometrical porosity (geometrical) was
calculated from the Archimedes densities of the pellets
using measured dimensions (diameter, height and mass)
relative to the theoretical densities. In order to determine
the contiguity values (C), which represent a frequency of
contacts between grains of one or various phases, in a
three-phase mixture the Gurland principle27,28 modified
by J. H. Lee et al.29 was used. SEM micrographs of the
given sample were lined (Figure 1) vertically and
horizontally. Along with these lines contacts between
different phases were counted. The final number of
contacts (N) was normalized to 1 mm length of sample.
Finally, contiguities were calculated using Equations (1),
(2) and (3):
C
N
N N NNi Ni
Ni Ni
Ni Ni Ni pore Ni SDC'
−
−
− − −
=
+ +
2
2
(1)
C
N
N N NSDC SDC
SDC SDC
SDC SDC SDC pore Ni SDC'
−
−
− − −
=
+ +
2
2
(2)
C
N
N N N NNi SDC
Ni SDC
Ni Ni Ni pore Ni SDC SDC SDC
−
−
− − − −
=
+ + +
2
2 2 + −N SDC pore
(3)
2.4 3D microstructure analysis
In addition to the conventional 2D microstructure
analysis, a FIB-SEM serial sectioning method was used
for a quantitative 3D examination of the anode micro-
structures. The method allows direct observation of
critical microstructure parameters like phase distribu-
tions, grain connectivity and triple phase boundary
density. The prepared Ni-SDC cermet was infiltrated
with epoxy resin (EpoFix, Struers) under vacuum in
order to easily distinguish the pores and avoid mistakes
due to the depth of field during SEM imaging. Sub-
sequently, a lamella with a thickness of 150 μm was cut
out, polished down to 50 μm and glued directly to an
aluminium stub to prevent sample drift during FIB ope-
ration. The configuration of the FIB-SEM setup is
presented in Figure 2. Initially, a layer of platinum was
"in situ" deposited on top of the region of interest to
protect the surface and prevent the curtaining effect of
the cross-section during milling. The volume of interest
was separated using an optimised U-pattern pre-milling
procedure to prevent material re-deposition and shadow-
ing of the signals used for imaging and microanalysis.
The sample was then serially sectioned using an auto-
mated slicing procedure with drift correction algorithms
to obtain a series of 2D images with narrow and repro-
ducible spacing between the individual image planes.30
Experimental milling and imaging parameters were
optimized in order to obtain a high-quality 3D recon-
struction (10 nm × 10 nm × 20 nm voxel resolution) with
phase-contrast information. Individual phases were
identified from EDXS elemental maps and further seg-
mented according to their grey level. Raw images were
pre-processed with "in-house" programing using ImageJ
followed by phase segmentation based on thresholding,
edge detection and a region growth algorithm.31,32 After
phase separation, the 3D microstructure with final di-
mensions of 10 μm × 10 μm × 10 μm was reconstructed
using the Amira 5.4.5. software package.33,34 The 3D
structure was additionally separated into particles based
upon a three-dimensional watershed algorithm applied to
a distance measurement.35,36 Volume fractions, feature
size and size distributions for each individual phase were
calculated directly from the 3D reconstruction. The
specific surface area was calculated inside the Amira
5.4.5. software using a triangular approximation based
on the marching cubes algorithm. A summary of the
results is listed in Table 1.
In the final step of the 3D quantification, the seg-
mented data cube with a 3D data matrix/grid of dimen-
sions Np × Np × Np, with Np = 400 at a resolution of 20
nm/pixel was exported for the TPB length calculation
using an "in house" programing algorithm based on the
centroid method.32,37,38 Each grid element (pixel) gets
associated with one of the three different phases: the
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782 777
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Schematic view of FIB-SEM settings and the measuring
procedure
Figure 1: In the left-hand part the definitions of FERETX, FERETY,
DCIRCLE and FCIRCLE are presented. Lined SEM micrograph in the
right-hand part (horizontal lines A, B, C, D, … and vertical lines a, b,
c, d, …) with marked Ni-SDC phase contact (mark) and listed num-
ber of Ni-SDC contacts for lines "I" (left) and "g" (right)
SDC phase, the Ni phase and the pore/void "phase",
respectively. We identify the three-phase boundary
(TPB) elements/pixels satisfying the following condi-
tion: an element belonging to the pore is a TPB element
if it has pixels of both other two phases among its first
neighbours on the grid and at least one of its first
neighbours also belongs to the TPB. In order to avoid
multiplication of the TPB points it is important to
assume that the boundary element should refer to one
phase only (pore in our case). If a given TPB element is
not connected to any other TPB element, it does not form
a TPB line. The TPB length can be obtained, to a
reasonable approximation, as a product of the number of
TPB points and the grid spacing, NTPB × d0, that
corresponds to the edge summation. The corresponding
TPB density is then simply  = NTPBNP3d0–2. Here we
define one cubic layer of neighbouring elements around
the tagged pixel represented either by a single element
(Figure 3), leading to  = 7.23 μm μm–3.
Table 1: Microstructural and morphological parameters of Ni-SDC
determined by 2D or 3D approach
bulk 2D 3D
Ni / % 36.1 37.7 34.7
SDC / % 36.1 42.8 37.4
feret / μm
x
Ni 0.55 0.49
SDC 0.42 0.40
y Ni 0.44 0.44
SDC 0.38 0.34
z Ni / 0.51
SDC / 0.40
d / μm
Ni 0.52 0.48
SDC 0.38 0.37
F CIRCLE /
Ni 0.90 0.85*
SDC 0.78 0.71*
 / %
geometrical 27.8 / /
microstructural 21.5 27.9
open / 27.2
closed / 0.7
C /
Ni-Ni 0.05 /
Ni-SDC 0.18 /
SDC-SDC 0.47 /
TPB density / μm μm–3 / 3.5
SSA / μm μm–3
Ni / 15.4
SDC / 18.7
pore / 15.7
SIA / μm μm–3
Ni-pore / 6.2
Ni-SDC / 9.2
SDC-pore / 9.5
* In xy projection
3 RESULTS AND DISCUSSION
The combustion synthesis of NiO-SDC preparation is
chosen due to its major advantages over the classic cal-
cination method. Namely, the citrate-nitrate combustion
is a one-step, relatively simple method with rapid mate-
rial synthesis inside the combustion zone. Typically, in
the head of the combustion zone, changes from ambient
temperature to combustion temperature happen in a
fraction of a second, while after the combustion the
product ash is also quickly cooled. In the case of
NiO-SDC preparation, the measured peak combustion
temperature is 1169.2 °C, the propagation wave velocity
is 0.6 mm s–1 and the temperature gradient in the head of
the reaction zone is calculated to be 1053 °C s–1. Such
rapid temperature variations limit significantly the time
span during which diffusion of species in the system
plays an important role. This leads to a typical nano-
scaled mixture of NiO and SDC phases (Figure 4),
which serves as the starting material for the final
Ni-SDC cermet preparation.
The as-synthesized NiO-SDC product is partially
agglomerated. The size of the average agglomerate size
is about 1350 mm. The formed agglomerates of
NiO-SDC are weak and can be easily reduced in size
during milling (Figure 5). The first and second milling
cycles produce particles with average sizes of 5.23 mm
and 2.80 μm, respectively, which are still small agglo-
merates of nano-sized NiO and SDC one-phase regions.
The sintering behaviour of NiO-SDC during a dyna-
mic temperature rise is presented in Figure 6.
Densification of the tablet starts slightly above 800 °C,
proceeds over two separate steps, and is practically
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
778 Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Left: an example of the TPB element defined as one central
pixel surrounded by one layer of neighbouring pixels (pixels are
represented by balls, dark grey balls denote the SDC phase, black the
Ni phase and light grey the pores). Right: an image showing a distri-
bution of the TPB points in the sample ( = 7.23 μm μm–3)
Figure 4: High-resolution TEM image of as synthesized product
(some individual grains are encircled) and corresponding selected area
electron diffraction pattern for the image (insert)
completed at 1400 °C. The final shrinkage reached at
1450 °C is 24 %. Some features found in the sintering
curve may be clearly associated with the state of the
reacting matter during the synthesis. For example, the
relatively early start of the densification is a consequence
of the highly sinterable nano-sized mixture. However,
due to the presence of some agglomerates, the densifi-
cation first proceeds mainly as an inter-agglomerate
sintering with the maximal sintering rate at 950 °C.
Intra-agglomerate sintering starts at rather higher tempe-
ratures (above 1100 °C), reaching the maximum sinter-
ing rate at 1300 °C.
Among the crucial criteria for high-quality SOFC
anodes are the maximization of the TPB density, as well
as of the electronic and ionic conductivities. This means
that sintering must provide very good contacts between
the particles. Good particle-to-particle contact should
appear both inside one phase as well as between both
phases. Furthermore, the continuity of the Ni and SDC
phases throughout the cermet additionally contributes to
improved conductivities (ionic and electronic). At the
same time, the TPB density can be increased by prevent-
ing excessive grain coarsening during sintering. From
this point of view, the best compromise for the sintering
temperature for final cermet preparation was found to be
1400 °C. Any sintering temperature below 1400 °C may
result in the formation of continuous phases inside the
agglomerates only, whereas the intra-agglomerate con-
nection remains poor. Sintering temperatures above
1400 °C, however, result in exaggerated grain growth.
Microstructures of sintered NiO-SDC and sub-
sequently reduced Ni-SDC cermet are shown in Figure
7. Such a final Ni-SDC cermet was submitted to an
in-depth microstructure analysis. The dark-grey particles
seen in Figures 7a and 7b are NiO or Ni in the sintered
or reduced samples, respectively. The light grey is the
SDC phase, whereas the pores are black. It is evident
that NiO or Ni grains are homogeneously distributed
between the SDC phase. Furthermore, pores are also
uniformly distributed, especially in the reduced sample.
The parameters important for ann exact cermet 2D
analysis were extracted from the SEM micrographs by a
detailed quantitative microstructure analysis (Table 1).
Each parameter was determined on 5–10 different
regions on the surface of the analysed tablets. As a
whole, any quantitative calculation was based on several
hundreds of grains. After sintering at 1400 °C, the
remaining porosity in the NiO-SDC was around 97 %,
which means that the sample was densely sintered. After
the reduction, the microstructural porosity was increased
to 21.5 %. The value for the porosity, however, depends
to some extent on the determination method. The
microstructural porosity, obtained from conventional 2D
image analysis, tends to be lower than the geometrical
porosity. This may be explained by the fact that micro-
structural porosity is determined on a limited number of
images (which may be considered also as limited number
of slices throughout the sample), while the geometrical
porosity is a bulk property. To minimize this discrepancy
one has to increase the number of images for 2D micro-
structure analysis. Similarly, approximate numbers are
also determined for the phase fraction values (Ni and
SDC) and all the morphological values (Feret-values, –
F-CIRCLE and d). An interesting phenomenon may be
deduced from the contiguity values’ calculation. Namely,
the relatively low CNi-Ni value may indicate a poor contact
between the Ni grains. Such poor contact is highly
undesired, since it also results in a lower electrical
conductivity. However, the relatively high Ni content in
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782 779
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Particle size distribution of ungrounded powder (a), powder
after 1st milling cycle (b) and after 2nd milling cycle (c)
Figure 6: Sintering curve of NiO-SDC material calcined at 800 °C
(full line) and derivative sintering curve (dashed line)
SDC phase, the Ni phase and the pore/void "phase",
respectively. We identify the three-phase boundary
(TPB) elements/pixels satisfying the following condi-
tion: an element belonging to the pore is a TPB element
if it has pixels of both other two phases among its first
neighbours on the grid and at least one of its first
neighbours also belongs to the TPB. In order to avoid
multiplication of the TPB points it is important to
assume that the boundary element should refer to one
phase only (pore in our case). If a given TPB element is
not connected to any other TPB element, it does not form
a TPB line. The TPB length can be obtained, to a
reasonable approximation, as a product of the number of
TPB points and the grid spacing, NTPB × d0, that
corresponds to the edge summation. The corresponding
TPB density is then simply  = NTPBNP3d0–2. Here we
define one cubic layer of neighbouring elements around
the tagged pixel represented either by a single element
(Figure 3), leading to  = 7.23 μm μm–3.
Table 1: Microstructural and morphological parameters of Ni-SDC
determined by 2D or 3D approach
bulk 2D 3D
Ni / % 36.1 37.7 34.7
SDC / % 36.1 42.8 37.4
feret / μm
x
Ni 0.55 0.49
SDC 0.42 0.40
y Ni 0.44 0.44
SDC 0.38 0.34
z Ni / 0.51
SDC / 0.40
d / μm
Ni 0.52 0.48
SDC 0.38 0.37
F CIRCLE /
Ni 0.90 0.85*
SDC 0.78 0.71*
 / %
geometrical 27.8 / /
microstructural 21.5 27.9
open / 27.2
closed / 0.7
C /
Ni-Ni 0.05 /
Ni-SDC 0.18 /
SDC-SDC 0.47 /
TPB density / μm μm–3 / 3.5
SSA / μm μm–3
Ni / 15.4
SDC / 18.7
pore / 15.7
SIA / μm μm–3
Ni-pore / 6.2
Ni-SDC / 9.2
SDC-pore / 9.5
* In xy projection
3 RESULTS AND DISCUSSION
The combustion synthesis of NiO-SDC preparation is
chosen due to its major advantages over the classic cal-
cination method. Namely, the citrate-nitrate combustion
is a one-step, relatively simple method with rapid mate-
rial synthesis inside the combustion zone. Typically, in
the head of the combustion zone, changes from ambient
temperature to combustion temperature happen in a
fraction of a second, while after the combustion the
product ash is also quickly cooled. In the case of
NiO-SDC preparation, the measured peak combustion
temperature is 1169.2 °C, the propagation wave velocity
is 0.6 mm s–1 and the temperature gradient in the head of
the reaction zone is calculated to be 1053 °C s–1. Such
rapid temperature variations limit significantly the time
span during which diffusion of species in the system
plays an important role. This leads to a typical nano-
scaled mixture of NiO and SDC phases (Figure 4),
which serves as the starting material for the final
Ni-SDC cermet preparation.
The as-synthesized NiO-SDC product is partially
agglomerated. The size of the average agglomerate size
is about 1350 mm. The formed agglomerates of
NiO-SDC are weak and can be easily reduced in size
during milling (Figure 5). The first and second milling
cycles produce particles with average sizes of 5.23 mm
and 2.80 μm, respectively, which are still small agglo-
merates of nano-sized NiO and SDC one-phase regions.
The sintering behaviour of NiO-SDC during a dyna-
mic temperature rise is presented in Figure 6.
Densification of the tablet starts slightly above 800 °C,
proceeds over two separate steps, and is practically
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
778 Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Left: an example of the TPB element defined as one central
pixel surrounded by one layer of neighbouring pixels (pixels are
represented by balls, dark grey balls denote the SDC phase, black the
Ni phase and light grey the pores). Right: an image showing a distri-
bution of the TPB points in the sample ( = 7.23 μm μm–3)
Figure 4: High-resolution TEM image of as synthesized product
(some individual grains are encircled) and corresponding selected area
electron diffraction pattern for the image (insert)
completed at 1400 °C. The final shrinkage reached at
1450 °C is 24 %. Some features found in the sintering
curve may be clearly associated with the state of the
reacting matter during the synthesis. For example, the
relatively early start of the densification is a consequence
of the highly sinterable nano-sized mixture. However,
due to the presence of some agglomerates, the densifi-
cation first proceeds mainly as an inter-agglomerate
sintering with the maximal sintering rate at 950 °C.
Intra-agglomerate sintering starts at rather higher tempe-
ratures (above 1100 °C), reaching the maximum sinter-
ing rate at 1300 °C.
Among the crucial criteria for high-quality SOFC
anodes are the maximization of the TPB density, as well
as of the electronic and ionic conductivities. This means
that sintering must provide very good contacts between
the particles. Good particle-to-particle contact should
appear both inside one phase as well as between both
phases. Furthermore, the continuity of the Ni and SDC
phases throughout the cermet additionally contributes to
improved conductivities (ionic and electronic). At the
same time, the TPB density can be increased by prevent-
ing excessive grain coarsening during sintering. From
this point of view, the best compromise for the sintering
temperature for final cermet preparation was found to be
1400 °C. Any sintering temperature below 1400 °C may
result in the formation of continuous phases inside the
agglomerates only, whereas the intra-agglomerate con-
nection remains poor. Sintering temperatures above
1400 °C, however, result in exaggerated grain growth.
Microstructures of sintered NiO-SDC and sub-
sequently reduced Ni-SDC cermet are shown in Figure
7. Such a final Ni-SDC cermet was submitted to an
in-depth microstructure analysis. The dark-grey particles
seen in Figures 7a and 7b are NiO or Ni in the sintered
or reduced samples, respectively. The light grey is the
SDC phase, whereas the pores are black. It is evident
that NiO or Ni grains are homogeneously distributed
between the SDC phase. Furthermore, pores are also
uniformly distributed, especially in the reduced sample.
The parameters important for ann exact cermet 2D
analysis were extracted from the SEM micrographs by a
detailed quantitative microstructure analysis (Table 1).
Each parameter was determined on 5–10 different
regions on the surface of the analysed tablets. As a
whole, any quantitative calculation was based on several
hundreds of grains. After sintering at 1400 °C, the
remaining porosity in the NiO-SDC was around 97 %,
which means that the sample was densely sintered. After
the reduction, the microstructural porosity was increased
to 21.5 %. The value for the porosity, however, depends
to some extent on the determination method. The
microstructural porosity, obtained from conventional 2D
image analysis, tends to be lower than the geometrical
porosity. This may be explained by the fact that micro-
structural porosity is determined on a limited number of
images (which may be considered also as limited number
of slices throughout the sample), while the geometrical
porosity is a bulk property. To minimize this discrepancy
one has to increase the number of images for 2D micro-
structure analysis. Similarly, approximate numbers are
also determined for the phase fraction values (Ni and
SDC) and all the morphological values (Feret-values, –
F-CIRCLE and d). An interesting phenomenon may be
deduced from the contiguity values’ calculation. Namely,
the relatively low CNi-Ni value may indicate a poor contact
between the Ni grains. Such poor contact is highly
undesired, since it also results in a lower electrical
conductivity. However, the relatively high Ni content in
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782 779
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Particle size distribution of ungrounded powder (a), powder
after 1st milling cycle (b) and after 2nd milling cycle (c)
Figure 6: Sintering curve of NiO-SDC material calcined at 800 °C
(full line) and derivative sintering curve (dashed line)
the sample (Ni = 37.7 %) and its homogeneous dis-
tribution throughout the cermet should guarantee a good
Ni-to-Ni particle contact. The latter suggests that a
contiguity determination should consider a real 3D
microstructure, since contacts between particles may
continue also in the direction perpendicular to the
investigated 2D plane. Generally, it is clear that after the
reduction the average particle sizes remain well beneath
the micrometre range; however, more precise micro-
structural quantization requires a more sophisticated
approach.
More precise quantitative microstructural evaluation
is possible by using the real 3D microstructure recon-
struction based on FIB-SEM serial sectioning technique.
It has to be stressed that for 3D microstructure recon-
struction the back-scattered electron detector (BSC-ED)
is more appropriate than the commonly used in-lens
secondary electron detector (IL-SED), since BSC-ED
successfully copes with the shadowing effect, which
sometimes appears on the images when using IL-SED.
Considering the obtained phase-extraction images
(Figure 8), it is evident that the distribution of all three
phases, namely, Ni, SDC and pores (pores may be
considered as an additional phase in the cermet), is
homogeneous. No larger areas of one-phase domination
can be found in the cermet. Moreover, average particle
sizes and pore diameters in any direction are in the
nanometre range. Similar to the 2D analytical approach,
the 3D microstructural description also allows the deter-
mination of morphological parameters (Table 1);
however, the 3D analysis additionally enables a descrip-
tion of microstructure also in the z axis. It appears that
the d and Feret-values (x, y and z) are somewhat lower
when determined by the 3D approach. There are two
main reasons for such a slight discrepancy in the results.
Firstly, the computational algorithms employed for grey
image segmentations and thus the separation of phases
are different; and secondly, 2D statistical calculation is
based on several hundreds of analysed grains, while the
3D statistics involves several tens of thousands of re-
cognized particles. This fact makes the 3D microstruc-
tural analysis more reliable. Similar is true also for the
Ni and SDC determination.
One of the big advantages of 3D microstructural
analyses over the 2D approach is the fact that 3D
microstructure reconstruction also enables an in-depth
porosity description. Due to the large voxel matrix
(2.6×108 voxels) in 3D reconstruction, the determined
value for the microstructural porosity comes very close
to the determined value of the geometrical porosity. The
discrepancy of the geometrical and microstructural is much
bigger for the 2D approach. Moreover, a detailed exami-
nation of the separated pore phase enables the overall
porosity microstructural to be separated into open- and
closed-porosity. According to the results only 0.7% of
the porosity belongs to closed pores, making the pre-
pared microstructure interesting for SOFC applications.
The contiguity values may no longer be calculated
from the 3D microstructure reconstruction. Instead,
using an appropriate computational algorithm the much
more interesting TPB density value (TPBL) can be
obtained. In the analysed Ni-SDC cermet the TPBL is
determined as 7.23 μm μm–3, making this material fully
comparable to other Ni-based cermets described in the
literature. Further interesting information that may also
be obtained from 3D microstructural reconstruction is
the specific surface areas SSA of the phases (i.e., surface
areas of individual phases normalized to the sample
volume). For the investigated cermet the SSA values for
Ni- SDC and pores are determined 15.4 μm2 μm–3,
18.7 μm2 μm–3, 15.7 μm2 μm–3, respectively. The obtained
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
780 Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: 3D microstructure reconstruction of Ni-SDC cermet and
extracted phases Ni, SDC and pores
Figure 7: Microstructures of a) sintered (1400 °C) and b) reduced
(900 °C)
values are rather high compared to some data available in
the literature and are a consequence of the relatively
small grains of each individual phase. Furthermore, each
phase in the cermet is in contact with the two other
phases, meaning that by simple combination of the
obtained SSA values specific interface areas SIA
(normalized to the sample volume) between the three
neighbouring phases can be calculated. For instance, the
SSA of the Ni phase is divided into Ni-pore SIA
6.2 μm2 μm–3 and Ni-SDC SIA 9.2 μm2 μm–3. Thus, about
40 % of the total Ni surface is exposed to the porous
phase and can be used for surface catalytic reaction with
fuel. Namely, oxidation of the fuel involves numerous
reaction steps, which are not necessarily located at the
TPB (i.e., adsorption and dissociation of hydrogen or
hydrocarbon reforming inside an anode layer of a
SOFC).
4 CONCLUSIONS
This work focused on a 2D and 3D microstructural
analysis of a Ni-SDC anode cermet material using SEM
analysis and FIB tomography, respectively. The Ni-SDC
sample was prepared by citrate-nitrate combustion
synthesis, followed by milling, sintering and NiO to Ni
reduction. The 2D microstructure investigation employed
grey image-analysis and made it possible to obtain
quantitative values of the fraction of phases, microstruc-
tural porosity, grain sizes of all phases, sphericity and
contiguity. On the other hand, the 3D analysis addition-
ally provided information on specific surface areas and
specific interface areas of all the phases, as well as the
triple phase boundary density. When comparing the 2D
and the 3D microstructure analysis, the latter turned out
to be more challenging to perform, difficult to interpret
and extra time consuming in the initial step when
quantification algorithms are established. Due to the fact
that 3D microstructural analysis takes into the calcula-
tion a huge number of individual particles, the obtained
information is also more reliable than in the case of 2D
analysis. However, the biggest advantage of 3D
microstructure analysis is the possibility of a complete
microstructural reconstruction and determination of
interface connected properties that directly determine the
cermet activity in an operating SOFC.
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G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782 781
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
the sample (Ni = 37.7 %) and its homogeneous dis-
tribution throughout the cermet should guarantee a good
Ni-to-Ni particle contact. The latter suggests that a
contiguity determination should consider a real 3D
microstructure, since contacts between particles may
continue also in the direction perpendicular to the
investigated 2D plane. Generally, it is clear that after the
reduction the average particle sizes remain well beneath
the micrometre range; however, more precise micro-
structural quantization requires a more sophisticated
approach.
More precise quantitative microstructural evaluation
is possible by using the real 3D microstructure recon-
struction based on FIB-SEM serial sectioning technique.
It has to be stressed that for 3D microstructure recon-
struction the back-scattered electron detector (BSC-ED)
is more appropriate than the commonly used in-lens
secondary electron detector (IL-SED), since BSC-ED
successfully copes with the shadowing effect, which
sometimes appears on the images when using IL-SED.
Considering the obtained phase-extraction images
(Figure 8), it is evident that the distribution of all three
phases, namely, Ni, SDC and pores (pores may be
considered as an additional phase in the cermet), is
homogeneous. No larger areas of one-phase domination
can be found in the cermet. Moreover, average particle
sizes and pore diameters in any direction are in the
nanometre range. Similar to the 2D analytical approach,
the 3D microstructural description also allows the deter-
mination of morphological parameters (Table 1);
however, the 3D analysis additionally enables a descrip-
tion of microstructure also in the z axis. It appears that
the d and Feret-values (x, y and z) are somewhat lower
when determined by the 3D approach. There are two
main reasons for such a slight discrepancy in the results.
Firstly, the computational algorithms employed for grey
image segmentations and thus the separation of phases
are different; and secondly, 2D statistical calculation is
based on several hundreds of analysed grains, while the
3D statistics involves several tens of thousands of re-
cognized particles. This fact makes the 3D microstruc-
tural analysis more reliable. Similar is true also for the
Ni and SDC determination.
One of the big advantages of 3D microstructural
analyses over the 2D approach is the fact that 3D
microstructure reconstruction also enables an in-depth
porosity description. Due to the large voxel matrix
(2.6×108 voxels) in 3D reconstruction, the determined
value for the microstructural porosity comes very close
to the determined value of the geometrical porosity. The
discrepancy of the geometrical and microstructural is much
bigger for the 2D approach. Moreover, a detailed exami-
nation of the separated pore phase enables the overall
porosity microstructural to be separated into open- and
closed-porosity. According to the results only 0.7% of
the porosity belongs to closed pores, making the pre-
pared microstructure interesting for SOFC applications.
The contiguity values may no longer be calculated
from the 3D microstructure reconstruction. Instead,
using an appropriate computational algorithm the much
more interesting TPB density value (TPBL) can be
obtained. In the analysed Ni-SDC cermet the TPBL is
determined as 7.23 μm μm–3, making this material fully
comparable to other Ni-based cermets described in the
literature. Further interesting information that may also
be obtained from 3D microstructural reconstruction is
the specific surface areas SSA of the phases (i.e., surface
areas of individual phases normalized to the sample
volume). For the investigated cermet the SSA values for
Ni- SDC and pores are determined 15.4 μm2 μm–3,
18.7 μm2 μm–3, 15.7 μm2 μm–3, respectively. The obtained
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
780 Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: 3D microstructure reconstruction of Ni-SDC cermet and
extracted phases Ni, SDC and pores
Figure 7: Microstructures of a) sintered (1400 °C) and b) reduced
(900 °C)
values are rather high compared to some data available in
the literature and are a consequence of the relatively
small grains of each individual phase. Furthermore, each
phase in the cermet is in contact with the two other
phases, meaning that by simple combination of the
obtained SSA values specific interface areas SIA
(normalized to the sample volume) between the three
neighbouring phases can be calculated. For instance, the
SSA of the Ni phase is divided into Ni-pore SIA
6.2 μm2 μm–3 and Ni-SDC SIA 9.2 μm2 μm–3. Thus, about
40 % of the total Ni surface is exposed to the porous
phase and can be used for surface catalytic reaction with
fuel. Namely, oxidation of the fuel involves numerous
reaction steps, which are not necessarily located at the
TPB (i.e., adsorption and dissociation of hydrogen or
hydrocarbon reforming inside an anode layer of a
SOFC).
4 CONCLUSIONS
This work focused on a 2D and 3D microstructural
analysis of a Ni-SDC anode cermet material using SEM
analysis and FIB tomography, respectively. The Ni-SDC
sample was prepared by citrate-nitrate combustion
synthesis, followed by milling, sintering and NiO to Ni
reduction. The 2D microstructure investigation employed
grey image-analysis and made it possible to obtain
quantitative values of the fraction of phases, microstruc-
tural porosity, grain sizes of all phases, sphericity and
contiguity. On the other hand, the 3D analysis addition-
ally provided information on specific surface areas and
specific interface areas of all the phases, as well as the
triple phase boundary density. When comparing the 2D
and the 3D microstructure analysis, the latter turned out
to be more challenging to perform, difficult to interpret
and extra time consuming in the initial step when
quantification algorithms are established. Due to the fact
that 3D microstructural analysis takes into the calcula-
tion a huge number of individual particles, the obtained
information is also more reliable than in the case of 2D
analysis. However, the biggest advantage of 3D
microstructure analysis is the possibility of a complete
microstructural reconstruction and determination of
interface connected properties that directly determine the
cermet activity in an operating SOFC.
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G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782 781
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
21 Z. Wu, M. Liu, Modelling of ambipolar transport properties of
composite mixed ionic-electronic conductors, Solid State Ionics, 93
(1996), 65–84, doi:10.1016/S0167-2738(96)00521-8
22 M. Marin{ek, S. Pejovnik, M. Ma~ek, Modelling of Electrical
Properties of Ni-YSZ Composites, J. Eur. Ceram. Soc., 27 (2007),
959–964, doi:10.1016/j.jeurceramsoc.2006.04.165
23 J. A. Quiblier, A new three-dimensional modeling technique for
studying porous media, J. Colloid Interface Sci., 98 (1984), 84–102,
doi:10.1016/0021-9797(84)90481-8
24 H. S. Hong, U. S. Chae, S. T. Choo, The effect of ball milling
parameters and Ni concentration on a YSZ-coated Ni composite for a
high temperature electrolysis cathode, J. Alloys Compd., 449 (2008),
331–334, doi:10.1016/j.jallcom.2006.01.131
25 J. R. Wilson, S. A. Barnett, Solid Oxide Fuel Cell Ni–YSZ Anodes:
Effect of Composition on Microstructure and Performance,
Electrochem. Solid-State Lett., 11 (2008), B181–B185, doi:10.1149/
1.2960528
26 J. R. Wilson, M. Gameiro, K. Mischaikow, W. Kalies, P. W.
Voorhees, S. A. Barnett, hree-Dimensional Analysis of Solid Oxide
Fuel Cell Ni-YSZ Anode Interconnectivity, Microsc. Microanal., 15
(2009), 71–77, doi:10.1017/S1431927609090096
27 J. Gurland, The measurements of grain contiguity in two phase
alloys, Trans. Metall. Soc., 212 (1958), 452-455
28 J. Gurland, An estimate of contact and contiguity of dispersions in
opaque sample, Trans. Metall. Soc., 236 (1966), 642–64
29 J. H. Lee, H. Moon, H. W. Lee, J. Kim, J. D. Kim, K. H. Yoon,
Quantitative analysis of microstructure and its related electrical
property of SOFC anode, Ni-YSZ cermet, Solid State Ionics, 148
(2002), 15–26, doi:10.1016/j.jpowsour.2011.08.083
30 M. Schaffer, J. Wagner, B. Schaffer, M. Schmied, H. Mulders,
Automated three-dimensional X-ray analysis using a dual-beam FIB,
Ultramic., 107 (2007), 587–597, doi:10.1016/j.ultramic.2006.11.007
31 J. M. Perez, J. Pascau, Image Processing with ImageJ, Packt
Publishing Ltd, 2013
32 J. Joos, M. Ender, I. Rotscholl, N. H. Menzler, E. Ivers-Tiffée,
Quantification of double-layer Ni/YSZ fuel cell anodes from focused
ion beam tomography data, J. Power Sources, 246 (2014), 819–830,
doi:10.1016/j.jpowsour.2013.08.021
33 F. Hess, P. Fürnstahl, L. M. Gallo, A. Schweizer, 3D Analysis of the
Proximal Interphalangeal Joint Kinematics during Flexion,
Computational and Mathematical Methods in Medicine, (2013), 7
pages, doi:10.1155/2013/138063
34 A. V. Nagasekhar, C. H. Cáceres, C. Kong, 3D characterization of
intermetallics in a high pressure die cast Mg alloy using focused ion
beam tomography, Mater. Charact., 61(2010), 1035–1042,
doi:10.1016/j.matchar.2010.06.007
35 F. Tariq, R. Haswell, P. D. Lee, D. W. McComb, Characterization of
hierarchical pore structures in ceramics using multiscale tomography,
Acta Mater., 59 (2011), 2109–2120, doi:10.1016/j.actamat.
2010.12.012
36 J. E. Cates, R. T. Whitaker, G. M. Jones, Case study: An evaluation
of user-assisted hierarchical watershed segmentation, Medical Image
Analysis, 9 (2005), 566–578, doi:10.1016/j.media.2005.04.007
37 H. Iwai, N. Shikazono, T. Matsui, H. Teshima, M. Kishimoto, R.
Kishida, H. Yoshida, Quantification of SOFC anode microstructure
based on dual beam FIB-SEM technique, J. Power Sources, 195
(2010), 955–961, doi:10.1016/j.jpowsour.2009.09.005
38 N. Shikazono, D. Kanno, K. Matsuzaki, H. Teshima, S. Sumino, N.
Kasagi, Numerical Assessment of SOFC Anode Polarization Based
on Three-Dimensional Model Microstructure Reconstructed from
FIB-SEM Images, J. Electrochem. Soc., 157 (2010), B665–B672,
doi:10.1149/1.3330568
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
782 Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
H. Y. JIN et al.: A FACILE METHOD TO PREPARE SUPER-HYDROPHOBIC SURFACES ON SILICONE RUBBERS
783–787
A FACILE METHOD TO PREPARE SUPER-HYDROPHOBIC
SURFACES ON SILICONE RUBBERS
PREPROSTA METODA ZA PRIPRAVO SUPERHIDROFOBNIH
POVR[IN PRI SILIKONSKIH GUMAH
Hai Yun Jin1, Yu Feng Li1, Shi Chao Nie1, Peng Zhang1, Nai Kui Gao1, Wen Li2
1Xi’an Jiaotong University, State Key Laboratory of Electrical Insulation and Power Equipment, Xi’an, 710049, China
2Hubei Polytechnic University, Hubei Key Laboratory of Mine Environmental Pollution Control & Remediation and School of Chemical and
Materials Engineering, Huangshi, 435003, PR China
lxfzzl@126.com
Prejem rokopisa – received: 2016-09-02; sprejem za objavo – accepted for publication: 2016-10-24
doi:10.17222/mit.2016.267
Fabricating large-scale super-hydrophobic surfaces for commercial applications is challenging due to certain limitations. In this
paper, a simple and inexpensive method is developed to fabricate super-hydrophobic surfaces on silicone rubbers. Rough micro-
structures were prepared on the mould’s inner surfaces and then sample super-hydrophobic surfaces on silicone rubbers with
different surface roughness were achieved using the standard moulding process. Furthermore, the effects of roughness on the
wettability were investigated. The results showed that by controlling the roughness, the fabricated surfaces exhibited a static
contact angle of 150.9° and a sliding angle of 8°. Finally, the property of hydrophobicity recovery for the silicone-rubber sam-
ples was also studied. The surfaces of the samples could recover well after a sand-blasting experiment. The proposed method is
low-cost, environmentally friendly and suggests promising industrial applications.
Keywords: thin films, coatings, chemical techniques, surface properties
Izdelava obse`nih superhidrofobnih povr{in za komercialne namene je zaradi dolo~enih omejitev zahtevna. V ~lanku je razvit
preprost in poceni na~in za izdelavo superhidrofobne povr{ine na silikonskih gumah. Grobe mikrostrukture so bile pripravljene
na notranjih povr{inah kalupa in nato je bila dose`ena vzor~na superhidrofobna povr{ina na silikonski gumi z razli~no hrapavo
povr{ino z uporabo standardnega postopka modeliranja. Nadalje so bili raziskani u~inki hrapavosti na omo~ljivost. Rezultati so
pokazali, da so s krmiljenjem hrapavosti izdelane povr{ine razstavljene pod stati~nim kontaktnim kotom 150.9° in drsnim kotom
8°. Nazadnje je bila preu~evana tudi lastnost okrevanja hidrofobnosti vzorcev silikonskih gum. Povr{ine vzorcev se lahko dobro
obnovijo tudi po poskusu s peskanjem. Predlagana metoda je poceni, okolju prijazna in ka`e obetavne mo`nosti apliciranja v
procese v industriji.
Klju~ne besede: tanke plasti, prevleke, kemijske tehnike, povr{inske lastnosti
1 INTRODUCTION
Superhydrophobic surfaces, with a water contact
angle (CA) greater than 150° and a sliding angle (SA)
less than 10°, have aroused increasing research interest
for their promising applications.1–6 By observing the
microstructure of a lotus leaf surface, it was found that
the combination of micro papilla and a thin wax film
leads to the self-cleaning properties.4 L. Jiang et al.1 dis-
covered that on top of the micro papilla also exist
branch-like nanostructures and pointed out that the
micro- and nanoscale hierarchical structures were the
fundamental mechanism for lotus leaf’s unique wetting
properties. There were two main measures to fabricate
superhydrophobic surfaces:
1) preparation of a rough surface followed by a low
free-energy material coating step,
2) creation of a rough surface from low-surface-energy
materials.1
Up to now, many approaches to fabricate rough sur-
faces had been developed, including the template
method7, phase separation8, self-assembly9, vapour-phase
deposition10, chemical etching11, laser etching12, hydro-
thermal method13, and electrospinning14. In the power
system, the surfaces of insulators always work under
high voltage and the sand-dust climate condition. How-
ever, some superhydrophobic coats (especially rubber-
based coatings) synthesized by techniques on large-area
and environmentally friendly cannot meet the operating
requirements of composite insulators.15–17 Therefore, it is
a good way to solve this problem by synthesizing a
superhydrophobic surface on the basis of not changing
the insulator material. Silicon rubber was a hydrophobic
material; therefore, a superhydrophobic surface could be
created by only forming a special nm-μm geometry
structure. The template method was a simple and con-
venient technique that could be used as a method to
prepare superhydrophobic surfaces, and the quality of
this method was easy to control; therefore, template
method was suited for industrial production.
In this study, to fabricate a superhydrophobic surface
on the composite insulators, a silicone-rubber super-
hydrophobic surface was developed with a facile tem-
plate method. The creation of a rough structure on the
inner surface of the mould and a special geotroy mor-
phology was prepared on the silicone-rubber surface
Materiali in tehnologije / Materials and technology 51 (2017) 5, 783–787 783
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:621.315.617:544.722.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)783(2017)
21 Z. Wu, M. Liu, Modelling of ambipolar transport properties of
composite mixed ionic-electronic conductors, Solid State Ionics, 93
(1996), 65–84, doi:10.1016/S0167-2738(96)00521-8
22 M. Marin{ek, S. Pejovnik, M. Ma~ek, Modelling of Electrical
Properties of Ni-YSZ Composites, J. Eur. Ceram. Soc., 27 (2007),
959–964, doi:10.1016/j.jeurceramsoc.2006.04.165
23 J. A. Quiblier, A new three-dimensional modeling technique for
studying porous media, J. Colloid Interface Sci., 98 (1984), 84–102,
doi:10.1016/0021-9797(84)90481-8
24 H. S. Hong, U. S. Chae, S. T. Choo, The effect of ball milling
parameters and Ni concentration on a YSZ-coated Ni composite for a
high temperature electrolysis cathode, J. Alloys Compd., 449 (2008),
331–334, doi:10.1016/j.jallcom.2006.01.131
25 J. R. Wilson, S. A. Barnett, Solid Oxide Fuel Cell Ni–YSZ Anodes:
Effect of Composition on Microstructure and Performance,
Electrochem. Solid-State Lett., 11 (2008), B181–B185, doi:10.1149/
1.2960528
26 J. R. Wilson, M. Gameiro, K. Mischaikow, W. Kalies, P. W.
Voorhees, S. A. Barnett, hree-Dimensional Analysis of Solid Oxide
Fuel Cell Ni-YSZ Anode Interconnectivity, Microsc. Microanal., 15
(2009), 71–77, doi:10.1017/S1431927609090096
27 J. Gurland, The measurements of grain contiguity in two phase
alloys, Trans. Metall. Soc., 212 (1958), 452-455
28 J. Gurland, An estimate of contact and contiguity of dispersions in
opaque sample, Trans. Metall. Soc., 236 (1966), 642–64
29 J. H. Lee, H. Moon, H. W. Lee, J. Kim, J. D. Kim, K. H. Yoon,
Quantitative analysis of microstructure and its related electrical
property of SOFC anode, Ni-YSZ cermet, Solid State Ionics, 148
(2002), 15–26, doi:10.1016/j.jpowsour.2011.08.083
30 M. Schaffer, J. Wagner, B. Schaffer, M. Schmied, H. Mulders,
Automated three-dimensional X-ray analysis using a dual-beam FIB,
Ultramic., 107 (2007), 587–597, doi:10.1016/j.ultramic.2006.11.007
31 J. M. Perez, J. Pascau, Image Processing with ImageJ, Packt
Publishing Ltd, 2013
32 J. Joos, M. Ender, I. Rotscholl, N. H. Menzler, E. Ivers-Tiffée,
Quantification of double-layer Ni/YSZ fuel cell anodes from focused
ion beam tomography data, J. Power Sources, 246 (2014), 819–830,
doi:10.1016/j.jpowsour.2013.08.021
33 F. Hess, P. Fürnstahl, L. M. Gallo, A. Schweizer, 3D Analysis of the
Proximal Interphalangeal Joint Kinematics during Flexion,
Computational and Mathematical Methods in Medicine, (2013), 7
pages, doi:10.1155/2013/138063
34 A. V. Nagasekhar, C. H. Cáceres, C. Kong, 3D characterization of
intermetallics in a high pressure die cast Mg alloy using focused ion
beam tomography, Mater. Charact., 61(2010), 1035–1042,
doi:10.1016/j.matchar.2010.06.007
35 F. Tariq, R. Haswell, P. D. Lee, D. W. McComb, Characterization of
hierarchical pore structures in ceramics using multiscale tomography,
Acta Mater., 59 (2011), 2109–2120, doi:10.1016/j.actamat.
2010.12.012
36 J. E. Cates, R. T. Whitaker, G. M. Jones, Case study: An evaluation
of user-assisted hierarchical watershed segmentation, Medical Image
Analysis, 9 (2005), 566–578, doi:10.1016/j.media.2005.04.007
37 H. Iwai, N. Shikazono, T. Matsui, H. Teshima, M. Kishimoto, R.
Kishida, H. Yoshida, Quantification of SOFC anode microstructure
based on dual beam FIB-SEM technique, J. Power Sources, 195
(2010), 955–961, doi:10.1016/j.jpowsour.2009.09.005
38 N. Shikazono, D. Kanno, K. Matsuzaki, H. Teshima, S. Sumino, N.
Kasagi, Numerical Assessment of SOFC Anode Polarization Based
on Three-Dimensional Model Microstructure Reconstructed from
FIB-SEM Images, J. Electrochem. Soc., 157 (2010), B665–B672,
doi:10.1149/1.3330568
G. KAPUN et al.: MICROSTRUCTURAL EVALUATION OF Ni-SDC CERMET FROM A REPRESENTATIVE 2D IMAGE ...
782 Materiali in tehnologije / Materials and technology 51 (2017) 5, 775–782
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
H. Y. JIN et al.: A FACILE METHOD TO PREPARE SUPER-HYDROPHOBIC SURFACES ON SILICONE RUBBERS
783–787
A FACILE METHOD TO PREPARE SUPER-HYDROPHOBIC
SURFACES ON SILICONE RUBBERS
PREPROSTA METODA ZA PRIPRAVO SUPERHIDROFOBNIH
POVR[IN PRI SILIKONSKIH GUMAH
Hai Yun Jin1, Yu Feng Li1, Shi Chao Nie1, Peng Zhang1, Nai Kui Gao1, Wen Li2
1Xi’an Jiaotong University, State Key Laboratory of Electrical Insulation and Power Equipment, Xi’an, 710049, China
2Hubei Polytechnic University, Hubei Key Laboratory of Mine Environmental Pollution Control & Remediation and School of Chemical and
Materials Engineering, Huangshi, 435003, PR China
lxfzzl@126.com
Prejem rokopisa – received: 2016-09-02; sprejem za objavo – accepted for publication: 2016-10-24
doi:10.17222/mit.2016.267
Fabricating large-scale super-hydrophobic surfaces for commercial applications is challenging due to certain limitations. In this
paper, a simple and inexpensive method is developed to fabricate super-hydrophobic surfaces on silicone rubbers. Rough micro-
structures were prepared on the mould’s inner surfaces and then sample super-hydrophobic surfaces on silicone rubbers with
different surface roughness were achieved using the standard moulding process. Furthermore, the effects of roughness on the
wettability were investigated. The results showed that by controlling the roughness, the fabricated surfaces exhibited a static
contact angle of 150.9° and a sliding angle of 8°. Finally, the property of hydrophobicity recovery for the silicone-rubber sam-
ples was also studied. The surfaces of the samples could recover well after a sand-blasting experiment. The proposed method is
low-cost, environmentally friendly and suggests promising industrial applications.
Keywords: thin films, coatings, chemical techniques, surface properties
Izdelava obse`nih superhidrofobnih povr{in za komercialne namene je zaradi dolo~enih omejitev zahtevna. V ~lanku je razvit
preprost in poceni na~in za izdelavo superhidrofobne povr{ine na silikonskih gumah. Grobe mikrostrukture so bile pripravljene
na notranjih povr{inah kalupa in nato je bila dose`ena vzor~na superhidrofobna povr{ina na silikonski gumi z razli~no hrapavo
povr{ino z uporabo standardnega postopka modeliranja. Nadalje so bili raziskani u~inki hrapavosti na omo~ljivost. Rezultati so
pokazali, da so s krmiljenjem hrapavosti izdelane povr{ine razstavljene pod stati~nim kontaktnim kotom 150.9° in drsnim kotom
8°. Nazadnje je bila preu~evana tudi lastnost okrevanja hidrofobnosti vzorcev silikonskih gum. Povr{ine vzorcev se lahko dobro
obnovijo tudi po poskusu s peskanjem. Predlagana metoda je poceni, okolju prijazna in ka`e obetavne mo`nosti apliciranja v
procese v industriji.
Klju~ne besede: tanke plasti, prevleke, kemijske tehnike, povr{inske lastnosti
1 INTRODUCTION
Superhydrophobic surfaces, with a water contact
angle (CA) greater than 150° and a sliding angle (SA)
less than 10°, have aroused increasing research interest
for their promising applications.1–6 By observing the
microstructure of a lotus leaf surface, it was found that
the combination of micro papilla and a thin wax film
leads to the self-cleaning properties.4 L. Jiang et al.1 dis-
covered that on top of the micro papilla also exist
branch-like nanostructures and pointed out that the
micro- and nanoscale hierarchical structures were the
fundamental mechanism for lotus leaf’s unique wetting
properties. There were two main measures to fabricate
superhydrophobic surfaces:
1) preparation of a rough surface followed by a low
free-energy material coating step,
2) creation of a rough surface from low-surface-energy
materials.1
Up to now, many approaches to fabricate rough sur-
faces had been developed, including the template
method7, phase separation8, self-assembly9, vapour-phase
deposition10, chemical etching11, laser etching12, hydro-
thermal method13, and electrospinning14. In the power
system, the surfaces of insulators always work under
high voltage and the sand-dust climate condition. How-
ever, some superhydrophobic coats (especially rubber-
based coatings) synthesized by techniques on large-area
and environmentally friendly cannot meet the operating
requirements of composite insulators.15–17 Therefore, it is
a good way to solve this problem by synthesizing a
superhydrophobic surface on the basis of not changing
the insulator material. Silicon rubber was a hydrophobic
material; therefore, a superhydrophobic surface could be
created by only forming a special nm-μm geometry
structure. The template method was a simple and con-
venient technique that could be used as a method to
prepare superhydrophobic surfaces, and the quality of
this method was easy to control; therefore, template
method was suited for industrial production.
In this study, to fabricate a superhydrophobic surface
on the composite insulators, a silicone-rubber super-
hydrophobic surface was developed with a facile tem-
plate method. The creation of a rough structure on the
inner surface of the mould and a special geotroy mor-
phology was prepared on the silicone-rubber surface
Materiali in tehnologije / Materials and technology 51 (2017) 5, 783–787 783
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:621.315.617:544.722.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)783(2017)
through a conventional moulding process. The surface
microstructures were investigated by a roughness tester
and scanning electron microscope (SEM). It was found
that the superhydrophobic silicone-rubber surface could
be attained by creating the appropriate roughness sur-
faces, which lead to the lotus-leaf-like surface morphol-
ogy. After samples suffered the artificial sandstorm
testing, the property of hydrophobicity recovery for
superhydrophobic and a common silicone-rubber surface
have also been studied in this paper.
2 EXPERIMENTAL PART
Templates with different rough structures were pre-
pared first. Silicon carbide particles of (63, 21, 15, and
10.5) μm in diameter were obtained by using test sieves
of 80, 240, 600, 800 and 1200 mesh, respectively. A uni-
form layer of resin binder-epoxy resin was sprayed on
the inner surface of a cubic mould with dimensions of 3
cm × 3 cm × 3 cm after treating with a silane coupling
agent N--aminoethyl--aminopropyl trimethoxysilane.
The prepared silicon carbide particles were evenly
sprayed on the inner surface of the mould and dried at
room temperature. Then rough silicon rubber samples
were prepared through a conventional moulding process.
Liquid silicone rubber was poured into the mould, the
mould was taken off after consolidation, and thus sili-
cone-rubber surfaces with different surface morphologies
were obtained. To avoid the superhydrophobic surface
being destroyed during the mould unloading, the moulds
were treated with a release agent(methyl-silicone oil).
For a comparison, the smooth silicone rubber was
prepared with a smooth mould.
To analyses the morphology and microstructures of
the sample surfaces, SEM and a roughness tester were
used. The morphological characterization was observed
with a SEM (VE-9800, Keyence). The surface roughness
Ra was measured with a handhold roughness tester
(TR200, Times Run Bao). The roughness tester was
placed at 20 different locations of the sample surface un-
der investigation and the average value was taken as the
surface roughness.
The water contact angles were investigated using a
measurement system (JC2000C4, POWEREACH). Be-
fore the test the samples were washed with deionized
water and ethanol, respectively, and then naturally dried.
Five different locations of the sample surface were mea-
sured, and the average value was used.
The property of hydrophobicity recovery for both the
superhydrophobic and the common silicone-rubber sur-
face were tested in the artificial experimental platform,
and the sand-dust sample was also prepared, as described
in 18. Samples of superhydrophobic silicone rubber and
the common silicone rubber were placed horizontally in
the experimental platform. The surfaces of the samples
and the wind direction were parallel. The wind speed
was 15m/s. The samples were in the artificial sandstorm
for 1 h. To accelerate the process of hydrophobicity re-
covery, the heat treatment for the sample was carried out
in an oven at 75 °C.
3 RESULTS AND DISCUSSION
Figure 1a to 1f show SEM micrographs of surface
morphologies for the different silicone-rubber surfaces.
A very flat microstructure can be seen from the surface
prepared with the smooth template (Figure 1a). After
preparing with templates, the surfaces have spongy-like
rough structures and this confirmed that the morphology
of the mould surface was successfully replicated. It can
be seen that rough structure units (number and decrease
in scale) increase with decreasing silicon carbide particle
size. The silicone-rubber surface prepared with a tem-
plate of particle size 63 μm has an average diameter of
200 μm irregular protuberances and pits (Figure 1b), the
surface distributes on average a diameter of 100 μm ir-
regular papilla (Figure 1c), and the surface prepared
with template of particle size of 15 μm in diameter has a
diameter of about 40 μm irregular pits and protuberances
(Figure 1d). Figure 1e to 1f are higher-resolution
micrographs of Figure 1d. Figure 1f and show that the
randomly distributed papilla morphology (observed in
Figure 1e) consists of sub-micron structures, which
implies two-length-scale hierarchical structure on the
surface. Therefore, the Cassie-Baxter superhydrophobic
state was well formed.
H. Y. JIN et al.: A FACILE METHOD TO PREPARE SUPER-HYDROPHOBIC SURFACES ON SILICONE RUBBERS
784 Materiali in tehnologije / Materials and technology 51 (2017) 5, 783–787
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: SEM micrographs of surface prepared with templates of dif-
ferent diameter silicon carbide particles: a) smooth, b) 63 μm,
c) 21 μm, d) 15 μm, e) the high magnification of d), and f) the high
magnification of (e)
Figure 2 shows the roughness of the surface mor-
phology. The surface roughness of the smooth silicone
rubber surface was 0.42 μm. When the templates are
made by different silicon carbide particles of (10.5, 15,
21, and 63) μm in diameter, the surface roughness of the
samples made by different templates are (10.88, 14.63,
17.32 and 33.61) μm, respectively. The results show that,
after roughening with a rough template, the surface
roughness of the silicone rubber increases with the in-
creasing particle size of the sprayed silicon carbide.
Figure 3 shows the relationship between the surface
wettability and the surface roughness. It is clear that the
contact angle of the smooth silicone rubber surface is
110.2° and the sliding angle is about 90°. While rough-
ening with rough templates, the surface wettability is
greatly improved (the contact angle is larger than 140°
and the sliding angle is from 8° to 50°). The higher
contact angle of the rough silicone-rubber surface can be
attributed to the formation of the Wenzel wetting state
(Figure 4a). According to Wenzel model (cos * = r cos
, where * is Wenzel apparent contact angle,  is the
intrinsic contact angle of the ideal smooth material,
which also called as Young’s contact angle, and r is the
ratio of the actual area to the projected area), rough
structures can amplify the natural wettability tendency of
the surface. Since silicone rubber is hydrophobic, the
surface hydrophobic property will be improved by
roughening the surface; its new contact angle becomes
greater than the intrinsic contact angle. The contact angle
reaches to 150.9° with a very small sliding angle of 8o
for samples with a surface roughness of 14.63 μm. This
can be explained by the formation of the Cassie-Baxter
wetting state (Figure 4b), for it has a lotus-leaf-like
surface morphology (Figure 1d to 1f). It is because of
that that there is such a micro-submicron hierarchical
structure on samples, which is like the structure of the
lotus leaf.1,4 If a water droplet is placed upon the surface,
air can be entrapped in both the hollows and interstices,
resulting in the formation of a high area fraction of
air-liquid interfaces. According to the Cassie-Baxter
equation (cos a = f1 cos – f2, where a is the apparent
contact angle on a rough surface, f1 is the area fraction of
the solid-liquid interface, and f2 is area fraction of the
air-liquid interface), the contact angle can larger than
150° when the rough surface attains a large area fraction
of air-liquid interface. With regards to the sliding angle,
a surface exhibiting the Wenzel wetting state usually
shows a larger sliding angle due to the pinning effects
(Figure 4a). In contrast, a surface following the regime
of the Cassie–Baxter wetting state (Figure 4b) allows a
water droplet to roll off easily with a sliding angle of less
than 10°.
Figure 3 shows that surface roughness plays an im-
portant role in the surface wettability. Since all the sili-
cone rubber samples are prepared with a similar template
structure (but different roughness), the samples present a
similar surface morphology (Figure 1b to 1d). For
samples with surface roughness Ra of 10.88 μm and
14.63 μm, the contact angle exceeds 150°, with a negli-
gible sliding angle of about 8°. The contact angles
reaches to a maximum value of 151.8°, when the surface
roughness is 10.88 μm. When the surface roughness is
larger than 10.88 μm, the contact angles decrease with
increasing surface roughness. For the samples with a
surface roughness (Ra) of 17.32 μm and 33.61 μm, the
contact angles are 146.9° and 140.8° respectively, while
both of them show a sliding angle of about 50°. It is
known that the micro-scale and nano-scale hierarchical
structures affect the contact angle greatly. The micro-
scale and nano-scale hierarchical structures can be
H. Y. JIN et al.: A FACILE METHOD TO PREPARE SUPER-HYDROPHOBIC SURFACES ON SILICONE RUBBERS
Materiali in tehnologije / Materials and technology 51 (2017) 5, 783–787 785
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Various wetting state model sketch of a droplet on a rough
surface: a) Wezel state, b) Cassie-Baxer state and c) combined
Cassie-Baxter/Wenzel state
Figure 3: Relationship between static contact angle and surface
roughness
Figure 2: Relationship between particle size and surface roughness Ra
through a conventional moulding process. The surface
microstructures were investigated by a roughness tester
and scanning electron microscope (SEM). It was found
that the superhydrophobic silicone-rubber surface could
be attained by creating the appropriate roughness sur-
faces, which lead to the lotus-leaf-like surface morphol-
ogy. After samples suffered the artificial sandstorm
testing, the property of hydrophobicity recovery for
superhydrophobic and a common silicone-rubber surface
have also been studied in this paper.
2 EXPERIMENTAL PART
Templates with different rough structures were pre-
pared first. Silicon carbide particles of (63, 21, 15, and
10.5) μm in diameter were obtained by using test sieves
of 80, 240, 600, 800 and 1200 mesh, respectively. A uni-
form layer of resin binder-epoxy resin was sprayed on
the inner surface of a cubic mould with dimensions of 3
cm × 3 cm × 3 cm after treating with a silane coupling
agent N--aminoethyl--aminopropyl trimethoxysilane.
The prepared silicon carbide particles were evenly
sprayed on the inner surface of the mould and dried at
room temperature. Then rough silicon rubber samples
were prepared through a conventional moulding process.
Liquid silicone rubber was poured into the mould, the
mould was taken off after consolidation, and thus sili-
cone-rubber surfaces with different surface morphologies
were obtained. To avoid the superhydrophobic surface
being destroyed during the mould unloading, the moulds
were treated with a release agent(methyl-silicone oil).
For a comparison, the smooth silicone rubber was
prepared with a smooth mould.
To analyses the morphology and microstructures of
the sample surfaces, SEM and a roughness tester were
used. The morphological characterization was observed
with a SEM (VE-9800, Keyence). The surface roughness
Ra was measured with a handhold roughness tester
(TR200, Times Run Bao). The roughness tester was
placed at 20 different locations of the sample surface un-
der investigation and the average value was taken as the
surface roughness.
The water contact angles were investigated using a
measurement system (JC2000C4, POWEREACH). Be-
fore the test the samples were washed with deionized
water and ethanol, respectively, and then naturally dried.
Five different locations of the sample surface were mea-
sured, and the average value was used.
The property of hydrophobicity recovery for both the
superhydrophobic and the common silicone-rubber sur-
face were tested in the artificial experimental platform,
and the sand-dust sample was also prepared, as described
in 18. Samples of superhydrophobic silicone rubber and
the common silicone rubber were placed horizontally in
the experimental platform. The surfaces of the samples
and the wind direction were parallel. The wind speed
was 15m/s. The samples were in the artificial sandstorm
for 1 h. To accelerate the process of hydrophobicity re-
covery, the heat treatment for the sample was carried out
in an oven at 75 °C.
3 RESULTS AND DISCUSSION
Figure 1a to 1f show SEM micrographs of surface
morphologies for the different silicone-rubber surfaces.
A very flat microstructure can be seen from the surface
prepared with the smooth template (Figure 1a). After
preparing with templates, the surfaces have spongy-like
rough structures and this confirmed that the morphology
of the mould surface was successfully replicated. It can
be seen that rough structure units (number and decrease
in scale) increase with decreasing silicon carbide particle
size. The silicone-rubber surface prepared with a tem-
plate of particle size 63 μm has an average diameter of
200 μm irregular protuberances and pits (Figure 1b), the
surface distributes on average a diameter of 100 μm ir-
regular papilla (Figure 1c), and the surface prepared
with template of particle size of 15 μm in diameter has a
diameter of about 40 μm irregular pits and protuberances
(Figure 1d). Figure 1e to 1f are higher-resolution
micrographs of Figure 1d. Figure 1f and show that the
randomly distributed papilla morphology (observed in
Figure 1e) consists of sub-micron structures, which
implies two-length-scale hierarchical structure on the
surface. Therefore, the Cassie-Baxter superhydrophobic
state was well formed.
H. Y. JIN et al.: A FACILE METHOD TO PREPARE SUPER-HYDROPHOBIC SURFACES ON SILICONE RUBBERS
784 Materiali in tehnologije / Materials and technology 51 (2017) 5, 783–787
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: SEM micrographs of surface prepared with templates of dif-
ferent diameter silicon carbide particles: a) smooth, b) 63 μm,
c) 21 μm, d) 15 μm, e) the high magnification of d), and f) the high
magnification of (e)
Figure 2 shows the roughness of the surface mor-
phology. The surface roughness of the smooth silicone
rubber surface was 0.42 μm. When the templates are
made by different silicon carbide particles of (10.5, 15,
21, and 63) μm in diameter, the surface roughness of the
samples made by different templates are (10.88, 14.63,
17.32 and 33.61) μm, respectively. The results show that,
after roughening with a rough template, the surface
roughness of the silicone rubber increases with the in-
creasing particle size of the sprayed silicon carbide.
Figure 3 shows the relationship between the surface
wettability and the surface roughness. It is clear that the
contact angle of the smooth silicone rubber surface is
110.2° and the sliding angle is about 90°. While rough-
ening with rough templates, the surface wettability is
greatly improved (the contact angle is larger than 140°
and the sliding angle is from 8° to 50°). The higher
contact angle of the rough silicone-rubber surface can be
attributed to the formation of the Wenzel wetting state
(Figure 4a). According to Wenzel model (cos * = r cos
, where * is Wenzel apparent contact angle,  is the
intrinsic contact angle of the ideal smooth material,
which also called as Young’s contact angle, and r is the
ratio of the actual area to the projected area), rough
structures can amplify the natural wettability tendency of
the surface. Since silicone rubber is hydrophobic, the
surface hydrophobic property will be improved by
roughening the surface; its new contact angle becomes
greater than the intrinsic contact angle. The contact angle
reaches to 150.9° with a very small sliding angle of 8o
for samples with a surface roughness of 14.63 μm. This
can be explained by the formation of the Cassie-Baxter
wetting state (Figure 4b), for it has a lotus-leaf-like
surface morphology (Figure 1d to 1f). It is because of
that that there is such a micro-submicron hierarchical
structure on samples, which is like the structure of the
lotus leaf.1,4 If a water droplet is placed upon the surface,
air can be entrapped in both the hollows and interstices,
resulting in the formation of a high area fraction of
air-liquid interfaces. According to the Cassie-Baxter
equation (cos a = f1 cos – f2, where a is the apparent
contact angle on a rough surface, f1 is the area fraction of
the solid-liquid interface, and f2 is area fraction of the
air-liquid interface), the contact angle can larger than
150° when the rough surface attains a large area fraction
of air-liquid interface. With regards to the sliding angle,
a surface exhibiting the Wenzel wetting state usually
shows a larger sliding angle due to the pinning effects
(Figure 4a). In contrast, a surface following the regime
of the Cassie–Baxter wetting state (Figure 4b) allows a
water droplet to roll off easily with a sliding angle of less
than 10°.
Figure 3 shows that surface roughness plays an im-
portant role in the surface wettability. Since all the sili-
cone rubber samples are prepared with a similar template
structure (but different roughness), the samples present a
similar surface morphology (Figure 1b to 1d). For
samples with surface roughness Ra of 10.88 μm and
14.63 μm, the contact angle exceeds 150°, with a negli-
gible sliding angle of about 8°. The contact angles
reaches to a maximum value of 151.8°, when the surface
roughness is 10.88 μm. When the surface roughness is
larger than 10.88 μm, the contact angles decrease with
increasing surface roughness. For the samples with a
surface roughness (Ra) of 17.32 μm and 33.61 μm, the
contact angles are 146.9° and 140.8° respectively, while
both of them show a sliding angle of about 50°. It is
known that the micro-scale and nano-scale hierarchical
structures affect the contact angle greatly. The micro-
scale and nano-scale hierarchical structures can be
H. Y. JIN et al.: A FACILE METHOD TO PREPARE SUPER-HYDROPHOBIC SURFACES ON SILICONE RUBBERS
Materiali in tehnologije / Materials and technology 51 (2017) 5, 783–787 785
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Various wetting state model sketch of a droplet on a rough
surface: a) Wezel state, b) Cassie-Baxer state and c) combined
Cassie-Baxter/Wenzel state
Figure 3: Relationship between static contact angle and surface
roughness
Figure 2: Relationship between particle size and surface roughness Ra
quantified by the surface roughness; therefore, the effect
of surface structures on the contact angle can be studied
by the surface roughness.19 Compared with the contact
angle, the sliding angle is more affected by the surface
roughness. This can be attributed to the formation of the
combined Cassie-Baxter and Wenzel wetting state model
(Figure 4c).20,21 As seen, the area fraction of air-liquid
interfaces from the combined Cassie-Baxter and Wenzel
wetting state model changes little (compared with
Cassie-Baxter wetting state model); therefore, the con-
tact angle will be larger than 140°. However, the sliding
angle is greatly increased due to a certain extent by the
pinning effects (Figure 4c). With respect to promising
applications, such as anti-icing of superhydrophobic sur-
faces, both the contact angle and the small sliding angle
are very important; therefore, the sliding angle also plays
a important role.22 Superhydrophobic surfaces with
negligible sliding angle can be obtained by controlling
the surface structure and the roughness.
Figure 5 shows the changes of the contact angle after
the sandstorm experiment and the hydrophobicity recov-
ery experiment. Because of the effects of the sand-dust,
both of the contact angles decrease by approximately
20°. At the beginning of the recovering process, the hyd-
rophobicity of samples recover quickly. Then the speed
of the recovery slows down. After 24 h in a high-tem-
perature treatment, the hydrophobicities of both super-
hydrophobic and common silicone rubber have basically
recovered to the level of the primary sample.
The sand-dust on the silicone rubber is hydrophilic,
which results in hydrophobicity degradation of the
samples. Inside the silicone rubber, the large molecules
polysiloxane are cracked into smaller molecules poly-
siloxane. And the high temperature can accelerate the
cracking process. The small molecules polysiloxane with
a low surface tension can migrate to the surface of the
silicone rubber and cover the sand-dust particles.23,24
Therefore, the hydrophobicity can be recovered after the
high-temperature treatment. Considering that the contact
angle of the superhydrophobic silicone rubber has recov-
ered to almost 150°, it shows that the micro-submicron
hierarchical structure of superhydrophobic surfaces can
almost not be destroyed.
4 CONCLUSIONS
Superhydrophobic surfaces were successfully fabri-
cated on silicone rubbers through a facile template
method. The contact angle for the surfaces can reach a
maximum of 151.8° (sliding angle about 8°) with a sur-
face roughness of 10.88 μm. When the surface roughness
is larger than 10.88 μm the contact angle decreased with
increasing roughness, and when the surface roughness is
less than 10.88 μm, the contact angle increased with
increasing roughness. The micro/sub-micron hierarchical
structures of the surface were responsible for the high
water contact angle and the low sliding angle. The
wetting ability of the various surface roughness was
different, and it could be explained by the different
wetting state model. The sandstorm testing and the
hydrophobicity experiment showed that the contact angle
of both superhydrophobic and the common silicone
rubber sample could recover.
Acknowledgments
This research was financially supported by the Na-
tional Natural Science Foundation of China (No.
51272208), the Natural Science Foundation of Hubei
Province (2010CDA026), the Key Program of Hubei
Provincial Department of Education (Z20104401), the
Outstanding Scientific and Technological Innovation
Team Project of Universities in Hubei Province
(T201423), and the Talent Program of Hubei Polytechnic
University (11yjz04R).
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290 (2000) 5499, 2130–2133, doi:10.1126/science.290.5499.2130
10 A. Hozumi, O. Takai, Preparation of ultra water-repellent films by
microwave plasma-enhanced CVD, Thin Solid Films, 303 (1997)
1–2, 222–225, doi:10.1016/S0040-6090(97)00076-X
11 D. Xie, W. Li. A novel simple approach to preparation of super-
hydrophobic surface of aluminum alloys, Applied Surface Science,
258 (2011) 3, 1004–1007, doi: 10.1016/j.apsusc.2011.07.104
12 T. Sun, G. Wang, L. Feng, B. Liu, Y. Ma, L. Jiang, D. Zhu,
Reversible switching between super-hydrophilicity and super-hydro-
phboicity, Angewandte Chemie International Edition, 43 (2004) 3,
357–360, doi:10.1002/anie.200352565
13 X. Feng, L. Feng, M. Jin, J. Zhai, L. Jiang, D. Zhu, Reversible
super-hydrophobicity to super-hydrophilicity transition of aligned
ZnO nanorod films, Journal of the American Chemical Society, 126
(2004) 1, 62–63, doi:10.1021/ja038636o
14 L. Jiang, Y. Zhao, J. Zhai, A lotus-leaf-like superhydrophobic sur-
face: a porous microsphere/nanofiber composite film prepared by
electrohydrodynamics, Angewandte Chemie International Edition,
43 (2004) 33, 4338–4341, doi:10.1002/anie.200460333
15 M. K. Tiwari, I. S. Bayer, G. M. Jursich, T. M. Schutzius, C. M.
Megaridis, Highly Liquid-Repellent, Large-Area, Nanostructured
Poly(vinylidene fluoride)/Poly(ethyl 2-cyanoacrylate) Composite
Coatings: Particle Filler Effects, ACS applied materials & interfaces,
2 (2010) 4, 1114–1119, doi:10.1021/am900894n
16 T. M. Schutzius, I. S. Bayer, J. Qin, D. Waldroup, C. M. Megaridis,
Water-Based, Nonfluorinated Dispersions for Environmentally Be-
nign, Large-Area, Superhydrophobic Coatings, ACS applied mate-
rials & interfaces, 5 (2013) 24, 13419–13425, doi:10.1021/
am4043307
17 E. Bormashenko, V. Goldshtein, R. Barayev, T. Stein, G. Whyman,
R. Pogreb, Z. Barkay, D. Aurbach, Robust method of manufacturing
rubber waste-based water repellent surfaces, Polymers for Advanced
Technologies, 20 (2009) 7, 650–653, doi:10.1002/pat.1323
18 B. He, G. Zhang, B. Chen, N. Gao, Y. Li, Z. Peng, H. Jin, The
influence of the sand-dust environment on air-gap breakdown
discharge characteristics of the plate-to-plate electrode, Science
China Physics, Mechanics and Astronomy, 53 (2010) 3, 458–464,
doi:10.1007/s11433-010-0078-1
19 X. Shi, T. A. Nguyen, Z. Suo, J. Wu, J. Gong, R. Avci, Electrochemi-
cal and mechanical properties of superhydrophobic aluminum sub-
strates modified with nano-silica and fluorosilane, Surface and Coat-
ings Technology, 206 (2012) 17, 3700–3713, doi:10.1016/j.surfcoat.
2012.02.058
20 X. Feng, L. Jiang, Design and creation of superwetting/antiwetting
surfaces, Advanced Material, 18 (2006) 23, 3063–3078, doi:10.1002/
adma.200501961
21 C. W. Yang, F. He, P. F. Hao, The apparent contact angle of water
droplet on the micro-structured hydrophobic surface, Science China
Chemistry, 53 (2010) 4, 912–916, doi:10.1007/s11426-010-0115-y
22 S. A. Kulinich, M. Farzaneh, How Wetting Hysteresis Influences Ice
Adhesion Strength on Superhydrophobic Surfaces. Langmuir, 25
(2009) 16, 8854–8856, doi:10.1021/la901439c
23 J. W. Chang, R. S. Gorur, Surface recovery of silicone rubber used
for HV outdoor insulation, IEEE Transactions on Dielectrics and
Electrical Insulation, 1 (1994) 6, 1039–1046, doi:10.1109/94.368659
24 T. Tokoro, R. Hackam, Loss and recovery of hydrophobicity and sur-
face energy of HTV silicone rubber, IEEE Transactions on Dielec-
trics and Electrical Insulation, 8 (2001) 6, 1088–1097,
doi:10.1109/94.971469
H. Y. JIN et al.: A FACILE METHOD TO PREPARE SUPER-HYDROPHOBIC SURFACES ON SILICONE RUBBERS
Materiali in tehnologije / Materials and technology 51 (2017) 5, 783–787 787
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
quantified by the surface roughness; therefore, the effect
of surface structures on the contact angle can be studied
by the surface roughness.19 Compared with the contact
angle, the sliding angle is more affected by the surface
roughness. This can be attributed to the formation of the
combined Cassie-Baxter and Wenzel wetting state model
(Figure 4c).20,21 As seen, the area fraction of air-liquid
interfaces from the combined Cassie-Baxter and Wenzel
wetting state model changes little (compared with
Cassie-Baxter wetting state model); therefore, the con-
tact angle will be larger than 140°. However, the sliding
angle is greatly increased due to a certain extent by the
pinning effects (Figure 4c). With respect to promising
applications, such as anti-icing of superhydrophobic sur-
faces, both the contact angle and the small sliding angle
are very important; therefore, the sliding angle also plays
a important role.22 Superhydrophobic surfaces with
negligible sliding angle can be obtained by controlling
the surface structure and the roughness.
Figure 5 shows the changes of the contact angle after
the sandstorm experiment and the hydrophobicity recov-
ery experiment. Because of the effects of the sand-dust,
both of the contact angles decrease by approximately
20°. At the beginning of the recovering process, the hyd-
rophobicity of samples recover quickly. Then the speed
of the recovery slows down. After 24 h in a high-tem-
perature treatment, the hydrophobicities of both super-
hydrophobic and common silicone rubber have basically
recovered to the level of the primary sample.
The sand-dust on the silicone rubber is hydrophilic,
which results in hydrophobicity degradation of the
samples. Inside the silicone rubber, the large molecules
polysiloxane are cracked into smaller molecules poly-
siloxane. And the high temperature can accelerate the
cracking process. The small molecules polysiloxane with
a low surface tension can migrate to the surface of the
silicone rubber and cover the sand-dust particles.23,24
Therefore, the hydrophobicity can be recovered after the
high-temperature treatment. Considering that the contact
angle of the superhydrophobic silicone rubber has recov-
ered to almost 150°, it shows that the micro-submicron
hierarchical structure of superhydrophobic surfaces can
almost not be destroyed.
4 CONCLUSIONS
Superhydrophobic surfaces were successfully fabri-
cated on silicone rubbers through a facile template
method. The contact angle for the surfaces can reach a
maximum of 151.8° (sliding angle about 8°) with a sur-
face roughness of 10.88 μm. When the surface roughness
is larger than 10.88 μm the contact angle decreased with
increasing roughness, and when the surface roughness is
less than 10.88 μm, the contact angle increased with
increasing roughness. The micro/sub-micron hierarchical
structures of the surface were responsible for the high
water contact angle and the low sliding angle. The
wetting ability of the various surface roughness was
different, and it could be explained by the different
wetting state model. The sandstorm testing and the
hydrophobicity experiment showed that the contact angle
of both superhydrophobic and the common silicone
rubber sample could recover.
Acknowledgments
This research was financially supported by the Na-
tional Natural Science Foundation of China (No.
51272208), the Natural Science Foundation of Hubei
Province (2010CDA026), the Key Program of Hubei
Provincial Department of Education (Z20104401), the
Outstanding Scientific and Technological Innovation
Team Project of Universities in Hubei Province
(T201423), and the Talent Program of Hubei Polytechnic
University (11yjz04R).
5 REFERENCES
1 L. Feng, S. Li, Y. Li, H. Li, L. Zhang, J. Zhai, Y. Song, B. Liu, L.
Jiang, D. Zhu, Super-hydrophobic surfaces: From natural to artificial,
Advanced Materials, 14 (2002) 24, 1857–1860, doi:10.1002/adma.
200290020
2 K. K. Varanasi, M. Hsu, N. Bhate, W. Yang, T. Deng, Spatial Control
in the Heterogeneous Nucleation of Water, Applied Physics Letters,
95 (2009) 9, 094101, doi:10.1063/1.3200951
3 L. Cao, A. K. Jones, V. K. Sikka, J. Z. Wu, D. Gao, Anti-Icing
Superhydrophobic Coatings, Langmuir, 25 (2009) 21, 12444–12448,
doi:10.1021/la902882b
4 W. Barthlott, C. Neinhuis, Purity of the sacred lotus, or escape from
contamination in biological surfaces, Planta, 202 (1997) 1, 1–8,
doi:10.1007/s004250050096
5 F. Arianpour, M. Farzaneh, S. A. Kulinich, Hydrophobic and ice-re-
tarding properties of doped silicone rubber coatings, Applied Surface
Science, 265 (2013), 546–552, doi:10.1016/j.apsusc.2012.11.042
6 J. Li, Y. Zhao, J. Hu, L. Shu, X. Shi, Anti-icing Performance of a
Superhydrophobic PDMS/Modified Nano-silica Hybrid Coating for
Insulators, Journal of Adhesion Science and Technology, 26 (2012)
4–5, 665–679, doi:10.1163/016942411X574826
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786 Materiali in tehnologije / Materials and technology 51 (2017) 5, 783–787
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: The recovery curve for the hydrophobicity of the sili-
cone-rubber samples
7 M. Jin, M. Liao, J. Zhai, L. Jiang, Super-hydrophobic polystyrene
films prepared via pattern transformation, Acta Chimica Sinica, 66
(2008) 1, 145–148, doi:10.3321/j.issn:0567-7351.2008.01.025
8 Q. Xie, G. Fan, N. Zhao, X. Guo, J. Xu, J. Dong, L. Zhang, Y.
Zhang, C. Han, Facile creation of a bionic super-hydrophobic block
copolymer surface, Advanced Materials, 16 (2004) 20, 1830–1833,
doi:10.1002/adma.200400074
9 J. Genzer, K. Efimenko, Creating long-lived super-hydrophobic poly-
mer surfaces through mechanically assembled monolayers, Science,
290 (2000) 5499, 2130–2133, doi:10.1126/science.290.5499.2130
10 A. Hozumi, O. Takai, Preparation of ultra water-repellent films by
microwave plasma-enhanced CVD, Thin Solid Films, 303 (1997)
1–2, 222–225, doi:10.1016/S0040-6090(97)00076-X
11 D. Xie, W. Li. A novel simple approach to preparation of super-
hydrophobic surface of aluminum alloys, Applied Surface Science,
258 (2011) 3, 1004–1007, doi: 10.1016/j.apsusc.2011.07.104
12 T. Sun, G. Wang, L. Feng, B. Liu, Y. Ma, L. Jiang, D. Zhu,
Reversible switching between super-hydrophilicity and super-hydro-
phboicity, Angewandte Chemie International Edition, 43 (2004) 3,
357–360, doi:10.1002/anie.200352565
13 X. Feng, L. Feng, M. Jin, J. Zhai, L. Jiang, D. Zhu, Reversible
super-hydrophobicity to super-hydrophilicity transition of aligned
ZnO nanorod films, Journal of the American Chemical Society, 126
(2004) 1, 62–63, doi:10.1021/ja038636o
14 L. Jiang, Y. Zhao, J. Zhai, A lotus-leaf-like superhydrophobic sur-
face: a porous microsphere/nanofiber composite film prepared by
electrohydrodynamics, Angewandte Chemie International Edition,
43 (2004) 33, 4338–4341, doi:10.1002/anie.200460333
15 M. K. Tiwari, I. S. Bayer, G. M. Jursich, T. M. Schutzius, C. M.
Megaridis, Highly Liquid-Repellent, Large-Area, Nanostructured
Poly(vinylidene fluoride)/Poly(ethyl 2-cyanoacrylate) Composite
Coatings: Particle Filler Effects, ACS applied materials & interfaces,
2 (2010) 4, 1114–1119, doi:10.1021/am900894n
16 T. M. Schutzius, I. S. Bayer, J. Qin, D. Waldroup, C. M. Megaridis,
Water-Based, Nonfluorinated Dispersions for Environmentally Be-
nign, Large-Area, Superhydrophobic Coatings, ACS applied mate-
rials & interfaces, 5 (2013) 24, 13419–13425, doi:10.1021/
am4043307
17 E. Bormashenko, V. Goldshtein, R. Barayev, T. Stein, G. Whyman,
R. Pogreb, Z. Barkay, D. Aurbach, Robust method of manufacturing
rubber waste-based water repellent surfaces, Polymers for Advanced
Technologies, 20 (2009) 7, 650–653, doi:10.1002/pat.1323
18 B. He, G. Zhang, B. Chen, N. Gao, Y. Li, Z. Peng, H. Jin, The
influence of the sand-dust environment on air-gap breakdown
discharge characteristics of the plate-to-plate electrode, Science
China Physics, Mechanics and Astronomy, 53 (2010) 3, 458–464,
doi:10.1007/s11433-010-0078-1
19 X. Shi, T. A. Nguyen, Z. Suo, J. Wu, J. Gong, R. Avci, Electrochemi-
cal and mechanical properties of superhydrophobic aluminum sub-
strates modified with nano-silica and fluorosilane, Surface and Coat-
ings Technology, 206 (2012) 17, 3700–3713, doi:10.1016/j.surfcoat.
2012.02.058
20 X. Feng, L. Jiang, Design and creation of superwetting/antiwetting
surfaces, Advanced Material, 18 (2006) 23, 3063–3078, doi:10.1002/
adma.200501961
21 C. W. Yang, F. He, P. F. Hao, The apparent contact angle of water
droplet on the micro-structured hydrophobic surface, Science China
Chemistry, 53 (2010) 4, 912–916, doi:10.1007/s11426-010-0115-y
22 S. A. Kulinich, M. Farzaneh, How Wetting Hysteresis Influences Ice
Adhesion Strength on Superhydrophobic Surfaces. Langmuir, 25
(2009) 16, 8854–8856, doi:10.1021/la901439c
23 J. W. Chang, R. S. Gorur, Surface recovery of silicone rubber used
for HV outdoor insulation, IEEE Transactions on Dielectrics and
Electrical Insulation, 1 (1994) 6, 1039–1046, doi:10.1109/94.368659
24 T. Tokoro, R. Hackam, Loss and recovery of hydrophobicity and sur-
face energy of HTV silicone rubber, IEEE Transactions on Dielec-
trics and Electrical Insulation, 8 (2001) 6, 1088–1097,
doi:10.1109/94.971469
H. Y. JIN et al.: A FACILE METHOD TO PREPARE SUPER-HYDROPHOBIC SURFACES ON SILICONE RUBBERS
Materiali in tehnologije / Materials and technology 51 (2017) 5, 783–787 787
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
H. Y. JIN et al.: INVESTIGATION OF THE STATIC ICING PROPERTY FOR SUPER-HYDROPHOBIC COATINGS ...
789–793
INVESTIGATION OF THE STATIC ICING PROPERTY FOR
SUPER-HYDROPHOBIC COATINGS ON ALUMINIUM
PREISKAVA LASTNOSTI STATI^NE ZALEDENITVE PRI
SUPERHIDROFOBNIH PREVLEKAH NA ALUMINIJU
Hai Yun Jin1, Shi Chao Nie1, Yu Feng Li1, Tian Feng Xu1, Peng Zhang1, Wen Li2
1Xi’an Jiaotong University Xi’an, State Key Laboratory of Electrical Insulation and Power Equipment, 710049, China
2Hubei Polytechnic University, Hubei Key Laboratory of Mine Environmental Pollution Control & Remediation and School of Chemical and
Materials Engineering, Huangshi, 435003, China
lxfzzl@126.com
Prejem rokopisa – received: 2016-09-02; sprejem za objavo – accepted for publication: 2016-10-24
doi:10.17222/mit.2016.266
An anti-icing aluminium with a super-hydrophobic surface was fabricated. A static icing system was developed to investigate
the process of static icing for samples with different contact angles (CAs), and the contact state was also researched. The results
showed that the water droplets’ solidification time on a super-hydrophobic surface was much longer than for other surfaces at
the same temperature. A water droplet on the super-hydrophobic surface had a much smaller contact area than on another
surfaces. In addition, the air gap between the droplet and the super-hydrophobic surface retarded and prevented any exchange of
heat. The temperature affected the transition from the Cassie state to the Wenzel state for super-hydrophobic aluminium.
Keywords: aluminium coating, static icing, solidification time, super hydrophobicity
Narejeno je bilo prepre~evanje zaledenitve aluminija s superhidrofobno plastjo. Sistem stati~ne zaledenitve je bil razvit za
preiskavo procesa stati~ne zaledenitve za vzorce z razli~nimi kontaktnimi koti (angl. CA), preiskano pa je bilo tudi stanje
kontakta. Rezultati so pokazali, da je ~as strjevanja vodnih kapljic na superhidrofobni povr{ini precej dalj{i kot na drugih
povr{inah pri enaki temperaturi. Vodna kapljica je imela na superhidrofobni povr{ini precej manj{o kontaktno povr{ino kot na
drugih povr{inah. Poleg tega je zra~na re`a med kapljicami in superhidrofobno povr{ino zaustavljala in prepre~evala vsako
izmenjavo toplote. Temperatura je vplivala na prehod iz stanja Cassie v stanje Wenzel za superhidrofobni aluminij.
Klju~ne besede: aluminijeva prevleka, stati~na zaledenitev, ~as strjevanja, superhidrofobnost
1 INTRODUCTION
Icing is a natural phenomenon that occurs on the
surfaces of objects in the bad conditions of low tem-
peratures and freezing rain. An ice disaster is a serious
natural disaster for a power system,1,2 especially for
abnormal weather with the El Nino and La Nina pheno-
mena, the ice disaster became more frequently. In the
past, for a long time, de-icing methods were active, and
external energy was required to deal with the icing on the
surfaces.3,4 In recent years, anti-icing coatings that could
improve the icing performance effectively received more
attention, such as electric heating coatings, photo-ther-
mal coatings and hydrophobic coatings. Two features are
required for promising anti-icing materials. One is that
overcooled water droplets can roll off from the surface
easily before crystallization. And the other, is that the ice
adhesion will be weaker when the ice is accumulated on
the surface.5 Super-hydrophobic coatings are promising
materials due to the retardation and prevention of icing
and have become a focus of study.6–16
The researches on super-hydrophobic materials for
anti-icing are abundant; however, there are few reports
about the process in detail of water condensing and then
freezing to form ice on a cold super-hydrophobic
surface.14,15 M. He et al.14 fabricated super-hydrophobic
surfaces with isotactic polypropylene; the results showed
that the micro- and nanometre structures led to a delay of
the solidification of liquid water at the three phase line
region. M. He et al.14 fabricated super-hydrophobic
surfaces with ZnO nanorod arrays; the process of icing at
different temperatures below the freezing point was
studied.15 However, the effect on the mechanism is un-
clear and needs more systemic and deep research.
Furthermore, there are few reports about super-hydro-
phobic aluminium used in a power transmission line.
Aluminium has good thermal conductivity, so different
phenomenon may occur in the process of icing. S. L.
Zheng et al.17,18 fabricated super-hydrophobic aluminium
coatings using the Electrochemical Corrosion Method.
The adhesion strength of ice on coatings was researched,
but the process of solidification of the droplets was not
mentioned.
In this paper, super-hydrophobic coatings were fabri-
cated using the Chemical Etching Method and a static
icing system was developed to study the ice crystalliza-
tion process of super-cooled water droplets on different
coatings and different temperatures below the freezing
point. The results are helpful to provide theoretical basis
effectively for the design and fabrication of super-hydro-
phobic transmission aluminium conductors in the future.
Materiali in tehnologije / Materials and technology 51 (2017) 5, 789–793 789
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.1:67.017:669.058 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)789(2017)
H. Y. JIN et al.: INVESTIGATION OF THE STATIC ICING PROPERTY FOR SUPER-HYDROPHOBIC COATINGS ...
789–793
INVESTIGATION OF THE STATIC ICING PROPERTY FOR
SUPER-HYDROPHOBIC COATINGS ON ALUMINIUM
PREISKAVA LASTNOSTI STATI^NE ZALEDENITVE PRI
SUPERHIDROFOBNIH PREVLEKAH NA ALUMINIJU
Hai Yun Jin1, Shi Chao Nie1, Yu Feng Li1, Tian Feng Xu1, Peng Zhang1, Wen Li2
1Xi’an Jiaotong University Xi’an, State Key Laboratory of Electrical Insulation and Power Equipment, 710049, China
2Hubei Polytechnic University, Hubei Key Laboratory of Mine Environmental Pollution Control & Remediation and School of Chemical and
Materials Engineering, Huangshi, 435003, China
lxfzzl@126.com
Prejem rokopisa – received: 2016-09-02; sprejem za objavo – accepted for publication: 2016-10-24
doi:10.17222/mit.2016.266
An anti-icing aluminium with a super-hydrophobic surface was fabricated. A static icing system was developed to investigate
the process of static icing for samples with different contact angles (CAs), and the contact state was also researched. The results
showed that the water droplets’ solidification time on a super-hydrophobic surface was much longer than for other surfaces at
the same temperature. A water droplet on the super-hydrophobic surface had a much smaller contact area than on another
surfaces. In addition, the air gap between the droplet and the super-hydrophobic surface retarded and prevented any exchange of
heat. The temperature affected the transition from the Cassie state to the Wenzel state for super-hydrophobic aluminium.
Keywords: aluminium coating, static icing, solidification time, super hydrophobicity
Narejeno je bilo prepre~evanje zaledenitve aluminija s superhidrofobno plastjo. Sistem stati~ne zaledenitve je bil razvit za
preiskavo procesa stati~ne zaledenitve za vzorce z razli~nimi kontaktnimi koti (angl. CA), preiskano pa je bilo tudi stanje
kontakta. Rezultati so pokazali, da je ~as strjevanja vodnih kapljic na superhidrofobni povr{ini precej dalj{i kot na drugih
povr{inah pri enaki temperaturi. Vodna kapljica je imela na superhidrofobni povr{ini precej manj{o kontaktno povr{ino kot na
drugih povr{inah. Poleg tega je zra~na re`a med kapljicami in superhidrofobno povr{ino zaustavljala in prepre~evala vsako
izmenjavo toplote. Temperatura je vplivala na prehod iz stanja Cassie v stanje Wenzel za superhidrofobni aluminij.
Klju~ne besede: aluminijeva prevleka, stati~na zaledenitev, ~as strjevanja, superhidrofobnost
1 INTRODUCTION
Icing is a natural phenomenon that occurs on the
surfaces of objects in the bad conditions of low tem-
peratures and freezing rain. An ice disaster is a serious
natural disaster for a power system,1,2 especially for
abnormal weather with the El Nino and La Nina pheno-
mena, the ice disaster became more frequently. In the
past, for a long time, de-icing methods were active, and
external energy was required to deal with the icing on the
surfaces.3,4 In recent years, anti-icing coatings that could
improve the icing performance effectively received more
attention, such as electric heating coatings, photo-ther-
mal coatings and hydrophobic coatings. Two features are
required for promising anti-icing materials. One is that
overcooled water droplets can roll off from the surface
easily before crystallization. And the other, is that the ice
adhesion will be weaker when the ice is accumulated on
the surface.5 Super-hydrophobic coatings are promising
materials due to the retardation and prevention of icing
and have become a focus of study.6–16
The researches on super-hydrophobic materials for
anti-icing are abundant; however, there are few reports
about the process in detail of water condensing and then
freezing to form ice on a cold super-hydrophobic
surface.14,15 M. He et al.14 fabricated super-hydrophobic
surfaces with isotactic polypropylene; the results showed
that the micro- and nanometre structures led to a delay of
the solidification of liquid water at the three phase line
region. M. He et al.14 fabricated super-hydrophobic
surfaces with ZnO nanorod arrays; the process of icing at
different temperatures below the freezing point was
studied.15 However, the effect on the mechanism is un-
clear and needs more systemic and deep research.
Furthermore, there are few reports about super-hydro-
phobic aluminium used in a power transmission line.
Aluminium has good thermal conductivity, so different
phenomenon may occur in the process of icing. S. L.
Zheng et al.17,18 fabricated super-hydrophobic aluminium
coatings using the Electrochemical Corrosion Method.
The adhesion strength of ice on coatings was researched,
but the process of solidification of the droplets was not
mentioned.
In this paper, super-hydrophobic coatings were fabri-
cated using the Chemical Etching Method and a static
icing system was developed to study the ice crystalliza-
tion process of super-cooled water droplets on different
coatings and different temperatures below the freezing
point. The results are helpful to provide theoretical basis
effectively for the design and fabrication of super-hydro-
phobic transmission aluminium conductors in the future.
Materiali in tehnologije / Materials and technology 51 (2017) 5, 789–793 789
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.1:67.017:669.058 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)789(2017)
2 EXPERIMENTAL PART
In order to fabricate samples with different hydro-
phobicities, a pre-treatment was necessary for the
electrically pure aluminium. First, 1000 mesh sandpaper
was used to polish the surface of the aluminium for the
purpose of removing dense oxide layer, then using water
and absolute ethyl alcohol to clean filth and grease on
the surface by ultrasonic cleaning for 10 min. After
pre-treatment, the aluminium can be used as a hydrophi-
lic aluminium surface samples (A1). Using stearic acid
to modify the aluminium surface, stearic acid is a kind of
long-chain fatty acid and has a lower free energy. As
shown in Figure 1, a dehydration reaction would happen
between hydroxyl groups on the surface of the
aluminium and the carboxyl group in stearic acid. Then a
tight and thin coating with low free energy would be
generated on the surface of the aluminium and hydro-
phobic aluminium surface samples (A2) could be
obtained. As for super-hydrophobic aluminium surface
samples, put the aluminium after pre-treatment to the
20 % of mass fractions of hydrochloric acid to etch for 1
min, after drying put it in the 1 % of mass fractions of
stearic acid to modify the surface energy for 15 min,
finally put the sample into the 90 °C oven to fabricate
super-hydrophobic aluminium surface samples (A3).19
Figure 2 shows that static icing system was made up
of a Peltier chip, the cooling fan, temperature sensors
and a temperature controller. The Peltier chip
(TEC-1206) was supplied with power by a 12V DC
powervsupply, the temperature difference between cold
end and hot end was 75 °C and the temperature on the
cold end could be –20 °C. The microprocessor in the
semiconductor refrigeration temperature controller
(TTC-B) was used to control the DC power in order to
realize the temperature control. The test of static icing
characteristics was completed at room temperature. In
order to compare the process of icing, three kinds of
typical aluminium wire samples (A1, A2 and A3) were
observed. Detailed test steps was as follows:
1) Setting the surrounding environment in a dark-
room, the entire Peltier cooling system was placed in a
simple studio, the illumination was backlit bright LED
projection mode, and the Peltier chip temperature was
kept constant by adjusting the parameters of the PID
control system (–5 °C± 0.5 °C, –3 °C ± 0.5 °C and –1 °C
± 0.5 °C), and recorded both the ambient temperature
and the humidity value;
2) Turning on the power of the Peltier cooling sys-
tem. When the temperature of the surface was constant,
we placed the samples on the Peltier surface, 15 μL
water droplets were injected onto the samples’ surface
rapidly by using a micro-syringe, and we recorded the
solidification process of droplets using a digital camera.
3 RESULTS AND DISCUSSION
Figure 3 shows the CA of droplets on 3 typical
aluminium samples at room temperature. The CA on A1,
A2 and A3 were 82°, 109° and 152°, respectively. Heat
transfer happened between the droplets and the cold
surface. Droplets placed on the surface of sample were
liquid at first, and because of the results of differences of
hydrophobicity, different contact states on three different
surfaces were observed. As shown in Figure 3, the
droplet on the surface of A1 was like a pan upside down,
it had largest contact area; the droplet on the surface of
A2 was like a semicircle, it had larger contact area and
droplet on the surface of A3 was like a ball, it had the
smallest contact area.
Figure 4 shows SEM images and 3D laser scanning
of A3, it can be seen that A3 had many micron structure
pits and protrusions, and the size of these pits and pro-
trusions was about 3 μm. From Figure 4b and 4d, a
rough structure similar with building blocks could be
observed, so the pits and protrusions connected with
each other. Besides, the micron pits and protrusions had
submicron even nano-sized rough structure on their
surfaces. This surface structure was easy to form super
hydrophobicity with a water droplet.20
Figure 5 showed droplets’ solidification process on
samples (surface temperature is –3 °C). Droplets on the
surface of samples were liquid at first, as the time pro-
H. Y. JIN et al.: INVESTIGATION OF THE STATIC ICING PROPERTY FOR SUPER-HYDROPHOBIC COATINGS ...
790 Materiali in tehnologije / Materials and technology 51 (2017) 5, 789–793
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Contact angle of droplets on 3 typical aluminium samples at
room temperature: a) A1, b) A2 and c) A3
Figure 1: Modification mechanism of stearic acid on the aluminium
surface
Figure 2: Schematic diagram of static icing test system
longed, gas/liquid/solid three phase contact line between
droplet and surface of sample expanded, and the contact
area between the droplet and the surface increased,
which meant the CA decreasing as the time prolonged. A
phase transition occurred in the droplets because of the
heat transfer between the cold surfaces and the droplets,
which meant the transition from liquid to solid for the
droplets. Also shown was the process of solidification
after the phase transition, a clear liquid/solid line could
be observed and this line grew up from bottom to the top.
Under the liquid/solid line, the droplets were already
frozen ice. Above the line, the droplets were still liquid.
The liquid/solid line grew up at different speeds on three
typical aluminium surfaces; the line grew up quickly on
the A1 and A2 but relatively slowly on the A3. The line
went up to the top of the droplet, and finally, a peak on
the top of droplet was generated, which meant the end of
solidification process.
Figure 5 showed the phase transition time on diffe-
rent surfaces. The phase transition starting times for A1,
A2 and A3 were 20 s, 34 s and 348 s, respectively. The
phase-transition time on A3 was much longer than the
other two surfaces. The solidification time for total drop-
let for A1, A2 and A3 were 38 s, 49 s and 404 s,
respectively, the solidification time increased with
increasing CA. The solidification time of the droplet for
A3 was 10.6 times that of for A1 and 8.2 times that of
for A2. Therefore, the solidification process of water
H. Y. JIN et al.: INVESTIGATION OF THE STATIC ICING PROPERTY FOR SUPER-HYDROPHOBIC COATINGS ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 789–793 791
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Microstructure of A3: a) The SEM image magnified 1000 times, b) the SEM image magnified 5000 times, c) the 3D laser scanning
magnified 10 times and d) the 3D laser scanning magnified 20 times
Figure 5: Droplets’ solidification process on the surfaces of 3 typical
aluminum wire samples, the red lines are the boundary of liquid and
solid: a) A1, b) A2 and c) A3
2 EXPERIMENTAL PART
In order to fabricate samples with different hydro-
phobicities, a pre-treatment was necessary for the
electrically pure aluminium. First, 1000 mesh sandpaper
was used to polish the surface of the aluminium for the
purpose of removing dense oxide layer, then using water
and absolute ethyl alcohol to clean filth and grease on
the surface by ultrasonic cleaning for 10 min. After
pre-treatment, the aluminium can be used as a hydrophi-
lic aluminium surface samples (A1). Using stearic acid
to modify the aluminium surface, stearic acid is a kind of
long-chain fatty acid and has a lower free energy. As
shown in Figure 1, a dehydration reaction would happen
between hydroxyl groups on the surface of the
aluminium and the carboxyl group in stearic acid. Then a
tight and thin coating with low free energy would be
generated on the surface of the aluminium and hydro-
phobic aluminium surface samples (A2) could be
obtained. As for super-hydrophobic aluminium surface
samples, put the aluminium after pre-treatment to the
20 % of mass fractions of hydrochloric acid to etch for 1
min, after drying put it in the 1 % of mass fractions of
stearic acid to modify the surface energy for 15 min,
finally put the sample into the 90 °C oven to fabricate
super-hydrophobic aluminium surface samples (A3).19
Figure 2 shows that static icing system was made up
of a Peltier chip, the cooling fan, temperature sensors
and a temperature controller. The Peltier chip
(TEC-1206) was supplied with power by a 12V DC
powervsupply, the temperature difference between cold
end and hot end was 75 °C and the temperature on the
cold end could be –20 °C. The microprocessor in the
semiconductor refrigeration temperature controller
(TTC-B) was used to control the DC power in order to
realize the temperature control. The test of static icing
characteristics was completed at room temperature. In
order to compare the process of icing, three kinds of
typical aluminium wire samples (A1, A2 and A3) were
observed. Detailed test steps was as follows:
1) Setting the surrounding environment in a dark-
room, the entire Peltier cooling system was placed in a
simple studio, the illumination was backlit bright LED
projection mode, and the Peltier chip temperature was
kept constant by adjusting the parameters of the PID
control system (–5 °C± 0.5 °C, –3 °C ± 0.5 °C and –1 °C
± 0.5 °C), and recorded both the ambient temperature
and the humidity value;
2) Turning on the power of the Peltier cooling sys-
tem. When the temperature of the surface was constant,
we placed the samples on the Peltier surface, 15 μL
water droplets were injected onto the samples’ surface
rapidly by using a micro-syringe, and we recorded the
solidification process of droplets using a digital camera.
3 RESULTS AND DISCUSSION
Figure 3 shows the CA of droplets on 3 typical
aluminium samples at room temperature. The CA on A1,
A2 and A3 were 82°, 109° and 152°, respectively. Heat
transfer happened between the droplets and the cold
surface. Droplets placed on the surface of sample were
liquid at first, and because of the results of differences of
hydrophobicity, different contact states on three different
surfaces were observed. As shown in Figure 3, the
droplet on the surface of A1 was like a pan upside down,
it had largest contact area; the droplet on the surface of
A2 was like a semicircle, it had larger contact area and
droplet on the surface of A3 was like a ball, it had the
smallest contact area.
Figure 4 shows SEM images and 3D laser scanning
of A3, it can be seen that A3 had many micron structure
pits and protrusions, and the size of these pits and pro-
trusions was about 3 μm. From Figure 4b and 4d, a
rough structure similar with building blocks could be
observed, so the pits and protrusions connected with
each other. Besides, the micron pits and protrusions had
submicron even nano-sized rough structure on their
surfaces. This surface structure was easy to form super
hydrophobicity with a water droplet.20
Figure 5 showed droplets’ solidification process on
samples (surface temperature is –3 °C). Droplets on the
surface of samples were liquid at first, as the time pro-
H. Y. JIN et al.: INVESTIGATION OF THE STATIC ICING PROPERTY FOR SUPER-HYDROPHOBIC COATINGS ...
790 Materiali in tehnologije / Materials and technology 51 (2017) 5, 789–793
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Contact angle of droplets on 3 typical aluminium samples at
room temperature: a) A1, b) A2 and c) A3
Figure 1: Modification mechanism of stearic acid on the aluminium
surface
Figure 2: Schematic diagram of static icing test system
longed, gas/liquid/solid three phase contact line between
droplet and surface of sample expanded, and the contact
area between the droplet and the surface increased,
which meant the CA decreasing as the time prolonged. A
phase transition occurred in the droplets because of the
heat transfer between the cold surfaces and the droplets,
which meant the transition from liquid to solid for the
droplets. Also shown was the process of solidification
after the phase transition, a clear liquid/solid line could
be observed and this line grew up from bottom to the top.
Under the liquid/solid line, the droplets were already
frozen ice. Above the line, the droplets were still liquid.
The liquid/solid line grew up at different speeds on three
typical aluminium surfaces; the line grew up quickly on
the A1 and A2 but relatively slowly on the A3. The line
went up to the top of the droplet, and finally, a peak on
the top of droplet was generated, which meant the end of
solidification process.
Figure 5 showed the phase transition time on diffe-
rent surfaces. The phase transition starting times for A1,
A2 and A3 were 20 s, 34 s and 348 s, respectively. The
phase-transition time on A3 was much longer than the
other two surfaces. The solidification time for total drop-
let for A1, A2 and A3 were 38 s, 49 s and 404 s,
respectively, the solidification time increased with
increasing CA. The solidification time of the droplet for
A3 was 10.6 times that of for A1 and 8.2 times that of
for A2. Therefore, the solidification process of water
H. Y. JIN et al.: INVESTIGATION OF THE STATIC ICING PROPERTY FOR SUPER-HYDROPHOBIC COATINGS ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 789–793 791
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Microstructure of A3: a) The SEM image magnified 1000 times, b) the SEM image magnified 5000 times, c) the 3D laser scanning
magnified 10 times and d) the 3D laser scanning magnified 20 times
Figure 5: Droplets’ solidification process on the surfaces of 3 typical
aluminum wire samples, the red lines are the boundary of liquid and
solid: a) A1, b) A2 and c) A3
droplet could be greatly delayed by the super-hydro-
phobic surface (A3).
The results show that the CA directly affects the
freezing process of the droplets. For A1 and A2, the con-
tact states between droplets and surfaces were the
Wenzel state. The contact angle on A2 was larger than
on A1, so the contact area between the droplet and A2
was smaller. For A3, the contact state of ball-like droplet
made the smallest contact area. It was known that the
heat-transfer rate was related to the contact area21, solidi-
fication time on A3 was the longest. At the same time,
the vertical height of the droplet on A3 was greater,
which meant more distance and time should be required
for the solidification process from the bottom to the top
of the droplet, therefore, the solidification time on A3
was the longest. On the other hand, for A3, the contact
state was the Cassie state. There were a lot of air films in
the contact interface and the air films played a role of
energy barrier, which greatly delayed the heat exchange
process between the droplets and surfaces, so the soli-
dification time of droplets for A3 surface was also much
longer than the other two surfaces.
After solidification, the top of the droplet would form
a peak, and there were two main reasons for this: one
was that the volume of droplets increased from liquid to
solid; the other was that the droplet contained some of
air, the dissolved air would form bubbles and released
out of droplet during the process of phase transition and
solidification. However, the peak of the droplet on the
super-hydrophobic surface was more acute, which
indicated that there was other air in the droplet besides
the dissolved air, the explanation for this phenomenon
could be the involvement of an air film between the
droplet and the surface of the super-hydrophobic alumi-
nium sample. Results might indirectly prove that the
contact state of the droplet and the super-hydrophobic
aluminium surface was in the Cassie contact state.
Figure 6 shows the relationship between the solidifi-
cation time and surface temperature (the data of solidifi-
cation time was average data of five experiments). When
the surface temperature was –1 °C, the solidification
times for A1, A2 and A3 were 41.7 s, 50.7 s and 803.7 s,
respectively. When the surface temperature of the
samples was –3 °C, the solidification time for A1, A2
and A3 were 36 s, 47.7 s and 390 s respectively. When
the surface temperature was –5 °C, the data were 31.7 s,
35 s and 186 s, respectively. From Figure 6 it can be
seen that with the reduction of surface temperature, the
solidification time of the droplets on three typical sam-
ples should decrease. This was because of the increase of
the energy exchange speed and the decrease of
nucleation barrier of surfaces with decreasing surface
temperature. At different surface temperatures the soli-
dification time of the droplets on A3 was much longer
than A1 and A2, which showed the excellent static
anti-icing performance of A3. When the surface tem-
peratures were –1 °C, –3 °C and –5 °C, the droplets’
solidification times on A3 were 803.7 s, 390.0 s and
186.0 s, respectively. With a decreasing surface tempe-
rature, the solidification time of the droplets on A3 fell
more rapidly than the other two typical surfaces. Due to
the instability of the air film between the droplets and
surface,22 the contact state change from the Cassie state
to the Wenzel state could become easier for A3, with tem-
perature decreasing (Figure 5c) for 0 s and 5 min 17 s.
4 CONCLUSIONS
In summary, the super-hydrophobic coatings on alu-
minium (CA=152°) were fabricated by Chemical
Etching and Modification Method. For the super-hydro-
phobic surface, the solidification and phase-change time
of the droplets were much longer than the hydrophilic
aluminium surface and the hydrophobic aluminium
surface. This was because for the super-hydrophobic
aluminium surfaces the freezing process could be
retarded and prevented by both smaller contact area and
air film between droplet and surface (Cassie contact
state), which influenced the thermal exchange between
the droplet and the surface.
Due to that the air film latched by the droplet on the
super-hydrophobic surface was not stable, with the
temperature decreasing, the droplet transited from the
Cassie contact state to the Wenzel contact state more
easily.
Acknowledgements
This research was financially supported by the
National Natural Science Foundation of China (No.
51272208), the Natural Science Foundation of Hubei
Province (2010CDA026), the Key Program of Hubei
Provincial Department of Education (Z20104401), the
Outstanding Scientific and Technological Innovation
H. Y. JIN et al.: INVESTIGATION OF THE STATIC ICING PROPERTY FOR SUPER-HYDROPHOBIC COATINGS ...
792 Materiali in tehnologije / Materials and technology 51 (2017) 5, 789–793
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Relationship between solidification time and surface tempe-
rature: the picture shows the initial liquid state and the final solid state
of the water droplets on A3 when the surface temperature is –3 °C
Team Project of Universities in Hubei Province
(T201423), and the Talent Program of Hubei Polytechnic
University (11yjz04R).
5 REFERENCES
1 Y. Hu, Analysis and Countermeasures Discussion for Large Area
Icing Accident on Power Grid, High Voltage Engineering, 34 (2008)
2, 215–219, doi:10.13336/j.1003-6520.hve.2008.02.020
2 Y. Yang, J. Li, J. Hu, Y. Zhao, X. Jiang, C. Sun, Experimental Study
on Icing Properties of Superhydrophobic Coatings on Insulators,
High Voltage Engineering, 36 (2010) 3, 621–626, doi:10.13336/
j.1003-6520.hve.2010.03.013
3 J. Laforte, M. Allaire, J. Laflamme, State-of-the-art on power line
de-icing, Atmospheric Research, 46 (1998) 1, 143–158, doi:10.1016/
S0169-8095(97)00057-4
4 O. Parent, A. Ilinca, Anti-icing and de-icing techniques for wind
turbines: Critical review, Cold Regions Science and Technology, 65
(2011) 1, 88–96, doi:10.1016/j.coldregions.2010.01.005
5 Y. Wang, J. Xue, Q. Wang, Q. Chen, J. Ding, Verification of Ice-
phobic/Anti-icing Properties of a Superhydrophobic Surface, ACS
applied materials & interfaces, 5 (2013) 8, 3370–3381,
doi:10.1021/am400429q
6 G. Fang, A. Amirfazli, Understanding the anti-icing behavior of
superhydrophobic surfaces, Surface Innovations, 2 (2014) 2, 94–102,
doi:10.1680/si.13.00046
7 S. Farhadi, M. Farzaneh, S. Kulinich. Anti-icing performance of
superhydrophobic surfaces, Applied Surface Science, 257 (2011) 14,
6264–6269, doi:10.1016/j.apsusc.2011.02.057
8 T. Kako, A. Nakajima, H. Irie, Z. Kato, K. Uematsu, T. Watanabe, K.
Hashimoto, Adhesion and sliding of wet snow on a super-hydro-
phobic surface with hydrophilic channels, Journal of Materials
Science, 39 (2004) 2, 547–555, doi:10.1023/B:JMSC.0000011510.
92644.3f
9 Y. Tang, Q. Zhang, X. Zhan, F. Chen, Superhydrophobic and
anti-icing properties under overcooled temperature of fluorinated
hybrid surface prepared via sol–gel process, Soft Matter, 11 (2015)
22, 4540–4550, doi:10.1039/c5sm00674k
10 P. Tourkine, M. Le Merrer, D. Quéré, Delayed freezing on water
repellent materials, Langmuir, 25 (2009) 13, 7214–7216,
doi:10.1021/la900929u
11 T. Bharathidasan, S.V. Kumar, M. Bobji, R. Chakradhar, B. J. Basu,
Effect of wettability and surface roughness on ice-adhesion strength
of hydrophilic,hydrophobic and superhydrophobic surfaces, Applied
Surface Science, 314 (2014), 241–250, doi:10.1016/j.apsusc.2014.
06.101
12 S. Kulinich, M. Farzaneh, How wetting hysteresis influences ice
adhesion strength on superhydrophobic surfaces, Langmuir, 25
(2009) 16, 8854–8856, doi:10.1021/la901439c
13 S. Kulinich, M. Farzaneh, Ice adhesion on super-hydrophobic
surfaces, Applied Surface Science, 255 (2009) 18, 8153–8157,
doi:10.1016/j.apsusc.2009.05.033
14 M. He, J. X. Wang, H. L. Li, X. L. Jin, J. J. Wang, B. Q. Liu, Y. L.
Song, Super-hydrophobic film retards frost formation, Soft Matter, 6
(2010) 11, 2396–2399, doi:10.1039/c0sm00024h
15 M. He, J. Wang, H. Li, Y. Song, Super-hydrophobic surfaces to
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retard ice/frost formation, Soft Matter, 7 (2011) 8, 3993–4000,
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16 J. Durret, N. Frolet, C. Gourgon. Hydrophobicity and anti-icing
performances of nanoimprinted and roughened fluoropolymers films
under overcooled temperature, Microelectronic Engineering, 155
(2016), 1–6, doi:10.1016/j.mee.2016.01.011
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H. Y. JIN et al.: INVESTIGATION OF THE STATIC ICING PROPERTY FOR SUPER-HYDROPHOBIC COATINGS ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 789–793 793
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
droplet could be greatly delayed by the super-hydro-
phobic surface (A3).
The results show that the CA directly affects the
freezing process of the droplets. For A1 and A2, the con-
tact states between droplets and surfaces were the
Wenzel state. The contact angle on A2 was larger than
on A1, so the contact area between the droplet and A2
was smaller. For A3, the contact state of ball-like droplet
made the smallest contact area. It was known that the
heat-transfer rate was related to the contact area21, solidi-
fication time on A3 was the longest. At the same time,
the vertical height of the droplet on A3 was greater,
which meant more distance and time should be required
for the solidification process from the bottom to the top
of the droplet, therefore, the solidification time on A3
was the longest. On the other hand, for A3, the contact
state was the Cassie state. There were a lot of air films in
the contact interface and the air films played a role of
energy barrier, which greatly delayed the heat exchange
process between the droplets and surfaces, so the soli-
dification time of droplets for A3 surface was also much
longer than the other two surfaces.
After solidification, the top of the droplet would form
a peak, and there were two main reasons for this: one
was that the volume of droplets increased from liquid to
solid; the other was that the droplet contained some of
air, the dissolved air would form bubbles and released
out of droplet during the process of phase transition and
solidification. However, the peak of the droplet on the
super-hydrophobic surface was more acute, which
indicated that there was other air in the droplet besides
the dissolved air, the explanation for this phenomenon
could be the involvement of an air film between the
droplet and the surface of the super-hydrophobic alumi-
nium sample. Results might indirectly prove that the
contact state of the droplet and the super-hydrophobic
aluminium surface was in the Cassie contact state.
Figure 6 shows the relationship between the solidifi-
cation time and surface temperature (the data of solidifi-
cation time was average data of five experiments). When
the surface temperature was –1 °C, the solidification
times for A1, A2 and A3 were 41.7 s, 50.7 s and 803.7 s,
respectively. When the surface temperature of the
samples was –3 °C, the solidification time for A1, A2
and A3 were 36 s, 47.7 s and 390 s respectively. When
the surface temperature was –5 °C, the data were 31.7 s,
35 s and 186 s, respectively. From Figure 6 it can be
seen that with the reduction of surface temperature, the
solidification time of the droplets on three typical sam-
ples should decrease. This was because of the increase of
the energy exchange speed and the decrease of
nucleation barrier of surfaces with decreasing surface
temperature. At different surface temperatures the soli-
dification time of the droplets on A3 was much longer
than A1 and A2, which showed the excellent static
anti-icing performance of A3. When the surface tem-
peratures were –1 °C, –3 °C and –5 °C, the droplets’
solidification times on A3 were 803.7 s, 390.0 s and
186.0 s, respectively. With a decreasing surface tempe-
rature, the solidification time of the droplets on A3 fell
more rapidly than the other two typical surfaces. Due to
the instability of the air film between the droplets and
surface,22 the contact state change from the Cassie state
to the Wenzel state could become easier for A3, with tem-
perature decreasing (Figure 5c) for 0 s and 5 min 17 s.
4 CONCLUSIONS
In summary, the super-hydrophobic coatings on alu-
minium (CA=152°) were fabricated by Chemical
Etching and Modification Method. For the super-hydro-
phobic surface, the solidification and phase-change time
of the droplets were much longer than the hydrophilic
aluminium surface and the hydrophobic aluminium
surface. This was because for the super-hydrophobic
aluminium surfaces the freezing process could be
retarded and prevented by both smaller contact area and
air film between droplet and surface (Cassie contact
state), which influenced the thermal exchange between
the droplet and the surface.
Due to that the air film latched by the droplet on the
super-hydrophobic surface was not stable, with the
temperature decreasing, the droplet transited from the
Cassie contact state to the Wenzel contact state more
easily.
Acknowledgements
This research was financially supported by the
National Natural Science Foundation of China (No.
51272208), the Natural Science Foundation of Hubei
Province (2010CDA026), the Key Program of Hubei
Provincial Department of Education (Z20104401), the
Outstanding Scientific and Technological Innovation
H. Y. JIN et al.: INVESTIGATION OF THE STATIC ICING PROPERTY FOR SUPER-HYDROPHOBIC COATINGS ...
792 Materiali in tehnologije / Materials and technology 51 (2017) 5, 789–793
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Relationship between solidification time and surface tempe-
rature: the picture shows the initial liquid state and the final solid state
of the water droplets on A3 when the surface temperature is –3 °C
Team Project of Universities in Hubei Province
(T201423), and the Talent Program of Hubei Polytechnic
University (11yjz04R).
5 REFERENCES
1 Y. Hu, Analysis and Countermeasures Discussion for Large Area
Icing Accident on Power Grid, High Voltage Engineering, 34 (2008)
2, 215–219, doi:10.13336/j.1003-6520.hve.2008.02.020
2 Y. Yang, J. Li, J. Hu, Y. Zhao, X. Jiang, C. Sun, Experimental Study
on Icing Properties of Superhydrophobic Coatings on Insulators,
High Voltage Engineering, 36 (2010) 3, 621–626, doi:10.13336/
j.1003-6520.hve.2010.03.013
3 J. Laforte, M. Allaire, J. Laflamme, State-of-the-art on power line
de-icing, Atmospheric Research, 46 (1998) 1, 143–158, doi:10.1016/
S0169-8095(97)00057-4
4 O. Parent, A. Ilinca, Anti-icing and de-icing techniques for wind
turbines: Critical review, Cold Regions Science and Technology, 65
(2011) 1, 88–96, doi:10.1016/j.coldregions.2010.01.005
5 Y. Wang, J. Xue, Q. Wang, Q. Chen, J. Ding, Verification of Ice-
phobic/Anti-icing Properties of a Superhydrophobic Surface, ACS
applied materials & interfaces, 5 (2013) 8, 3370–3381,
doi:10.1021/am400429q
6 G. Fang, A. Amirfazli, Understanding the anti-icing behavior of
superhydrophobic surfaces, Surface Innovations, 2 (2014) 2, 94–102,
doi:10.1680/si.13.00046
7 S. Farhadi, M. Farzaneh, S. Kulinich. Anti-icing performance of
superhydrophobic surfaces, Applied Surface Science, 257 (2011) 14,
6264–6269, doi:10.1016/j.apsusc.2011.02.057
8 T. Kako, A. Nakajima, H. Irie, Z. Kato, K. Uematsu, T. Watanabe, K.
Hashimoto, Adhesion and sliding of wet snow on a super-hydro-
phobic surface with hydrophilic channels, Journal of Materials
Science, 39 (2004) 2, 547–555, doi:10.1023/B:JMSC.0000011510.
92644.3f
9 Y. Tang, Q. Zhang, X. Zhan, F. Chen, Superhydrophobic and
anti-icing properties under overcooled temperature of fluorinated
hybrid surface prepared via sol–gel process, Soft Matter, 11 (2015)
22, 4540–4550, doi:10.1039/c5sm00674k
10 P. Tourkine, M. Le Merrer, D. Quéré, Delayed freezing on water
repellent materials, Langmuir, 25 (2009) 13, 7214–7216,
doi:10.1021/la900929u
11 T. Bharathidasan, S.V. Kumar, M. Bobji, R. Chakradhar, B. J. Basu,
Effect of wettability and surface roughness on ice-adhesion strength
of hydrophilic,hydrophobic and superhydrophobic surfaces, Applied
Surface Science, 314 (2014), 241–250, doi:10.1016/j.apsusc.2014.
06.101
12 S. Kulinich, M. Farzaneh, How wetting hysteresis influences ice
adhesion strength on superhydrophobic surfaces, Langmuir, 25
(2009) 16, 8854–8856, doi:10.1021/la901439c
13 S. Kulinich, M. Farzaneh, Ice adhesion on super-hydrophobic
surfaces, Applied Surface Science, 255 (2009) 18, 8153–8157,
doi:10.1016/j.apsusc.2009.05.033
14 M. He, J. X. Wang, H. L. Li, X. L. Jin, J. J. Wang, B. Q. Liu, Y. L.
Song, Super-hydrophobic film retards frost formation, Soft Matter, 6
(2010) 11, 2396–2399, doi:10.1039/c0sm00024h
15 M. He, J. Wang, H. Li, Y. Song, Super-hydrophobic surfaces to
condensed micro-droplets at temperatures below the freezing point
retard ice/frost formation, Soft Matter, 7 (2011) 8, 3993–4000,
doi:10.1039/c0sm01504k
16 J. Durret, N. Frolet, C. Gourgon. Hydrophobicity and anti-icing
performances of nanoimprinted and roughened fluoropolymers films
under overcooled temperature, Microelectronic Engineering, 155
(2016), 1–6, doi:10.1016/j.mee.2016.01.011
17 S. Zheng, C. Li, Q. Fu, M. Li, W. Hu, Q. Wang, M. Du, X. Liu, Z.
Chen, Fabrication of self-cleaning superhydrophobic surface on
aluminum alloys with excellent corrosion resistance, Surface &
Coatings Technology, 276 (2015), 341–348, doi:10.1016/j.surfcoat.
2015.07.002
18 S. Zheng, C. Li, Q. Fu, W. Hu, T. Xiang, Q. Wang, M. Du, X. Liu, Z.
Chen, Development of stable superhydrophobic coatings on alumi-
num surface for corrosion-resistant, self-cleaning,and anti-icing
applications, Materials and Design, 93 (2016), 261–270,
doi:10.1016/j.matdes.2015.12.155
19 Z. Huang, Y. Li, P. Jin, M. Hu, B. He, N. Gao, H. Zhang, H. Jin,
Fabrication of Aluminum Superhydrophobic Surface with Facile
Chemical Etching Method, Materials Science Forum, 804 (2015),
103–106, doi:10.4028/www.scientific.net/MSF.804.103
20 L. Jiang, Nanostructured Materials with Super-hydrophobic Sur-
face-from Nature to Biomimesis, Chemical Industry and Engineering
Progress, 22 (2003) 12, 1258–1264, doi:10.3321/j.issn:1000-6613.
2003.12.002
21 O. R. Enriquez, A.G. Marin, K. G. Winkels, J. H. Snoeijer, Freezing
singularities in water drops, Physics of Fluids, 24 (2012) 9, 091102,
doi:10.1063/1.4747185
22 H. Y. Jin, Y. F. Li, P. Zhang, S. C. Nie, N. K. Gao, The investigation
of the wetting behavior on the red rose petal, Applied physics Letter,
108 (2016) 15, 151605, doi:10.1063/1.4947057
H. Y. JIN et al.: INVESTIGATION OF THE STATIC ICING PROPERTY FOR SUPER-HYDROPHOBIC COATINGS ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 789–793 793
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
795–803
EFFECT OF BALL MILLING ON THE PROPERTIES OF THE
POROUS Ti–26Nb ALLOY FOR BIOMEDICAL APPLICATIONS
VPLIV KROGLI^NEGA MLETJA NA LASTNOSTI POROZNE
ZLITINE Ti–26Nb ZA BIOMEDICINSKE APLIKACIJE
Grzegorz Dercz, Izabela Matu³a
University of Silesia, Institute of Materials Science, 75 Pu³ku Piechoty Street 1 A, 41-500 Chorzów, Poland
grzegorz.dercz@us.edu.pl
Prejem rokopisa – received: 2016-09-09; sprejem za objavo – accepted for publication: 2017-02-21
doi:10.17222/mit.2016.269
The current study investigated the effect of the ball-milling time on the structural characteristics and pore morphology of a
biomedical porous Ti–26Nb (a/%) alloy. An open-cell porous material was synthesized by mechanical alloying and sintering.
Commercially available elemental metal powders of Ti and Nb were used as the starting materials. Elemental metal powders
with a nominal composition of Ti–26Nb (a/%) were milled in a planetary ball mill. During the testing, the powders were milled
for two milling times: 50 h and 70 h. During the shorter milling time, the speed of 200 min–1 was applied, and during the longer
milling time, the speed of 400 min–1 was applied. The powders were cold pressed under a pressure of 750 MPa and then sintered
at 1000 °C for 24 h. The effects of the powder milling time on the microstructure and mechanical properties of the porous
structure were investigated with optic microscopy (OM), scanning electron microscopy (SEM), X-ray diffraction (XRD),
microhardness, nanoindentation and nanomachining tests. The hardness and elastic modulus were calculated from the
load-displacement data obtained with nanoindentation using a three-sided, pyramidal, diamond (Berkovich) indenter tip. The
X-ray diffraction results confirmed the presence of the  and  phases. An analysis of diffraction patterns revealed a decrease in
the lattice parameters for the 70 h milling. In summary, it should be pointed out that the material has a hierarchical structure,
obtained during the successive stages of milling. The observed grains are composed of many smaller grains, the size of which
decreased with an increase in the milling time.
Keywords: mechanical alloying, porosity, nanoindentation, Ti–26Nb alloy
V {tudiji so preiskovali vpliv razli~no dolgega ~asa krogli~nega mletja na strukturne zna~ilnosti in morfologijo por biomedi-
cinske porozne Ti-26Nb zlitine (a/%). Porozni material z odprtimi celicami je bil sintetiziran z mehanskim legiranjem in
sintranjem. Osnovni metalur{ki prahovi z nominalno vsebnostjo Ti-26Nb (a/%) so bili mleti v planetarnem krogelnem mlinu.
Med testiranjem so bili prahovi mleti pri dveh razli~nih ~asih: 50 h in 70 h. Med kraj{im ~asom mletja je bila uporabljena hitrost
mletja 200 min–1 in med dalj{im ~asom mletja je bila uporabljena hitrost 400 min–1. Prahovi so bili hladno stiskani pod tlakom
750 MPa in nato sintrani pri 1000 °C za 24 h. Preiskovali so vpliv razli~nega ~asa mletja prahu na mikrostrukturo in na
mehanske lastnosti porozne strukture z opti~nim mikroskopom, s SEM- in XRD-analizo ter dolo~evanjem mikrotrdote in
mehanske nanoobdelave. Trdoto in modul elasti~nosti so izra~unali iz meritev odvisnosti med obremenitvijo in globino vtisa
tristranske diamantne konice (po Berkovichu). Rezultati rentgenske difrakcije so potrdili prisotnost  in  faz. Skratka, poudariti
je treba, da ima material hierarhi~no strukturo, ki je posledica faze mletja. Pregledana zrna so sestavljena iz mnogo manj{ih
podzrn, katerih velikost se je zmanj{evala s podalj{evanjem ~asa mletja.
Klju~ne besede: mehansko legiranje, poroznost, nanovtiskovanje, zlitina Ti-26Nb
1 INTRODUCTION
Titanium-based alloys seem to be the preferred ma-
terials for orthopaedic applications. They are characte-
rized by high biocompatibility, great corrosion resis-
tance, light weight and good mechanical properties and
have a relatively low modulus in comparison to the other
materials used as implants.1 However, titanium-based
alloys used for implants have important disadvantages.
Firstly, popular alloying additives (Al, V) are dangerous
for the health and life of a patient.2 Secondly, the metals
used for implants have too high an elastic modulus in
comparison to the mechanical properties of the bone.
This difference between the mechanical properties in-
duces stress between the implant and the bone.3 In recent
years, due to intensive investigation, we have seen the
production and development of various Ti-based alloys,
among which there are Ti-Mo-based, Ti-Nb-based,
Ti-Zr-based and Ti-Ta-based alloys.2
This work focuses on niobium as the additive to
titanium because it has been recognized as an extremely
promising element. Nb is highly biocompatible, has a
low ionic in-vitro cytotoxicity, and most importantly –
there is no evidence of mutagenicity or carcinogenicity.4
The Ti-Nb-based alloys are attracting more researchers
to study them due to their low modulus, good biocom-
patibility and shape-memory effect. Medical applications
of SME (shape-memory effect) alloys will be especially
increased in the near future, so the development and
basic or extensive research of the Ti-based alloys are
necessary.1,2,4,5 A series of studies confirmed the presence
of the shape-memory effect and superelasticity in the
alloys with the composition of Ti-26Nb (a/%) and also in
approximate compositions of the elements.4–7 Most of
Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803 795
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 621.926:669.293:669.295:620.192.47 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)795(2017)
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
795–803
EFFECT OF BALL MILLING ON THE PROPERTIES OF THE
POROUS Ti–26Nb ALLOY FOR BIOMEDICAL APPLICATIONS
VPLIV KROGLI^NEGA MLETJA NA LASTNOSTI POROZNE
ZLITINE Ti–26Nb ZA BIOMEDICINSKE APLIKACIJE
Grzegorz Dercz, Izabela Matu³a
University of Silesia, Institute of Materials Science, 75 Pu³ku Piechoty Street 1 A, 41-500 Chorzów, Poland
grzegorz.dercz@us.edu.pl
Prejem rokopisa – received: 2016-09-09; sprejem za objavo – accepted for publication: 2017-02-21
doi:10.17222/mit.2016.269
The current study investigated the effect of the ball-milling time on the structural characteristics and pore morphology of a
biomedical porous Ti–26Nb (a/%) alloy. An open-cell porous material was synthesized by mechanical alloying and sintering.
Commercially available elemental metal powders of Ti and Nb were used as the starting materials. Elemental metal powders
with a nominal composition of Ti–26Nb (a/%) were milled in a planetary ball mill. During the testing, the powders were milled
for two milling times: 50 h and 70 h. During the shorter milling time, the speed of 200 min–1 was applied, and during the longer
milling time, the speed of 400 min–1 was applied. The powders were cold pressed under a pressure of 750 MPa and then sintered
at 1000 °C for 24 h. The effects of the powder milling time on the microstructure and mechanical properties of the porous
structure were investigated with optic microscopy (OM), scanning electron microscopy (SEM), X-ray diffraction (XRD),
microhardness, nanoindentation and nanomachining tests. The hardness and elastic modulus were calculated from the
load-displacement data obtained with nanoindentation using a three-sided, pyramidal, diamond (Berkovich) indenter tip. The
X-ray diffraction results confirmed the presence of the  and  phases. An analysis of diffraction patterns revealed a decrease in
the lattice parameters for the 70 h milling. In summary, it should be pointed out that the material has a hierarchical structure,
obtained during the successive stages of milling. The observed grains are composed of many smaller grains, the size of which
decreased with an increase in the milling time.
Keywords: mechanical alloying, porosity, nanoindentation, Ti–26Nb alloy
V {tudiji so preiskovali vpliv razli~no dolgega ~asa krogli~nega mletja na strukturne zna~ilnosti in morfologijo por biomedi-
cinske porozne Ti-26Nb zlitine (a/%). Porozni material z odprtimi celicami je bil sintetiziran z mehanskim legiranjem in
sintranjem. Osnovni metalur{ki prahovi z nominalno vsebnostjo Ti-26Nb (a/%) so bili mleti v planetarnem krogelnem mlinu.
Med testiranjem so bili prahovi mleti pri dveh razli~nih ~asih: 50 h in 70 h. Med kraj{im ~asom mletja je bila uporabljena hitrost
mletja 200 min–1 in med dalj{im ~asom mletja je bila uporabljena hitrost 400 min–1. Prahovi so bili hladno stiskani pod tlakom
750 MPa in nato sintrani pri 1000 °C za 24 h. Preiskovali so vpliv razli~nega ~asa mletja prahu na mikrostrukturo in na
mehanske lastnosti porozne strukture z opti~nim mikroskopom, s SEM- in XRD-analizo ter dolo~evanjem mikrotrdote in
mehanske nanoobdelave. Trdoto in modul elasti~nosti so izra~unali iz meritev odvisnosti med obremenitvijo in globino vtisa
tristranske diamantne konice (po Berkovichu). Rezultati rentgenske difrakcije so potrdili prisotnost  in  faz. Skratka, poudariti
je treba, da ima material hierarhi~no strukturo, ki je posledica faze mletja. Pregledana zrna so sestavljena iz mnogo manj{ih
podzrn, katerih velikost se je zmanj{evala s podalj{evanjem ~asa mletja.
Klju~ne besede: mehansko legiranje, poroznost, nanovtiskovanje, zlitina Ti-26Nb
1 INTRODUCTION
Titanium-based alloys seem to be the preferred ma-
terials for orthopaedic applications. They are characte-
rized by high biocompatibility, great corrosion resis-
tance, light weight and good mechanical properties and
have a relatively low modulus in comparison to the other
materials used as implants.1 However, titanium-based
alloys used for implants have important disadvantages.
Firstly, popular alloying additives (Al, V) are dangerous
for the health and life of a patient.2 Secondly, the metals
used for implants have too high an elastic modulus in
comparison to the mechanical properties of the bone.
This difference between the mechanical properties in-
duces stress between the implant and the bone.3 In recent
years, due to intensive investigation, we have seen the
production and development of various Ti-based alloys,
among which there are Ti-Mo-based, Ti-Nb-based,
Ti-Zr-based and Ti-Ta-based alloys.2
This work focuses on niobium as the additive to
titanium because it has been recognized as an extremely
promising element. Nb is highly biocompatible, has a
low ionic in-vitro cytotoxicity, and most importantly –
there is no evidence of mutagenicity or carcinogenicity.4
The Ti-Nb-based alloys are attracting more researchers
to study them due to their low modulus, good biocom-
patibility and shape-memory effect. Medical applications
of SME (shape-memory effect) alloys will be especially
increased in the near future, so the development and
basic or extensive research of the Ti-based alloys are
necessary.1,2,4,5 A series of studies confirmed the presence
of the shape-memory effect and superelasticity in the
alloys with the composition of Ti-26Nb (a/%) and also in
approximate compositions of the elements.4–7 Most of
Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803 795
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 621.926:669.293:669.295:620.192.47 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)795(2017)
these alloys are obtained with a conventional method
such as arc melting. One of the alternatives to the widely
used arc melting is powder metallurgy (PM). In the
scientific literature, there are evidences of successful
preparations of titanium-based materials with this pro-
duction method.8–10 This production method has a
number of significant advantages. It allows creating
alloys of materials with large differences between their
melting temperatures or densities. The most important
advantages of PM for the production of biomaterials are
its low cost and the possibility to produce highly porous
materials.4,11
The aim of the present paper is to determine the
possibility to produce a Ti–26Nb (a/%) alloy using the
powder-metallurgy method. Simultaneously, this re-
search is to develop a technology of producing Ti-based
alloys and estimate the influence of changing parameters
during the mechanical alloying. Such research is the
starting point for further development of nickel-free
alloys with the shape-memory effect based on titanium
and niobium. In this article, the influence of the milling
time on the microstructure and mechanical properties of
the produced alloy is investigated.
2 EXPERIMENTAL PART
Commercial elemental powders of Ti (Atlantic
Equipment Engineers (AEE), 99.7 %, <20 μm) and Nb
(Atlantic Equipment Engineers (AEE), 99.8 %, < 5 μm)
were used as the initial materials for the synthesis of the
alloy. Elemental metal powders with the nominal com-
position of Ti-26Nb (a/%) were milled for two different
milling times, 50 h and 70 h, in a planetary ball mill
Fritch PULVERISETTE 7 premium line. To prevent the
powder from oxidation as much as possible, the process
was carried out in an argon-protective atmosphere.
During the shorter milling time, the speed of 200 min–1
was applied, and for the longer milling time (with
additional 20 h), the speed of 400 min–1 was applied. The
powders were cold isostatically pressed under a 750 MPa
pressure without any substances that improve the
porosity (e.g., space holders). Then the material was
sintered at 1000 °C for 24 h and cooled in the furnace to
room temperature.
The crystalline structure and phase content of the
sintered materials were tested with X-ray diffraction.
The refinement of the X-ray diffraction pattern was
carried out using Rietveld’s whole X-ray profile fitting
technique with the DBWS 9807a program.12 The profile
function used to adjust the calculated diffractograms to
the observed ones was the pseudo-Voigt one.13,14 The
weight fraction of each component was determined
based on the optimized scale factors with the use of the
relation proposed by Hill and Howard.15,16 Based on the
Williamson–Hall method, the crystallite sizes (D) and
lattice strain <a/a> for the phases appearing in the
material at various milling times were determined. The
apparatus factors were eliminated using LaB6 (SRM
660a).
The morphology of the initial powders and the
microstructure of the sintered material were tested using
an optic microscope and scanning microscope JEOL
JSM 6480 with an accelerating voltage of 20 kV.
Additionally, a chemical analysis was performed using
an EDS detector manufactured by IXRF using the
standard calibration method.
An analysis of the microstructure was performed on
microscopic images with the use of ImageJ, which is the
software used for image processing and analysis. The
size and shape of the grains were determined using the
planimetric method. The study also included determining
the inhomogeneity of the size of the grains. A quantita-
tive analysis of the morphology of the material was
performed using the stereological parameters presented
below:
a) the size of the area of a section of grain a (μm2),
b) dimensionless lengthening factor (Equation (1):
f
h
wg1
= (1)
where: h – height and w – width of the smallest
rectangle described on the object,
c) dimensionless shape factor – the circularity:
f
F
Lg 2 2
4
=
π
(2)
where: F – area of the analyzed object; L – perimeter of
the analyzed object,
d) grain-size change-ability factor:


= x
a
(3)
where: x – grain-size standard deviation; a – grain
mean value,
the number of analyzed elements per area unit of the
image: NA [ ]1 2/mm .
The thermal behavior of the martensitic transforma-
tion was studied using a differential scanning calorimeter
(DSC) Mettler Toledo DSC-1. Transformation tempera-
tures were determined from the thermograms measured
at a heating rate of 10 deg/min and a thermal range of
–120 °C to 600 °C.
The surface morphology, hardness and elastic modu-
lus were tested using a Hysitron Triboindenter Ti950
with AFM QScope 250. The measurements were
calculated from the load-displacement data obtained with
the nanoindentation, using a three-sided, pyramidal,
diamond (Berkovich) indenter tip. An area of 40 μm ×
40 μm was analyzed at a scanning frequency of 0.25 Hz.
The microhardness measurement was carried out parallel
to the nanoindentation tests. The Vickers microhardness
measurement was taken at a load of 500 N and a loading
time of 10 s on a 401MVD microhardness tester.
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
796 Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
3 RESULTS AND DISCUSSION
SEM micrographs of the initial metal powders are
presented in Figure 1. The titanium-powder morphology
is irregular, having sharp corners and rough surfaces.
The dispersion of the Ti particles is wide and ranges
from a few μm to even 50 μm. The niobium powder has a
much wider dispersion of the particles in comparison to
the Ti powder.
Some of the particles have the size below 10 μm, and
the shape of these particles is irregular. Simultaneously,
there are also very large particles, the sizes of which are
greater even than 50 μm and these particles have an
irregular, polyhedral shape. As a result of the work-hard-
ening effect and high density of defects such as dislo-
cations and vacancies created during long ball-milling
times, the fracture and welding of powders reach an
equilibrium state. This state leads to the formation of
rather equiaxed and small particles.16,17 However, due to
the higher bonding strength of the finer particles, their
ability to undergo further plastic deformation is de-
creased and a higher force is required to fracture the
small particles during the ball-milling process.18 It is
worth mentioning that the reduction in the particle size at
a given milling time can also be influenced by the
milling technique.
The morphologies of the Ti-26Nb (a/%) powders at
different milling stages are shown in Figure 2. The
powder after the shortest milling time shows globular
particles. In the picture with the higher magnification
(Figure 2b), some cracks were observed, which is the
beginning of the cyclic process of material milling
during high-energy ball milling. In contrast, in the
samples after 70 h of milling, the particles became larger
and the shape of the particles became less globular.
Moreover, the particles of this sample are cold welded
and the powder particles, in particular, are composed of
the welded layers of the material. This behavior of the
material during the high-energy milling is consistent
with the conduct and following steps of mechanical
alloying.19 As shown by the other studies, the mechani-
cal-alloying method leads to a decreased particle size,
the -Ti phase being stabilized by Nb, and an increased
surface area between Ti and Nb. This promotes good
homogeneity of the solid solution.20
Figure 3 shows a summary of diffraction patterns of
the material after the mechanical synthesis, depending on
the milling time of the initial powders. The X-ray
qualitative analysis showed that only  and  phases are
present in the material after milling for 50 h and 70 h.
Change of the profile lines of individual diffraction
patterns are clearly visible.
The broadening of the particle-size distribution at
longer ball-milling times is a typical behavior of high-
energy ball-milling process.17,19,21,22 For the  and 
phases, it was observed that a decrease in the crystallite
size took place simultaneously with an increase in the
milling time (Table 1). In addition, in the case of the 
phase there was a tendency towards a decrease in the
lattice strain <a/a> (Table 1). The estimated size of the
crystallites for the powder after milling for 70 h is 9(1)
nm and 16(2) nm for the  and  phases, respectively.
The observed <a/a> lattice strain values indicate
growth in the  phase (3.43E-03 %) and in the  phase
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803 797
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: X-ray diffraction patterns of the powder after 50 h and 70 h
of the milling process
Figure 1: SEM micrographs of the initial Ti and Nb powders at
different magnifications
Figure 2: SEM micrographs of particles after different milling times:
a), b) 50 h of milling and c), d) 70 h
these alloys are obtained with a conventional method
such as arc melting. One of the alternatives to the widely
used arc melting is powder metallurgy (PM). In the
scientific literature, there are evidences of successful
preparations of titanium-based materials with this pro-
duction method.8–10 This production method has a
number of significant advantages. It allows creating
alloys of materials with large differences between their
melting temperatures or densities. The most important
advantages of PM for the production of biomaterials are
its low cost and the possibility to produce highly porous
materials.4,11
The aim of the present paper is to determine the
possibility to produce a Ti–26Nb (a/%) alloy using the
powder-metallurgy method. Simultaneously, this re-
search is to develop a technology of producing Ti-based
alloys and estimate the influence of changing parameters
during the mechanical alloying. Such research is the
starting point for further development of nickel-free
alloys with the shape-memory effect based on titanium
and niobium. In this article, the influence of the milling
time on the microstructure and mechanical properties of
the produced alloy is investigated.
2 EXPERIMENTAL PART
Commercial elemental powders of Ti (Atlantic
Equipment Engineers (AEE), 99.7 %, <20 μm) and Nb
(Atlantic Equipment Engineers (AEE), 99.8 %, < 5 μm)
were used as the initial materials for the synthesis of the
alloy. Elemental metal powders with the nominal com-
position of Ti-26Nb (a/%) were milled for two different
milling times, 50 h and 70 h, in a planetary ball mill
Fritch PULVERISETTE 7 premium line. To prevent the
powder from oxidation as much as possible, the process
was carried out in an argon-protective atmosphere.
During the shorter milling time, the speed of 200 min–1
was applied, and for the longer milling time (with
additional 20 h), the speed of 400 min–1 was applied. The
powders were cold isostatically pressed under a 750 MPa
pressure without any substances that improve the
porosity (e.g., space holders). Then the material was
sintered at 1000 °C for 24 h and cooled in the furnace to
room temperature.
The crystalline structure and phase content of the
sintered materials were tested with X-ray diffraction.
The refinement of the X-ray diffraction pattern was
carried out using Rietveld’s whole X-ray profile fitting
technique with the DBWS 9807a program.12 The profile
function used to adjust the calculated diffractograms to
the observed ones was the pseudo-Voigt one.13,14 The
weight fraction of each component was determined
based on the optimized scale factors with the use of the
relation proposed by Hill and Howard.15,16 Based on the
Williamson–Hall method, the crystallite sizes (D) and
lattice strain <a/a> for the phases appearing in the
material at various milling times were determined. The
apparatus factors were eliminated using LaB6 (SRM
660a).
The morphology of the initial powders and the
microstructure of the sintered material were tested using
an optic microscope and scanning microscope JEOL
JSM 6480 with an accelerating voltage of 20 kV.
Additionally, a chemical analysis was performed using
an EDS detector manufactured by IXRF using the
standard calibration method.
An analysis of the microstructure was performed on
microscopic images with the use of ImageJ, which is the
software used for image processing and analysis. The
size and shape of the grains were determined using the
planimetric method. The study also included determining
the inhomogeneity of the size of the grains. A quantita-
tive analysis of the morphology of the material was
performed using the stereological parameters presented
below:
a) the size of the area of a section of grain a (μm2),
b) dimensionless lengthening factor (Equation (1):
f
h
wg1
= (1)
where: h – height and w – width of the smallest
rectangle described on the object,
c) dimensionless shape factor – the circularity:
f
F
Lg 2 2
4
=
π
(2)
where: F – area of the analyzed object; L – perimeter of
the analyzed object,
d) grain-size change-ability factor:


= x
a
(3)
where: x – grain-size standard deviation; a – grain
mean value,
the number of analyzed elements per area unit of the
image: NA [ ]1 2/mm .
The thermal behavior of the martensitic transforma-
tion was studied using a differential scanning calorimeter
(DSC) Mettler Toledo DSC-1. Transformation tempera-
tures were determined from the thermograms measured
at a heating rate of 10 deg/min and a thermal range of
–120 °C to 600 °C.
The surface morphology, hardness and elastic modu-
lus were tested using a Hysitron Triboindenter Ti950
with AFM QScope 250. The measurements were
calculated from the load-displacement data obtained with
the nanoindentation, using a three-sided, pyramidal,
diamond (Berkovich) indenter tip. An area of 40 μm ×
40 μm was analyzed at a scanning frequency of 0.25 Hz.
The microhardness measurement was carried out parallel
to the nanoindentation tests. The Vickers microhardness
measurement was taken at a load of 500 N and a loading
time of 10 s on a 401MVD microhardness tester.
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
796 Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
3 RESULTS AND DISCUSSION
SEM micrographs of the initial metal powders are
presented in Figure 1. The titanium-powder morphology
is irregular, having sharp corners and rough surfaces.
The dispersion of the Ti particles is wide and ranges
from a few μm to even 50 μm. The niobium powder has a
much wider dispersion of the particles in comparison to
the Ti powder.
Some of the particles have the size below 10 μm, and
the shape of these particles is irregular. Simultaneously,
there are also very large particles, the sizes of which are
greater even than 50 μm and these particles have an
irregular, polyhedral shape. As a result of the work-hard-
ening effect and high density of defects such as dislo-
cations and vacancies created during long ball-milling
times, the fracture and welding of powders reach an
equilibrium state. This state leads to the formation of
rather equiaxed and small particles.16,17 However, due to
the higher bonding strength of the finer particles, their
ability to undergo further plastic deformation is de-
creased and a higher force is required to fracture the
small particles during the ball-milling process.18 It is
worth mentioning that the reduction in the particle size at
a given milling time can also be influenced by the
milling technique.
The morphologies of the Ti-26Nb (a/%) powders at
different milling stages are shown in Figure 2. The
powder after the shortest milling time shows globular
particles. In the picture with the higher magnification
(Figure 2b), some cracks were observed, which is the
beginning of the cyclic process of material milling
during high-energy ball milling. In contrast, in the
samples after 70 h of milling, the particles became larger
and the shape of the particles became less globular.
Moreover, the particles of this sample are cold welded
and the powder particles, in particular, are composed of
the welded layers of the material. This behavior of the
material during the high-energy milling is consistent
with the conduct and following steps of mechanical
alloying.19 As shown by the other studies, the mechani-
cal-alloying method leads to a decreased particle size,
the -Ti phase being stabilized by Nb, and an increased
surface area between Ti and Nb. This promotes good
homogeneity of the solid solution.20
Figure 3 shows a summary of diffraction patterns of
the material after the mechanical synthesis, depending on
the milling time of the initial powders. The X-ray
qualitative analysis showed that only  and  phases are
present in the material after milling for 50 h and 70 h.
Change of the profile lines of individual diffraction
patterns are clearly visible.
The broadening of the particle-size distribution at
longer ball-milling times is a typical behavior of high-
energy ball-milling process.17,19,21,22 For the  and 
phases, it was observed that a decrease in the crystallite
size took place simultaneously with an increase in the
milling time (Table 1). In addition, in the case of the 
phase there was a tendency towards a decrease in the
lattice strain <a/a> (Table 1). The estimated size of the
crystallites for the powder after milling for 70 h is 9(1)
nm and 16(2) nm for the  and  phases, respectively.
The observed <a/a> lattice strain values indicate
growth in the  phase (3.43E-03 %) and in the  phase
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803 797
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: X-ray diffraction patterns of the powder after 50 h and 70 h
of the milling process
Figure 1: SEM micrographs of the initial Ti and Nb powders at
different magnifications
Figure 2: SEM micrographs of particles after different milling times:
a), b) 50 h of milling and c), d) 70 h
(2.31E-03 %). The analysis of the diffraction patters
obtained using the Rietveld method showed that that
high-energy milling has an influence on the lattice
parameters.
Table 1: Changes of the average crystallite sizes (D) and lattice
distortions (<a/a>) of the  and  phases of the milled powders
Phase Parameters
Milling time (h)
50 70

D (nm) 22(2) 16(1)
<a/a> (%) 2.76E-03 3.43E-03

D (nm) 15(2) 9(1)
<a/a> (%) 2.04E-03 2.31E-03
Table 2 presents lattice parameters determined for
individual phases and the corresponding ICDD data
sheets. In the case of the  phase, a slight deviation of
the a0 lattice parameters from the catalogue data was
observed. A similar type of cell contraction was
discovered for the a0 parameter of the  phase and for
the c0 parameter, but these changes are smaller, resulting
from the shift systems for the phases of the hexagonal
system. It should be stressed that for the sample after 70
h of milling, the deviation of both lattice parameters can
be noted. It is probably the effect of the alternate cold
welding and deposition of the particles during the
mechanical synthesis.
Table 2: Lattice parameters and contents of  and  phases of the
milled powders
Phase
Lattice parameters (nm)
ICDD*

a0 0.2970
c0 0.4720
 a0 0.3307
* International Centre for Diffraction Data® – a scientific organization
dedicated to collecting, editing, publishing and distributing powder-
diffraction data for the identification of materials
The quantitative phase analysis of the material after
each stage show a successive progressive synthesis of the
as-prepared powders in relation to the  phase (77.5(11))
% mass fraction and 82.6(12) % mass fraction, for 50 h
and 70 h milling times, respectively).
The X-ray qualitative analysis showed that only 
and  phases are present in the samples annealed at
1000 °C for 24 h from the powder previously milled for
50 h and 70 h, as shown in Figure 4. The XRD shows
that there were no obvious diffraction peaks of elemental
Nb remaining in the sintered Ti–26Nb (a/%). The
diffraction peaks attributable to the  and  phases from
the pattern confirm a duplex microstructure. M. Tahara et
al.23 confirmed only the  phase for Ti-26Nb (a/%)
produced by arc melting. The predominant phase in all
the samples was the body-centered-cubic (bcc)  phase.
This indicates that Nb was diffused into Ti, thereby lead-
ing to the formation of  +  phases after the sintering.
This is significant because two-phase Ti alloys, with the
major  phase and the minor  phase, possess good com-
prehensive properties including high yield strength, ex-
cellent corrosion resistance and good fracture toughness.
The presence of unalloyed Ti particles in materials
may arise from different facts. According to S. N.
Patankar and F. H. Froes24, it could be correlated with the
brittleness of the Nb particles, as compared to the Ti
particles, since these brittle particles adhere to the sur-
faces of ductile particles. Another reason could be the
limited solubility of Nb in -Ti. It is known that iso-
morphous -stabilizer elements, such as Nb, exhibit
slower diffusivities in Ti than the eutectoid elements of
Fe, Co and Ni.25
The structural analysis of the diffraction patters
obtained using the Rietveld method showed that the
lattice parameters for all the tested samples become
reduced. In comparison to the ICDD data, it was found
that high-energy milling and annealing have the smallest,
but significant effect on the reduction of the size of a unit
cell of the  and  phases. The quantitative phase anal-
ysis (Table 3) of the material after annealing showed a
significant synthesis of the as-prepared powders in rela-
tion to the  phase 96.6(12) w/% and 98.8(12) w/% for
samples previously milled for 50 h and 70 h, respecti-
vely).
Table 3: Lattice parameters and contents of  and  phases after
annealing (A) the samples at 1000 °C for 24 h from the powder
previously milled for 50 h and 70 h
Phase
Lattice parameters (nm)
Phase
Contents (w/%)
ICDD*
Rietveld
50 h +
A
70 h +
A
50 h +
A
70 h +
A

a0 0.2970 0.2898(2)
0.2891
(2) 
3.4
(6)
1.2
(4)
c0 0.4720 0.4715(4)
0.4710
(4)

96.6
(12)
98.8
(12)
 a0 0.3307 0.3287(2)
0.3286
(2)
* International Centre for Diffraction Data®– a scientific organization
dedicated to collecting, editing, publishing and distributing
powder-diffraction data for the identification of materials
Optical and SEM micrographs of the studied samples
are presented in Figures 5 and 6. They clearly show the
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
798 Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: X-ray diffraction patterns for the samples annealed (+A) at
1000 °C for 24 h from the powder previously milled for 50 h and 70 h
changes in the morphology of the material after sintering
with respect to various milling times. The common
feature of the samples is a hierarchical structure. After
the milling, large-sized particles caused the formation of
a porous material. The pores were large and intercon-
nected (Figures 5a and 5c) in spite of the applied
pressure of 750 MPa during the isostatic pressing.
In the case of the sample milled for 50 h (Figure 5b),
an anomalous grain growth was observed. SEM micro-
graphs (Figure 6) revealed further differences between
the obtained microstructures. The powders ball-milled
for a long time cannot move and deform effectively
during the compaction. This may be primarily due to the
effects of the decrease in the particle size, solid-solution
strengthening and work hardening of the particles during
the ball-milling process.26,27 Amongst these factors, the
work-hardening effect appears to be the main deter-
minant in governing the plastic behavior of the powder
particles during the compaction. In the ball-milling
process, work hardening is caused by the interaction of
dislocations with each other through the continuous
mechanical impact of the balls on the powder particles.16
A quantitative analysis of the microstructure was
conducted including the determination of the cross-sec-
tional area of the grains and the subgrains, as well as the
average circularity of the subgrains of the microsections
based on microscopic images, the results of which are
presented in Tables 4 to 6. Significant changes between
the samples in the average a of the grain cross-sectional
area can also be observed. For the samples milled for
50 h and 70 h, the average section area of grains was
9.09 μm2 and 3.55 μm2, respectively (Table 4). A change
in the time and rotations per minute during milling
causes a difference in the parameter.
Table 4: Results for the cross-sectional area of the grains for the
samples milled for different times
Parameter Sample Minimumvalue
Maximum
value
Average
value Std. dev.
a (μm2)
50 h 0.51 267.57 9.09 12.68
70 h 0.51 32.35 3.55 3.45
Table 5: Results for the circularity of the grains for the samples milled
for different times
Parameter Sample Minimumvalue
Maximum
value
Average
value Std. dev.
fg2
50 h 0.10 1.00 0.60 0.15
70 h 0.11 0.96 0.59 0.15
Table 6: Results for the dimensionless lengthening factor of the grains
for the samples milled for different times
Parameter Sample Minimumvalue
Maximum
value
Average
value Std. dev.
fg1
50 h 0.32 3.31 1.06 0.37
70 h 0.28 3.30 1.06 0.37
Histograms of the size of the area of the section of
the grains in the samples milled for 50 h and 70 h are
presented in Figures 7 and 8, respectively. For the sam-
ple milled for 70 h, a larger fraction of the grains with
the section area in the range of 0.5–35 μm can be ob-
served. Milling for 70 h caused a significant fragmen-
tation of the grains with the section area below 30 μm2.
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803 799
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: a), b) SEM micrographs of Ti-26Nb (a/%) after different
milling times – 50 h and c), d) 70 h and annealing
Figure 5: a), b) Optical micrographs of Ti-26Nb (a/%) at different
magnifications, produced in different milling times – 50 h and c), d)
70 h and annealing
Figure 7: Histogram of the grain average area for the sample milled
for 50 h and annealed
(2.31E-03 %). The analysis of the diffraction patters
obtained using the Rietveld method showed that that
high-energy milling has an influence on the lattice
parameters.
Table 1: Changes of the average crystallite sizes (D) and lattice
distortions (<a/a>) of the  and  phases of the milled powders
Phase Parameters
Milling time (h)
50 70

D (nm) 22(2) 16(1)
<a/a> (%) 2.76E-03 3.43E-03

D (nm) 15(2) 9(1)
<a/a> (%) 2.04E-03 2.31E-03
Table 2 presents lattice parameters determined for
individual phases and the corresponding ICDD data
sheets. In the case of the  phase, a slight deviation of
the a0 lattice parameters from the catalogue data was
observed. A similar type of cell contraction was
discovered for the a0 parameter of the  phase and for
the c0 parameter, but these changes are smaller, resulting
from the shift systems for the phases of the hexagonal
system. It should be stressed that for the sample after 70
h of milling, the deviation of both lattice parameters can
be noted. It is probably the effect of the alternate cold
welding and deposition of the particles during the
mechanical synthesis.
Table 2: Lattice parameters and contents of  and  phases of the
milled powders
Phase
Lattice parameters (nm)
ICDD*

a0 0.2970
c0 0.4720
 a0 0.3307
* International Centre for Diffraction Data® – a scientific organization
dedicated to collecting, editing, publishing and distributing powder-
diffraction data for the identification of materials
The quantitative phase analysis of the material after
each stage show a successive progressive synthesis of the
as-prepared powders in relation to the  phase (77.5(11))
% mass fraction and 82.6(12) % mass fraction, for 50 h
and 70 h milling times, respectively).
The X-ray qualitative analysis showed that only 
and  phases are present in the samples annealed at
1000 °C for 24 h from the powder previously milled for
50 h and 70 h, as shown in Figure 4. The XRD shows
that there were no obvious diffraction peaks of elemental
Nb remaining in the sintered Ti–26Nb (a/%). The
diffraction peaks attributable to the  and  phases from
the pattern confirm a duplex microstructure. M. Tahara et
al.23 confirmed only the  phase for Ti-26Nb (a/%)
produced by arc melting. The predominant phase in all
the samples was the body-centered-cubic (bcc)  phase.
This indicates that Nb was diffused into Ti, thereby lead-
ing to the formation of  +  phases after the sintering.
This is significant because two-phase Ti alloys, with the
major  phase and the minor  phase, possess good com-
prehensive properties including high yield strength, ex-
cellent corrosion resistance and good fracture toughness.
The presence of unalloyed Ti particles in materials
may arise from different facts. According to S. N.
Patankar and F. H. Froes24, it could be correlated with the
brittleness of the Nb particles, as compared to the Ti
particles, since these brittle particles adhere to the sur-
faces of ductile particles. Another reason could be the
limited solubility of Nb in -Ti. It is known that iso-
morphous -stabilizer elements, such as Nb, exhibit
slower diffusivities in Ti than the eutectoid elements of
Fe, Co and Ni.25
The structural analysis of the diffraction patters
obtained using the Rietveld method showed that the
lattice parameters for all the tested samples become
reduced. In comparison to the ICDD data, it was found
that high-energy milling and annealing have the smallest,
but significant effect on the reduction of the size of a unit
cell of the  and  phases. The quantitative phase anal-
ysis (Table 3) of the material after annealing showed a
significant synthesis of the as-prepared powders in rela-
tion to the  phase 96.6(12) w/% and 98.8(12) w/% for
samples previously milled for 50 h and 70 h, respecti-
vely).
Table 3: Lattice parameters and contents of  and  phases after
annealing (A) the samples at 1000 °C for 24 h from the powder
previously milled for 50 h and 70 h
Phase
Lattice parameters (nm)
Phase
Contents (w/%)
ICDD*
Rietveld
50 h +
A
70 h +
A
50 h +
A
70 h +
A

a0 0.2970 0.2898(2)
0.2891
(2) 
3.4
(6)
1.2
(4)
c0 0.4720 0.4715(4)
0.4710
(4)

96.6
(12)
98.8
(12)
 a0 0.3307 0.3287(2)
0.3286
(2)
* International Centre for Diffraction Data®– a scientific organization
dedicated to collecting, editing, publishing and distributing
powder-diffraction data for the identification of materials
Optical and SEM micrographs of the studied samples
are presented in Figures 5 and 6. They clearly show the
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
798 Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: X-ray diffraction patterns for the samples annealed (+A) at
1000 °C for 24 h from the powder previously milled for 50 h and 70 h
changes in the morphology of the material after sintering
with respect to various milling times. The common
feature of the samples is a hierarchical structure. After
the milling, large-sized particles caused the formation of
a porous material. The pores were large and intercon-
nected (Figures 5a and 5c) in spite of the applied
pressure of 750 MPa during the isostatic pressing.
In the case of the sample milled for 50 h (Figure 5b),
an anomalous grain growth was observed. SEM micro-
graphs (Figure 6) revealed further differences between
the obtained microstructures. The powders ball-milled
for a long time cannot move and deform effectively
during the compaction. This may be primarily due to the
effects of the decrease in the particle size, solid-solution
strengthening and work hardening of the particles during
the ball-milling process.26,27 Amongst these factors, the
work-hardening effect appears to be the main deter-
minant in governing the plastic behavior of the powder
particles during the compaction. In the ball-milling
process, work hardening is caused by the interaction of
dislocations with each other through the continuous
mechanical impact of the balls on the powder particles.16
A quantitative analysis of the microstructure was
conducted including the determination of the cross-sec-
tional area of the grains and the subgrains, as well as the
average circularity of the subgrains of the microsections
based on microscopic images, the results of which are
presented in Tables 4 to 6. Significant changes between
the samples in the average a of the grain cross-sectional
area can also be observed. For the samples milled for
50 h and 70 h, the average section area of grains was
9.09 μm2 and 3.55 μm2, respectively (Table 4). A change
in the time and rotations per minute during milling
causes a difference in the parameter.
Table 4: Results for the cross-sectional area of the grains for the
samples milled for different times
Parameter Sample Minimumvalue
Maximum
value
Average
value Std. dev.
a (μm2)
50 h 0.51 267.57 9.09 12.68
70 h 0.51 32.35 3.55 3.45
Table 5: Results for the circularity of the grains for the samples milled
for different times
Parameter Sample Minimumvalue
Maximum
value
Average
value Std. dev.
fg2
50 h 0.10 1.00 0.60 0.15
70 h 0.11 0.96 0.59 0.15
Table 6: Results for the dimensionless lengthening factor of the grains
for the samples milled for different times
Parameter Sample Minimumvalue
Maximum
value
Average
value Std. dev.
fg1
50 h 0.32 3.31 1.06 0.37
70 h 0.28 3.30 1.06 0.37
Histograms of the size of the area of the section of
the grains in the samples milled for 50 h and 70 h are
presented in Figures 7 and 8, respectively. For the sam-
ple milled for 70 h, a larger fraction of the grains with
the section area in the range of 0.5–35 μm can be ob-
served. Milling for 70 h caused a significant fragmen-
tation of the grains with the section area below 30 μm2.
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803 799
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: a), b) SEM micrographs of Ti-26Nb (a/%) after different
milling times – 50 h and c), d) 70 h and annealing
Figure 5: a), b) Optical micrographs of Ti-26Nb (a/%) at different
magnifications, produced in different milling times – 50 h and c), d)
70 h and annealing
Figure 7: Histogram of the grain average area for the sample milled
for 50 h and annealed
The circularity of grains was 0.60 for 50 h of milling
and 0.59 for 70 h of milling, which corresponds to the
dimensionless lengthening factor for both samples
(Tables 5 to 6). The coefficient of the grain variability
was also determined. For all the samples milled for 50 h,
it was 1.4, and for the sample milled for 70 h, it was 0.9.
It means that for the sample milled for a shorter time, the
size of the cross-sectional area is much more diversified
in comparison to the sample milled for 70 h. The longer
milling time (70 h) caused a greater fragmentation of the
material. The difference in the average section area of
the grains was significant, especially when compared to
the size of the powder after milling. No significant diffe-
rences in the size of the particles depending on the
milling time were noticed. The number of the elements
that were tested per area unit of the image was also
determined; for the samples annealed and previously
milled for 50 h and 70 h, the number of elements per
area was estimated to be 82379 [ ]1 2/mm and 199951
[ ]1 2/mm respectively.
Distribution maps of the elements for the sintered
samples revealed single regions rich in titanium in the
case of the material milled for 50 h (Figure 9) Also, a
slightly higher concentration of Ti was observed inside
the grains. In contrast, milling for 70 h allowed us to
obtain a material with an even distribution of the ele-
ments on its surface (Figure 10). In the case of this alloy,
the increase in the milling time allows a more
homogeneous structure.
Figures 11 and 12 show the effect of an analysis with
the Hysitron Tribointender Ti950 with AFM QScope
250, which allowed us to compare the surface topogra-
phies of the produced samples. For the material milled
for 50 h and sintered, a higher diversification of the
sample surface was observed. On the sample, initially
synthetized for 70 h of milling, we can observe smaller
and more regular grains in comparison to the sample ex-
posed to the shorter milling time, which was confirmed
with a stereological analysis.
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
800 Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: Histograms of the grain average area for the sample milled
for 70 h and annealed. The upper histogram presents the distribution
of the average area of the grains smaller than 36 μm2 for this sample
Figure 10: Distribution maps of the elements for the sample milled for 70 h and annealed
Figure 9: Distribution maps of the elements for the sample milled for 50 h and annealed
Another important aspect of the produced material is
its properties. DSC results showed (Figure 13) the
presence of an endothermic peak on the heating curve for
both samples. In the case of the sample milled for 70 h,
the peak was slightly shifted toward higher temperatures.
The peaks occurred at about 490 °C and 520 °C for the
samples milled for 50 h and 70 h, respectively. The pre-
sence of the peaks can indicate the presence of a partial
phase transformation   . Moreover, the milling time
has an impact on the temperature of transformation.
However, no significant changes were noticed on the
cooling curves.
The nanoindentation method allowed the determina-
tion of the reduced Young’s modulus and hardness of the
material, and the results of this analysis of the material
after different milling times and sintering are presented
in Table 7.
Table 7: Results of the nanoindentation analysis – hardness and
modulus of both samples
Samples Contact depth(nm)
Hardness
(GPa) Modulus (GPa)
50 h 89(14) 0.99(0.31) 48(19)
70 h 42(12) 3.86(1.71) 95(32)
The measurement is based on the curves of the load
and the penetration depth of the indenter (Figures 14 and
15). After 50 h of milling, the material obtained much
lower values of hardness and modulus in comparison to
the material milled for 70 h. For the samples milled for
50 h and 70 h, the hardness was 0.99±0.31 GPa and
3.86±1.71GPa, respectively. A similar situation applied
to the values of the modulus. For the material milled for
50 h, it was 48±19 GPa, and for the material milled for
70 h, it was 95±32 GPa. Such low modulus values,
especially in the case of the sample with the shorter
milling time, are very promising as they are close to the
value of the modulus of a human bone.
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803 801
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 13: DSC analysis of the samples milled for 50 h and 70 h and
also annealed
Figure 12: Surface morphology of the area of 40 μm × 40 μm for the
sample milled for 70 h and annealed
Figure 11: Surface morphology of the area of 40 μm × 40 μm for the
sample milled for 50 h and annealed
The circularity of grains was 0.60 for 50 h of milling
and 0.59 for 70 h of milling, which corresponds to the
dimensionless lengthening factor for both samples
(Tables 5 to 6). The coefficient of the grain variability
was also determined. For all the samples milled for 50 h,
it was 1.4, and for the sample milled for 70 h, it was 0.9.
It means that for the sample milled for a shorter time, the
size of the cross-sectional area is much more diversified
in comparison to the sample milled for 70 h. The longer
milling time (70 h) caused a greater fragmentation of the
material. The difference in the average section area of
the grains was significant, especially when compared to
the size of the powder after milling. No significant diffe-
rences in the size of the particles depending on the
milling time were noticed. The number of the elements
that were tested per area unit of the image was also
determined; for the samples annealed and previously
milled for 50 h and 70 h, the number of elements per
area was estimated to be 82379 [ ]1 2/mm and 199951
[ ]1 2/mm respectively.
Distribution maps of the elements for the sintered
samples revealed single regions rich in titanium in the
case of the material milled for 50 h (Figure 9) Also, a
slightly higher concentration of Ti was observed inside
the grains. In contrast, milling for 70 h allowed us to
obtain a material with an even distribution of the ele-
ments on its surface (Figure 10). In the case of this alloy,
the increase in the milling time allows a more
homogeneous structure.
Figures 11 and 12 show the effect of an analysis with
the Hysitron Tribointender Ti950 with AFM QScope
250, which allowed us to compare the surface topogra-
phies of the produced samples. For the material milled
for 50 h and sintered, a higher diversification of the
sample surface was observed. On the sample, initially
synthetized for 70 h of milling, we can observe smaller
and more regular grains in comparison to the sample ex-
posed to the shorter milling time, which was confirmed
with a stereological analysis.
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
800 Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: Histograms of the grain average area for the sample milled
for 70 h and annealed. The upper histogram presents the distribution
of the average area of the grains smaller than 36 μm2 for this sample
Figure 10: Distribution maps of the elements for the sample milled for 70 h and annealed
Figure 9: Distribution maps of the elements for the sample milled for 50 h and annealed
Another important aspect of the produced material is
its properties. DSC results showed (Figure 13) the
presence of an endothermic peak on the heating curve for
both samples. In the case of the sample milled for 70 h,
the peak was slightly shifted toward higher temperatures.
The peaks occurred at about 490 °C and 520 °C for the
samples milled for 50 h and 70 h, respectively. The pre-
sence of the peaks can indicate the presence of a partial
phase transformation   . Moreover, the milling time
has an impact on the temperature of transformation.
However, no significant changes were noticed on the
cooling curves.
The nanoindentation method allowed the determina-
tion of the reduced Young’s modulus and hardness of the
material, and the results of this analysis of the material
after different milling times and sintering are presented
in Table 7.
Table 7: Results of the nanoindentation analysis – hardness and
modulus of both samples
Samples Contact depth(nm)
Hardness
(GPa) Modulus (GPa)
50 h 89(14) 0.99(0.31) 48(19)
70 h 42(12) 3.86(1.71) 95(32)
The measurement is based on the curves of the load
and the penetration depth of the indenter (Figures 14 and
15). After 50 h of milling, the material obtained much
lower values of hardness and modulus in comparison to
the material milled for 70 h. For the samples milled for
50 h and 70 h, the hardness was 0.99±0.31 GPa and
3.86±1.71GPa, respectively. A similar situation applied
to the values of the modulus. For the material milled for
50 h, it was 48±19 GPa, and for the material milled for
70 h, it was 95±32 GPa. Such low modulus values,
especially in the case of the sample with the shorter
milling time, are very promising as they are close to the
value of the modulus of a human bone.
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803 801
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 13: DSC analysis of the samples milled for 50 h and 70 h and
also annealed
Figure 12: Surface morphology of the area of 40 μm × 40 μm for the
sample milled for 70 h and annealed
Figure 11: Surface morphology of the area of 40 μm × 40 μm for the
sample milled for 50 h and annealed
The nanosized indentation measurements of the
mechanical properties are from the selected specific
areas, which neutralize the effects of the pores or defects
on the results. A large measurement uncertainty is even
more surprising. Assuming that the material is mechani-
cally widely heterogeneous, it seems necessary to apply
heat treatment to the material. The microhardness
measurement of the sample milled for 50 h revealed a
higher value of the microhardness and it is 288.6(±31.0)
HV0.5, which is around 2.82 GPa. For the sample milled
for 70 h, the revealed value is 256.2(±14.6) HV0.5,
which is approximately 2.51 GPa. However, we observed
a relatively high uncertainty of the measurement, which
made it difficult to clearly evaluate the differences in the
microhardness measurement.
4 CONCLUSIONS
Based on the investigation into the microstructure
and mechanical properties of the Ti-26Nb (a/%) alloy,
the main conclusions can be drawn:
• It was proved that the microstructure, the degree of
porosity and mechanical properties can be adjusted
by means of mechanical alloying.
• A longer milling time (70 h) resulted in an increased
degree of structure fragmentation. Young’s modulus
increased twice as well.
• The porous samples sintered from the powders milled
for 50 hours have a very low elastic modulus of
48(19) GPa, which is similar to that of natural bones.
• Ball-milling for 50 h led to the formation of a rela-
tively equiaxed shape and powder particles of a
smaller size and with a more asymmetrical particle-
size distribution than that of the powders ball milled
for 70 h.
• Based on XRD, it was found that the process of
high-energy milling for 50 h and 70 h makes it pos-
sible to obtain a solid nanocrystalline solution  + .
Acknowledgments
This work was supported by the Polish National
Science Centre (Polish: Narodowe Centrum Nauki, abbr.
NCN) under the research project no. UMO-2011/03/
D/ST8/04884
5 REFERENCES
1 M. Long, H. J. Rack, Titanium alloys in total joint replacement – a
materials science perspective, Biomaterials, 19 (1998) 18,
1621–1639, doi:10.1016/S0142-9612(97)00146-4
2 Y. Li, C. Yang, H. Zhao, S. Qu, X. Li, Y. Li, New Developments of
Ti-Based Alloys for Biomedical Applications, Mater., 7 (2014) 3,
1709–1800, doi:10.3390/ma7031709
3 M. Niinomi, Mechanical biocompatibilities of titanium alloys for
biomedical application, J. Mech. Behav. Biomed. Mater., 1 (2008) 1,
30–42, doi:10.1016/j.jmbbm.2007.07.001
4 A. Biesiekierski, J. Wang, M. Gepreel, C. Wena, A new look at
biomedical Ti-based shape memory alloys, Acta Biomater. 8 (2012)
5, 1661–1669, doi:10.1016/j.actbio.2012.01.018
5 S. Miyazaki, H.Y. Kim, Basic characteristics of titanium–nickel
(Ti–Ni)-based and titanium–niobium (Ti–Nb)-based alloys, Shape
Memory and Superelastic Alloys, Appl. Technol., (2011), 15–42,
doi:10.1016/B978-1-84569-707-5.50002-X
6 M. Tahara, H. Y. Kim, H. Hosoda, S. Miyazaki, Cyclic deformation
behavior of Ti-26 at.% Nb alloy, Acta Mater., 25 (2009), 2461–2469,
doi:10.1016/j.actamat.2009.01.037
7 H. Y. Kim, Y. Ikehara, J. I. Kim, H. Hosoda, S. Miyazaki, Marten-
sitic transformation, shape memory effect and superelasticity of
Ti–Nb binary alloys, Acta Mater. 54 (2006), 2419–2429,
doi:10.1016/j.actamat.2006.01.019
8 G. Dercz, I. Matu³a, M. Zubko, A. Liberska, Structure characteri-
zation of biomedical Ti-Mo-Sn prepared by mechanical alloying
method, Acta Phys. Pol. A, 130 (2016), 1029–1032, doi:10.12693/
APhysPolA.130.1029
9 I. Matu³a, G. Dercz, M. Zubko, L. Paj¹k, Influence of high energy
milling time on the Ti-50Ta biomedical alloy structure, Acta Phys.
Pol. A, 130 (2016), 1033–1036, doi:10.12693/APhysPolA.130.1033
10 X. Rao, C. L. Chu, Y. Y. Zheng, Phase composition, microstructure,
and mechanical properties of porous Ti-Nb-Zr alloys prepared by a
two-step foaming powder metallurgy method, J. Mech. Behav.
Biomed., 34 (2014), 27–36, doi:10.1016/j.jmbbm.2014.02.001
11 G. Ryan, A. Pandit, D. P. Apatsidis, Fabrication methods of porous
metals for use in orthopaedic applications, Biomaterials, 27 (2006)
13, 2651–2670, doi:10.1016/j.biomaterials.2005.12.002
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
802 Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 15: Diagram of reduced modulus to contact depth for both
samples, after 50 h and 70 h of milling
Figure 14: Diagram of hardness to contact depth for both samples,
after 50 h and 70 h of milling
12 H. M. Rietveld, A Profile Refinement Method for Nuclear and
Magnetic Structures, J. Appl. Cryst., 3 (1969), 65–69, doi:10.1107/
S0021889869006558
13 R. J. Hill, C. J. Howard, Quantitative phase analysis from neutron
powder diffraction data using the Rietveld method. J. Appl. Cryst.,
20 (1987), 467–474, doi:10.1107/S0021889887086199
14 G. Dercz, D. Oleszak, K. Prusik, L. Paj¹k, Rietveld-based quanti-
tative analysis of multiphase powders with nanocrystalline NiAl and
FeAl phases, Rev. Adv. Mater. Sci., 8 (2008), 764–768
15 Y. Li, Y. Cui, F. Zhanga, H. Xua, Shape memory behavior in Ti–Zr
alloys, Scripta Mater., 64 (2011) 6, 584–587, doi:/10.1016/
j.scriptamat.2010.11.048
16 A. Nouri, P. D. Hodgson, C. Wen, Effect of ball-milling time on the
structural characteristics of biomedical porous Ti–Sn–Nb alloy,
Mater. Sci. Eng. C, 31 (2011), 921–928, doi:10.1016/j.msec.2011.
02.011
17 G. Dercz, I. Matu³a, M. Zubko, J. Dercz, Phase composition and
microstructure of new Ti–Ta–Nb–Zr biomedical alloys prepared by
mechanical alloying method, Powder Diffr., (2017), 1–7,
doi:10.1017/S0885715617000045
18 L. Lü, M. O. Lai, Mechanical Alloying, Kluwer Academic Pub-
lishers, Boston, 1998
19 C. Suryanarayana, Mechanical alloying and milling, Prog. Mater.
Sci., 46 (2001) 1–2, 1–184, doi:10.1016/S0079-6425(99)00010-9
20 A. Omran, K. Woo, H. B. Lee, Mechanical properties of –Ti-
35Nb-2.5Sn alloy synthesized by mechanical alloying and pulsed
current activated sintering, Metall. Mater. Trans. A, 43 (2012) 12,
4866–4874, doi:10.1007/s11661-012-1298-y
21 G. Dercz, B. Formanek, K. Prusik, L. Paj¹k, Microstructure of
Ni(Cr)-TiC-Cr3C2-Cr7C3 composite powder, J. Mater. Process. Tech.,
162 (2005), 15–19, doi:10.1016/jmatprotec.2005.02.004
22 G. Dercz, L. Paj¹k, B. Formanek, Dispersion analysis of NiAl-
TiC-Al2O3 composite powder ground in a high-energy attritorial mill,
J. Mater. Process. Tech., 175 (2006), 334–337, doi:10.1016/
j.jmatprotec.2005.04.060
23 M. Tahara, H. Y. Kim, H. Hosoda, S. Miyazaki, Cyclic deformation
behavior of a Ti–26 at.% Nb alloy, Acta Mater., 57 (2009),
2461–2469, doi:10.1016/j.actamat.2009.01.037
24 E. K. Molchanova, Phase Diagrams of Titanium Alloys (Translation
of Atlas Diagram Sostoyaniya Titanovyk Splavov), Israel Program
for Scientific Translations, Jerusalem, 1965, 116
25 S. N. Patankar, F. H. Froes, Transformation of mechanically alloyed
Nb-Sn powder to Nb3Sn, Metall. Mater. Trans. A, 35 (2004),
3009–3012, doi:10.1007/s11661-004-0248-8
26 P. J. James, Particle Deformation During Cold Isostatic Pressing of
Metal Powders, Powder Metall., 20 (1977), 199–204, doi:10.1179/
pom.1977.20.4.199
27 L. Lu, M. O. Lai, G. Li, Influence of sintering process on the mecha-
nical property and microstructure of ball milled composite compacts,
Mater. Res. Bull., 31 (1996), 453–464, doi:10.1016/S0025-5408(96)
00024-4
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803 803
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
The nanosized indentation measurements of the
mechanical properties are from the selected specific
areas, which neutralize the effects of the pores or defects
on the results. A large measurement uncertainty is even
more surprising. Assuming that the material is mechani-
cally widely heterogeneous, it seems necessary to apply
heat treatment to the material. The microhardness
measurement of the sample milled for 50 h revealed a
higher value of the microhardness and it is 288.6(±31.0)
HV0.5, which is around 2.82 GPa. For the sample milled
for 70 h, the revealed value is 256.2(±14.6) HV0.5,
which is approximately 2.51 GPa. However, we observed
a relatively high uncertainty of the measurement, which
made it difficult to clearly evaluate the differences in the
microhardness measurement.
4 CONCLUSIONS
Based on the investigation into the microstructure
and mechanical properties of the Ti-26Nb (a/%) alloy,
the main conclusions can be drawn:
• It was proved that the microstructure, the degree of
porosity and mechanical properties can be adjusted
by means of mechanical alloying.
• A longer milling time (70 h) resulted in an increased
degree of structure fragmentation. Young’s modulus
increased twice as well.
• The porous samples sintered from the powders milled
for 50 hours have a very low elastic modulus of
48(19) GPa, which is similar to that of natural bones.
• Ball-milling for 50 h led to the formation of a rela-
tively equiaxed shape and powder particles of a
smaller size and with a more asymmetrical particle-
size distribution than that of the powders ball milled
for 70 h.
• Based on XRD, it was found that the process of
high-energy milling for 50 h and 70 h makes it pos-
sible to obtain a solid nanocrystalline solution  + .
Acknowledgments
This work was supported by the Polish National
Science Centre (Polish: Narodowe Centrum Nauki, abbr.
NCN) under the research project no. UMO-2011/03/
D/ST8/04884
5 REFERENCES
1 M. Long, H. J. Rack, Titanium alloys in total joint replacement – a
materials science perspective, Biomaterials, 19 (1998) 18,
1621–1639, doi:10.1016/S0142-9612(97)00146-4
2 Y. Li, C. Yang, H. Zhao, S. Qu, X. Li, Y. Li, New Developments of
Ti-Based Alloys for Biomedical Applications, Mater., 7 (2014) 3,
1709–1800, doi:10.3390/ma7031709
3 M. Niinomi, Mechanical biocompatibilities of titanium alloys for
biomedical application, J. Mech. Behav. Biomed. Mater., 1 (2008) 1,
30–42, doi:10.1016/j.jmbbm.2007.07.001
4 A. Biesiekierski, J. Wang, M. Gepreel, C. Wena, A new look at
biomedical Ti-based shape memory alloys, Acta Biomater. 8 (2012)
5, 1661–1669, doi:10.1016/j.actbio.2012.01.018
5 S. Miyazaki, H.Y. Kim, Basic characteristics of titanium–nickel
(Ti–Ni)-based and titanium–niobium (Ti–Nb)-based alloys, Shape
Memory and Superelastic Alloys, Appl. Technol., (2011), 15–42,
doi:10.1016/B978-1-84569-707-5.50002-X
6 M. Tahara, H. Y. Kim, H. Hosoda, S. Miyazaki, Cyclic deformation
behavior of Ti-26 at.% Nb alloy, Acta Mater., 25 (2009), 2461–2469,
doi:10.1016/j.actamat.2009.01.037
7 H. Y. Kim, Y. Ikehara, J. I. Kim, H. Hosoda, S. Miyazaki, Marten-
sitic transformation, shape memory effect and superelasticity of
Ti–Nb binary alloys, Acta Mater. 54 (2006), 2419–2429,
doi:10.1016/j.actamat.2006.01.019
8 G. Dercz, I. Matu³a, M. Zubko, A. Liberska, Structure characteri-
zation of biomedical Ti-Mo-Sn prepared by mechanical alloying
method, Acta Phys. Pol. A, 130 (2016), 1029–1032, doi:10.12693/
APhysPolA.130.1029
9 I. Matu³a, G. Dercz, M. Zubko, L. Paj¹k, Influence of high energy
milling time on the Ti-50Ta biomedical alloy structure, Acta Phys.
Pol. A, 130 (2016), 1033–1036, doi:10.12693/APhysPolA.130.1033
10 X. Rao, C. L. Chu, Y. Y. Zheng, Phase composition, microstructure,
and mechanical properties of porous Ti-Nb-Zr alloys prepared by a
two-step foaming powder metallurgy method, J. Mech. Behav.
Biomed., 34 (2014), 27–36, doi:10.1016/j.jmbbm.2014.02.001
11 G. Ryan, A. Pandit, D. P. Apatsidis, Fabrication methods of porous
metals for use in orthopaedic applications, Biomaterials, 27 (2006)
13, 2651–2670, doi:10.1016/j.biomaterials.2005.12.002
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
802 Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 15: Diagram of reduced modulus to contact depth for both
samples, after 50 h and 70 h of milling
Figure 14: Diagram of hardness to contact depth for both samples,
after 50 h and 70 h of milling
12 H. M. Rietveld, A Profile Refinement Method for Nuclear and
Magnetic Structures, J. Appl. Cryst., 3 (1969), 65–69, doi:10.1107/
S0021889869006558
13 R. J. Hill, C. J. Howard, Quantitative phase analysis from neutron
powder diffraction data using the Rietveld method. J. Appl. Cryst.,
20 (1987), 467–474, doi:10.1107/S0021889887086199
14 G. Dercz, D. Oleszak, K. Prusik, L. Paj¹k, Rietveld-based quanti-
tative analysis of multiphase powders with nanocrystalline NiAl and
FeAl phases, Rev. Adv. Mater. Sci., 8 (2008), 764–768
15 Y. Li, Y. Cui, F. Zhanga, H. Xua, Shape memory behavior in Ti–Zr
alloys, Scripta Mater., 64 (2011) 6, 584–587, doi:/10.1016/
j.scriptamat.2010.11.048
16 A. Nouri, P. D. Hodgson, C. Wen, Effect of ball-milling time on the
structural characteristics of biomedical porous Ti–Sn–Nb alloy,
Mater. Sci. Eng. C, 31 (2011), 921–928, doi:10.1016/j.msec.2011.
02.011
17 G. Dercz, I. Matu³a, M. Zubko, J. Dercz, Phase composition and
microstructure of new Ti–Ta–Nb–Zr biomedical alloys prepared by
mechanical alloying method, Powder Diffr., (2017), 1–7,
doi:10.1017/S0885715617000045
18 L. Lü, M. O. Lai, Mechanical Alloying, Kluwer Academic Pub-
lishers, Boston, 1998
19 C. Suryanarayana, Mechanical alloying and milling, Prog. Mater.
Sci., 46 (2001) 1–2, 1–184, doi:10.1016/S0079-6425(99)00010-9
20 A. Omran, K. Woo, H. B. Lee, Mechanical properties of –Ti-
35Nb-2.5Sn alloy synthesized by mechanical alloying and pulsed
current activated sintering, Metall. Mater. Trans. A, 43 (2012) 12,
4866–4874, doi:10.1007/s11661-012-1298-y
21 G. Dercz, B. Formanek, K. Prusik, L. Paj¹k, Microstructure of
Ni(Cr)-TiC-Cr3C2-Cr7C3 composite powder, J. Mater. Process. Tech.,
162 (2005), 15–19, doi:10.1016/jmatprotec.2005.02.004
22 G. Dercz, L. Paj¹k, B. Formanek, Dispersion analysis of NiAl-
TiC-Al2O3 composite powder ground in a high-energy attritorial mill,
J. Mater. Process. Tech., 175 (2006), 334–337, doi:10.1016/
j.jmatprotec.2005.04.060
23 M. Tahara, H. Y. Kim, H. Hosoda, S. Miyazaki, Cyclic deformation
behavior of a Ti–26 at.% Nb alloy, Acta Mater., 57 (2009),
2461–2469, doi:10.1016/j.actamat.2009.01.037
24 E. K. Molchanova, Phase Diagrams of Titanium Alloys (Translation
of Atlas Diagram Sostoyaniya Titanovyk Splavov), Israel Program
for Scientific Translations, Jerusalem, 1965, 116
25 S. N. Patankar, F. H. Froes, Transformation of mechanically alloyed
Nb-Sn powder to Nb3Sn, Metall. Mater. Trans. A, 35 (2004),
3009–3012, doi:10.1007/s11661-004-0248-8
26 P. J. James, Particle Deformation During Cold Isostatic Pressing of
Metal Powders, Powder Metall., 20 (1977), 199–204, doi:10.1179/
pom.1977.20.4.199
27 L. Lu, M. O. Lai, G. Li, Influence of sintering process on the mecha-
nical property and microstructure of ball milled composite compacts,
Mater. Res. Bull., 31 (1996), 453–464, doi:10.1016/S0025-5408(96)
00024-4
G. DERCZ, I. MATU£A: EFFECT OF BALL MILLING ON THE PROPERTIES OF THE POROUS Ti–26Nb ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 795–803 803
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
805–811
EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE
MECHANICAL PROPERTIES AND EROSIVE-WEAR BEHAVIOR
OF A WOOD-DUST-PARTICLE-REINFORCED PHENOL
FORMALDEHYDE COMPOSITE
VPLIVI DODATKA KOKOSOVIH VLAKEN
FENOL-FORMALDEHIDNEMU KOMPOZITU, OJA^ANEM Z
LESNIM PRAHOM, NA NJEGOVE MEHANSKE LASTNOSTI IN
EROZIJSKO OBRABO
Arul Sujin Jose1, Ayyanar Athijayamani2, Kalimuthu Ramanathan3,
Susaiyappan Sidhardhan4
1Lourdes Mount College of Engineering and Technology, Department of Mechanical Engineering, Kanyakumari, Tamilnadu, India
2Government College of Engineering, Department of Mechanical Engineering, Bodinayakkanur, Tamilnadu, India
3Alagappa Chettiar College of Engineering and Technology, Department of Mechanical Engineering, Karaikudi, Tamilnadu, India
4Government College of Engineering, Department of Civil Engineering, Tirunelveli, Tamilnadu, India
athimania@gmail.com
Prejem rokopisa – received: 2016-09-23; sprejem za objavo – accepted for publication: 2017-01-22
doi:10.17222/mit.2016.284
Several attempts were made to investigate the effects of various process parameters on the mechanical properties and wear
behavior of synthetic and natural cellulosic fibers and also particle-reinforced polymer composites. However, very few studies
were carried out on the effects of various process parameters on the mechanical and wear behavior of phenol formaldehyde (PF)
composites reinforced with natural cellulosic fibers and particles. Therefore, in the present study, an attempt was made to
observe the effects of various process parameters on the mechanical and wear behavior of wood-dust (WD) and coir-pith (CP)
particle-reinforced resole-type PF composites. First, the mechanical properties of a WD/PF composite were studied based on the
content of CP particles. Then, the erosive-wear behavior of the WD/PF composite was studied with respect to five different
parameters such particle content, erodent size, impact velocity, impingement angle, and standoff distance. The erosive
experiments were carried out for five different parameters based on the Taguchi experimental design (L27). The results show that
the mechanical properties of the WD/PF composite increase with an addition of CP particles. The increment in the composite
modulus was higher than that of the composite strength. The erosive test results indicate that the erosion-wear rate is affected by
the particle content, impingement angle, erodent size and impact velocity. Brittle-erosion behavior was identified on the surface
of the composite with a heavy erosive wear occurring at a 60° impingement angle.
Keywords: biowaste particles, phenol formaldehyde, composites, mechanical properties, erosive-wear resistance, Taguchi
method
Izvedenih je bilo `e kar nekaj poizkusov v zvezi z u~inki razli~nih procesnih parametrov na mehanske lastnosti in obrabo
polimernih kompozitov oja~anih s sinteti~nimi in naravnimi celuloznimi vlakni in/ali delci. Toda zelo malo raziskav je bilo
izvedenih glede vpliva razli~nih procesnih parametrov na mehanske lastnosti in obrabo fenolformaldehidnih (angl. PF)
kompozitov, oja~anih z naravnimi celuloznimi vlakni in delci. Tako je v pri~ujo~em delu predstavljen vpliv razli~nih procesnih
parametrov na mehanske lastnosti in obrabo PF kompozitov, ki so bili oja~ani z delci lesnega prahu (angl. WD) in delci kokosa
(angl. CP). Te vrste kompozitov se uporabljajo za izdelavo podplatov ~evljev. Najprej so bile dolo~ene mehanske lastnosti
WD/PF kompozitov glede na vsebnost CP delcev. Sledili so preizkusi in analize erozijske obrabe WD/PF kompozitov glede na
vsebnost (koli~ino) delcev v kompozitu, velikost, hitrost in razdaljo u~inkovanja erozijskega sredsta ter njegov vpadni kot.
Preizkusi so temeljili na analizi s Taguchijevo metodo (L27) s petimi razli~nimi parametri. Rezultati so pokazali, da se mehanske
lastnosti WD/PF kompozitov izbolj{ujejo z dodajanjem CP delcev. Povi{anje modula kompozitov je bilo ve~je od pove~anja
trdnosti kompozita. Erozijski testi ka`ejo, da je hitrost erozijske obrabe posledica vseh procesnih parametrov, to je: vsebnosti
delcev, udarnega kota, hitrosti in velikosti delcev izbranega erozijskega sredstva. Najve~ja obraba zaradi erozije je bila dose`ena
(ugotovljena) pri 60 stopinjskem vpadnem kotu abrazijskega sredstva z nastalimi po{kodbami krhkega zna~aja.
Klju~ne besede: delci bioodpadkov, fenolformaldehid, kompoziti, mehanske lastnosti, odpornost proti erozijski obrabi, Taguchi
metoda
1 INTRODUCTION
Recently, polymer composites reinforced with syn-
thetic materials have been replaced with polymer com-
posites reinforced with bio-based natural materials. The
bio-based natural materials have many advantages over
the synthetic materials like renewability, biodegrada-
bility, abundant availability, low costs, etc.1–3 In India,
particularly in the South Indian region, the biowaste
materials like wood dust, coir pith, groundnut shell,
coconut shell, cashew nut shell, etc. are abundantly
available because in that region, coconut, groundnut and
cashew nut are cultivated in large amounts. A number of
timber and oil mills are also available in the Southern
region of Tamilnadu, India. Therefore, bioparticles are
thrown away after producing useful materials and
Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811 805
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:620.1:620.163 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)805(2017)
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
805–811
EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE
MECHANICAL PROPERTIES AND EROSIVE-WEAR BEHAVIOR
OF A WOOD-DUST-PARTICLE-REINFORCED PHENOL
FORMALDEHYDE COMPOSITE
VPLIVI DODATKA KOKOSOVIH VLAKEN
FENOL-FORMALDEHIDNEMU KOMPOZITU, OJA^ANEM Z
LESNIM PRAHOM, NA NJEGOVE MEHANSKE LASTNOSTI IN
EROZIJSKO OBRABO
Arul Sujin Jose1, Ayyanar Athijayamani2, Kalimuthu Ramanathan3,
Susaiyappan Sidhardhan4
1Lourdes Mount College of Engineering and Technology, Department of Mechanical Engineering, Kanyakumari, Tamilnadu, India
2Government College of Engineering, Department of Mechanical Engineering, Bodinayakkanur, Tamilnadu, India
3Alagappa Chettiar College of Engineering and Technology, Department of Mechanical Engineering, Karaikudi, Tamilnadu, India
4Government College of Engineering, Department of Civil Engineering, Tirunelveli, Tamilnadu, India
athimania@gmail.com
Prejem rokopisa – received: 2016-09-23; sprejem za objavo – accepted for publication: 2017-01-22
doi:10.17222/mit.2016.284
Several attempts were made to investigate the effects of various process parameters on the mechanical properties and wear
behavior of synthetic and natural cellulosic fibers and also particle-reinforced polymer composites. However, very few studies
were carried out on the effects of various process parameters on the mechanical and wear behavior of phenol formaldehyde (PF)
composites reinforced with natural cellulosic fibers and particles. Therefore, in the present study, an attempt was made to
observe the effects of various process parameters on the mechanical and wear behavior of wood-dust (WD) and coir-pith (CP)
particle-reinforced resole-type PF composites. First, the mechanical properties of a WD/PF composite were studied based on the
content of CP particles. Then, the erosive-wear behavior of the WD/PF composite was studied with respect to five different
parameters such particle content, erodent size, impact velocity, impingement angle, and standoff distance. The erosive
experiments were carried out for five different parameters based on the Taguchi experimental design (L27). The results show that
the mechanical properties of the WD/PF composite increase with an addition of CP particles. The increment in the composite
modulus was higher than that of the composite strength. The erosive test results indicate that the erosion-wear rate is affected by
the particle content, impingement angle, erodent size and impact velocity. Brittle-erosion behavior was identified on the surface
of the composite with a heavy erosive wear occurring at a 60° impingement angle.
Keywords: biowaste particles, phenol formaldehyde, composites, mechanical properties, erosive-wear resistance, Taguchi
method
Izvedenih je bilo `e kar nekaj poizkusov v zvezi z u~inki razli~nih procesnih parametrov na mehanske lastnosti in obrabo
polimernih kompozitov oja~anih s sinteti~nimi in naravnimi celuloznimi vlakni in/ali delci. Toda zelo malo raziskav je bilo
izvedenih glede vpliva razli~nih procesnih parametrov na mehanske lastnosti in obrabo fenolformaldehidnih (angl. PF)
kompozitov, oja~anih z naravnimi celuloznimi vlakni in delci. Tako je v pri~ujo~em delu predstavljen vpliv razli~nih procesnih
parametrov na mehanske lastnosti in obrabo PF kompozitov, ki so bili oja~ani z delci lesnega prahu (angl. WD) in delci kokosa
(angl. CP). Te vrste kompozitov se uporabljajo za izdelavo podplatov ~evljev. Najprej so bile dolo~ene mehanske lastnosti
WD/PF kompozitov glede na vsebnost CP delcev. Sledili so preizkusi in analize erozijske obrabe WD/PF kompozitov glede na
vsebnost (koli~ino) delcev v kompozitu, velikost, hitrost in razdaljo u~inkovanja erozijskega sredsta ter njegov vpadni kot.
Preizkusi so temeljili na analizi s Taguchijevo metodo (L27) s petimi razli~nimi parametri. Rezultati so pokazali, da se mehanske
lastnosti WD/PF kompozitov izbolj{ujejo z dodajanjem CP delcev. Povi{anje modula kompozitov je bilo ve~je od pove~anja
trdnosti kompozita. Erozijski testi ka`ejo, da je hitrost erozijske obrabe posledica vseh procesnih parametrov, to je: vsebnosti
delcev, udarnega kota, hitrosti in velikosti delcev izbranega erozijskega sredstva. Najve~ja obraba zaradi erozije je bila dose`ena
(ugotovljena) pri 60 stopinjskem vpadnem kotu abrazijskega sredstva z nastalimi po{kodbami krhkega zna~aja.
Klju~ne besede: delci bioodpadkov, fenolformaldehid, kompoziti, mehanske lastnosti, odpornost proti erozijski obrabi, Taguchi
metoda
1 INTRODUCTION
Recently, polymer composites reinforced with syn-
thetic materials have been replaced with polymer com-
posites reinforced with bio-based natural materials. The
bio-based natural materials have many advantages over
the synthetic materials like renewability, biodegrada-
bility, abundant availability, low costs, etc.1–3 In India,
particularly in the South Indian region, the biowaste
materials like wood dust, coir pith, groundnut shell,
coconut shell, cashew nut shell, etc. are abundantly
available because in that region, coconut, groundnut and
cashew nut are cultivated in large amounts. A number of
timber and oil mills are also available in the Southern
region of Tamilnadu, India. Therefore, bioparticles are
thrown away after producing useful materials and
Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811 805
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:620.1:620.163 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)805(2017)
dumped on the land of the village nearest to these in-
dustries.
Many studies have already reported on the properties
of different bio-based natural-fiber-reinforced polymer
composites in different conditions.3–7 But, only few
reports are available on the properties of polymer com-
posites filled with biowaste particles.8–11 In this investi-
gation, an attempt was made to study the mechanical and
wear behavior of PF composites reinforced with
biowaste particles (WD and CP). Mechanical properties
of wood-dust-particle-reinforced PF composites were
evaluated based on the content of coir pith. The erosive-
wear behavior of the composites was studied using five
different parameters such as particle content, erodent
size, impingement angle, impact velocity and standoff
distance. The erosive experiments were conducted as per
the Taguchi experimental design. The parameters used
for erosive-wear tests were also analyzed using an
analysis of variance with the wear rate.
2 MATERIALS AND METHODOLOGY
2.1 Materials
Wood-dust particles were collected from the Kumar
Timber and Sawmill, Karaikudi, Tamilnadu, India.
Coir-pith particles were collected from the Coir Industry,
Sozhavanthan, Tamilnadu, India. From the collected
wood-dust and CP particles, microparticles with the
average size of 800 microns were separated using a
sieving machine available in our composite laboratory.
The resole-type PF liquid resin was procured, together
with a cross-linking agent (divinylbenzene) and acidic
catalyst (hydrochloric acid), from POOJA Chemicals,
Madurai, Tamilnadu, India.
2.2 Preparation of the composites
A hardboard mold box with dimensions of 150 mm ×
150 mm × 3 mm was used to prepare the wood-dust and
coir-pith-particle composite plates using the hand lay-up
technique. Wood dust/coir pith/phenol formaldehyde
composites were fabricated at three different concentra-
tions of wood-dust and coir-pith particles, i.e., (20, 30 and
40) % mass fractions. The amount of WD particles was
maintained at a fixed level of 20 % mass fraction. Three
different amounts of CP particles (0–20 % mass frac-
tions) were hybridized with the constant amount of WD
particles, i.e., 20WD/0CP, 20WD/10CP, 20WD/20CP.
The weight percentage of WD and CP particles and
designation of the composites are given in Table 1. Prior
to the process, the particles were dried in sunlight for
12 h. The PF resin with the particles was mixed with a
mechanical stirrer at room temperature for 30 min. Then,
the cross-linking agent and acidic catalyst were also
mixed into the mixture of phenol formaldehyde/particles
and once again stirred with the mechanical stirrer for 15
min. After that, the mixture was poured into the mold
box and allowed to cure at room temperature for 48 h.
2.3 Testing composite specimens
Composite specimens were characterized using me-
chanical tests such as tensile, flexural and impact tests.
The tensile tests were conducted on an FIE universal
testing machine (UTE 40 HGFL) in accordance with
ASTM D638-10.12 The flexural tests were performed on
the same testing machine in accordance with ASTM
D790-10.13 The impact tests were carried out on an Izod
impact machine according to ISO 180.14 All the tests
were conducted at room temperature and atmospheric
pressure.
2.4 Taguchi experimental design
The erosive behavior of the WD/CP/PF composite
was studied based on the Taguchi method and analysis of
variance techniques. Experiments were performed as per
Taguchi experimental design (an orthogonal array)
because it is a systematic and efficient approach to get
the optimum range of process parameters with a good
performance. The number of experiments can be reduced
due to the constructed orthogonal array, which provides a
set of well-balanced experiments.15 The results obtained
with this experimental design are transformed into
signal-to-noise (S/N) ratios, which serve as objective
functions for the optimization of parameters and help
with the result analysis. There are three S/N ratios
available for the optimization of several static problems:
the smaller-the-better (used to minimize the response),
the nominal-the-better (used whenever an ideal quality is
equated with a particular nominal value.) and the larger-
the-better ratio (used to maximize the response). Among
these three characteristics, the minimum erosion rate
comes under the smaller-the-better characteristic, which
can be expressed as Equation (1):
S/N = –10 Log10 (1)
(the mean of the sum of squares of the measured data)
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
806 Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Table 1: The weight percentage of WD and CP particles and designation of the composites
Total weight percentage
of particles in the
composites
Weight percentage of
resin
Weight percentage of
WD particles
Weight percentage of
CP particles
Designation of
composites
20 80 20 0 20WD/0CP
30 70 20 10 20WD/10CP
40 60 20 20 20WD/20CP
The five different process parameters at three levels
are used in this study to observe the erosive behavior of
the WD/CP/PF composite. Therefore, the actual number
of experiments, based on the traditional experimental
design, should be 243 (35). But, this number is reduced
to 27 experiments using the Taguchi technique. The pro-
cess parameters and their setting levels for the erosion
test of the WD/CP/PF composite are presented in Table
2. In these experiments, the following parameters are
fixed throughout the process: the type of erodent is
silica, the erodent feed rate is 10.0±1.0 g/min, the nozzle
length is 80 mm, and the nozzle diameter is 3 mm.
Table 2: The erosive process parameters with their designation and
setting levels
Process Parameters and their
designation Level I Level II Level III
Particle content: (A) wt% 20 30 40
Impact velocity: (B) m/sec 41 52 63
Impingement angle: (C) degree 30 60 90
Erodent size: (D) ìm 300 500 700
Stand-off distance: (E) mm 80 120 160
2.5 Erosion test
The erosive tests of the WD/CP/PF composite speci-
mens were conducted as shown in the schematic diagram
of the erosion process (Figure 1). The main components
of the erosion-test apparatus are the erodent feeder box,
erodent feeder nozzle, mixing chamber, nozzle of the
mixing chamber, air-flow vent, sample holder and ero-
dent collector. Dry silica sand with three different sizes
(300, 500 and 700) μm was used as the erodent in the
erosion tests. After the test, the composite samples were
taken from the apparatus and cleaned with acetone.
Then, the cleaned composite specimens were dried and
weighed using a precision digital balance at an accuracy
of ±0.1 mg. The composite samples were weighed before
and after the erosion tests and their difference is termed
as the weight loss. Then, the weight loss was recorded
and used for the erosion-rate calculation. Generally, the
erosion rate can be obtained as the ratio of the weight
loss of samples to the weight of the eroding particle. The
process was repeated until the steady-state erosion was
reached.
3 RESULTS AND DISCUSSION
3.1 Mechanical properties of the composites
Mechanical tests were carried out on the WD/CP/PF
composites and their results are presented in Figure 2a.
The neat-resin sample had a tensile strength of 29.8
MPa, tensile modulus of 1168.4 MPa, flexural strength
of 34.7 MPa, flexural modulus of 1257.4 MPa, and
impact strength of 1.24 KJ/m2. It can be seen that the
tensile strength and modulus of the PF composite in-
crease with an increase in the particle content. The
tensile strength of the 20WD/PF composite is almost the
same as that of the neat-resin sample. It shows that the
addition of WD particles enhances the strength of the PF
composite. The WD/PF composite without the addition
of CP particles has a tensile strength of 30.4 MPa and
this value increases to 41.7 MPa with the incorporation
of 10 % mass fraction of CP particles; after that, it de-
creases to 36.8 MPa with the addition of 20 % mass
fraction of CP particles. This may be due to a poor
interfacial bonding between the particles and the matrix,
i.e., a weak transfer of stress. Moreover, the stress con-
centration in the PF matrix may be created due to the
corner edges of the irregularly shaped WD and CP parti-
cles.
Due to the addition of 10 % mass fraction and 20 %
mass fraction of CP particles, the tensile strength of the
WD/PF composite increases by about 37.17 % and
21.1 %, respectively. Figure 2a also shows the tensile-
modulus values of the WD/CP/PF composites with res-
pect to the particle content. The composite also reached
the tensile-modulus value of the neat-resin sample with
the particle addition of 20 % mass fraction. The tensile-
modulus value of the WD/PF composite increased with
the further addition of CP particles. The maximum mo-
dulus value was observed at 40 % mass fraction of the
particles.
The results of the flexural tests of the WD/CP/PF
composites with respect to the particle content are given
in Figure 2b. It is interesting to note that the flexural
strength and modulus of the WD/PF composite increase
with the addition of CP particles. The flexural strength of
the WD/PF composite is slightly lower than the value of
the neat-resin sample. The maximum values of the flexu-
ral strength and modulus were identified at the 40 %
addition. The flexural strength of the WD/PF composite
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811 807
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Schematic diagram of the erosive process of the WD/CP/PF
composite
dumped on the land of the village nearest to these in-
dustries.
Many studies have already reported on the properties
of different bio-based natural-fiber-reinforced polymer
composites in different conditions.3–7 But, only few
reports are available on the properties of polymer com-
posites filled with biowaste particles.8–11 In this investi-
gation, an attempt was made to study the mechanical and
wear behavior of PF composites reinforced with
biowaste particles (WD and CP). Mechanical properties
of wood-dust-particle-reinforced PF composites were
evaluated based on the content of coir pith. The erosive-
wear behavior of the composites was studied using five
different parameters such as particle content, erodent
size, impingement angle, impact velocity and standoff
distance. The erosive experiments were conducted as per
the Taguchi experimental design. The parameters used
for erosive-wear tests were also analyzed using an
analysis of variance with the wear rate.
2 MATERIALS AND METHODOLOGY
2.1 Materials
Wood-dust particles were collected from the Kumar
Timber and Sawmill, Karaikudi, Tamilnadu, India.
Coir-pith particles were collected from the Coir Industry,
Sozhavanthan, Tamilnadu, India. From the collected
wood-dust and CP particles, microparticles with the
average size of 800 microns were separated using a
sieving machine available in our composite laboratory.
The resole-type PF liquid resin was procured, together
with a cross-linking agent (divinylbenzene) and acidic
catalyst (hydrochloric acid), from POOJA Chemicals,
Madurai, Tamilnadu, India.
2.2 Preparation of the composites
A hardboard mold box with dimensions of 150 mm ×
150 mm × 3 mm was used to prepare the wood-dust and
coir-pith-particle composite plates using the hand lay-up
technique. Wood dust/coir pith/phenol formaldehyde
composites were fabricated at three different concentra-
tions of wood-dust and coir-pith particles, i.e., (20, 30 and
40) % mass fractions. The amount of WD particles was
maintained at a fixed level of 20 % mass fraction. Three
different amounts of CP particles (0–20 % mass frac-
tions) were hybridized with the constant amount of WD
particles, i.e., 20WD/0CP, 20WD/10CP, 20WD/20CP.
The weight percentage of WD and CP particles and
designation of the composites are given in Table 1. Prior
to the process, the particles were dried in sunlight for
12 h. The PF resin with the particles was mixed with a
mechanical stirrer at room temperature for 30 min. Then,
the cross-linking agent and acidic catalyst were also
mixed into the mixture of phenol formaldehyde/particles
and once again stirred with the mechanical stirrer for 15
min. After that, the mixture was poured into the mold
box and allowed to cure at room temperature for 48 h.
2.3 Testing composite specimens
Composite specimens were characterized using me-
chanical tests such as tensile, flexural and impact tests.
The tensile tests were conducted on an FIE universal
testing machine (UTE 40 HGFL) in accordance with
ASTM D638-10.12 The flexural tests were performed on
the same testing machine in accordance with ASTM
D790-10.13 The impact tests were carried out on an Izod
impact machine according to ISO 180.14 All the tests
were conducted at room temperature and atmospheric
pressure.
2.4 Taguchi experimental design
The erosive behavior of the WD/CP/PF composite
was studied based on the Taguchi method and analysis of
variance techniques. Experiments were performed as per
Taguchi experimental design (an orthogonal array)
because it is a systematic and efficient approach to get
the optimum range of process parameters with a good
performance. The number of experiments can be reduced
due to the constructed orthogonal array, which provides a
set of well-balanced experiments.15 The results obtained
with this experimental design are transformed into
signal-to-noise (S/N) ratios, which serve as objective
functions for the optimization of parameters and help
with the result analysis. There are three S/N ratios
available for the optimization of several static problems:
the smaller-the-better (used to minimize the response),
the nominal-the-better (used whenever an ideal quality is
equated with a particular nominal value.) and the larger-
the-better ratio (used to maximize the response). Among
these three characteristics, the minimum erosion rate
comes under the smaller-the-better characteristic, which
can be expressed as Equation (1):
S/N = –10 Log10 (1)
(the mean of the sum of squares of the measured data)
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
806 Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Table 1: The weight percentage of WD and CP particles and designation of the composites
Total weight percentage
of particles in the
composites
Weight percentage of
resin
Weight percentage of
WD particles
Weight percentage of
CP particles
Designation of
composites
20 80 20 0 20WD/0CP
30 70 20 10 20WD/10CP
40 60 20 20 20WD/20CP
The five different process parameters at three levels
are used in this study to observe the erosive behavior of
the WD/CP/PF composite. Therefore, the actual number
of experiments, based on the traditional experimental
design, should be 243 (35). But, this number is reduced
to 27 experiments using the Taguchi technique. The pro-
cess parameters and their setting levels for the erosion
test of the WD/CP/PF composite are presented in Table
2. In these experiments, the following parameters are
fixed throughout the process: the type of erodent is
silica, the erodent feed rate is 10.0±1.0 g/min, the nozzle
length is 80 mm, and the nozzle diameter is 3 mm.
Table 2: The erosive process parameters with their designation and
setting levels
Process Parameters and their
designation Level I Level II Level III
Particle content: (A) wt% 20 30 40
Impact velocity: (B) m/sec 41 52 63
Impingement angle: (C) degree 30 60 90
Erodent size: (D) ìm 300 500 700
Stand-off distance: (E) mm 80 120 160
2.5 Erosion test
The erosive tests of the WD/CP/PF composite speci-
mens were conducted as shown in the schematic diagram
of the erosion process (Figure 1). The main components
of the erosion-test apparatus are the erodent feeder box,
erodent feeder nozzle, mixing chamber, nozzle of the
mixing chamber, air-flow vent, sample holder and ero-
dent collector. Dry silica sand with three different sizes
(300, 500 and 700) μm was used as the erodent in the
erosion tests. After the test, the composite samples were
taken from the apparatus and cleaned with acetone.
Then, the cleaned composite specimens were dried and
weighed using a precision digital balance at an accuracy
of ±0.1 mg. The composite samples were weighed before
and after the erosion tests and their difference is termed
as the weight loss. Then, the weight loss was recorded
and used for the erosion-rate calculation. Generally, the
erosion rate can be obtained as the ratio of the weight
loss of samples to the weight of the eroding particle. The
process was repeated until the steady-state erosion was
reached.
3 RESULTS AND DISCUSSION
3.1 Mechanical properties of the composites
Mechanical tests were carried out on the WD/CP/PF
composites and their results are presented in Figure 2a.
The neat-resin sample had a tensile strength of 29.8
MPa, tensile modulus of 1168.4 MPa, flexural strength
of 34.7 MPa, flexural modulus of 1257.4 MPa, and
impact strength of 1.24 KJ/m2. It can be seen that the
tensile strength and modulus of the PF composite in-
crease with an increase in the particle content. The
tensile strength of the 20WD/PF composite is almost the
same as that of the neat-resin sample. It shows that the
addition of WD particles enhances the strength of the PF
composite. The WD/PF composite without the addition
of CP particles has a tensile strength of 30.4 MPa and
this value increases to 41.7 MPa with the incorporation
of 10 % mass fraction of CP particles; after that, it de-
creases to 36.8 MPa with the addition of 20 % mass
fraction of CP particles. This may be due to a poor
interfacial bonding between the particles and the matrix,
i.e., a weak transfer of stress. Moreover, the stress con-
centration in the PF matrix may be created due to the
corner edges of the irregularly shaped WD and CP parti-
cles.
Due to the addition of 10 % mass fraction and 20 %
mass fraction of CP particles, the tensile strength of the
WD/PF composite increases by about 37.17 % and
21.1 %, respectively. Figure 2a also shows the tensile-
modulus values of the WD/CP/PF composites with res-
pect to the particle content. The composite also reached
the tensile-modulus value of the neat-resin sample with
the particle addition of 20 % mass fraction. The tensile-
modulus value of the WD/PF composite increased with
the further addition of CP particles. The maximum mo-
dulus value was observed at 40 % mass fraction of the
particles.
The results of the flexural tests of the WD/CP/PF
composites with respect to the particle content are given
in Figure 2b. It is interesting to note that the flexural
strength and modulus of the WD/PF composite increase
with the addition of CP particles. The flexural strength of
the WD/PF composite is slightly lower than the value of
the neat-resin sample. The maximum values of the flexu-
ral strength and modulus were identified at the 40 %
addition. The flexural strength of the WD/PF composite
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811 807
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Schematic diagram of the erosive process of the WD/CP/PF
composite
was increased by about 68.99 % due to the incorporation
of 20 % mass fraction of CP particles.
The impact-strength values of the WD/CP/PF com-
posites after the impact tests are presented in Figure 2c.
It can be seen that the impact strength of the WD/PF
composite is slightly lower than the value of the neat-
resin sample. It is also shown that the impact strength of
the WD/PF composite increases with the addition of 10
% mass fraction of CP particles, but it decreases with the
incorporation of 20 % mass fraction of CP particles. This
may be due to a poor adhesion between the particles and
the matrix. It may also be due to the stress concentration
of the resin matrix. It is also observed that the incor-
poration of 10 % mass fraction and 20 % mass fraction
of CP particles shows 15.57 % and 8.19 % higher impact
values compared to the WD/PF composite.
3.2 Steady-state erosion: effects of the impingement
angle
Generally, the erosive-wear behavior of polymer
composite materials can be categorized as brittle and
ductile. A ductile erosive situation is created in thermo-
plastic polymer composites, whereas the a brittle erosive
situation may be created in thermosetting polymer
composites. For the steady-state-erosion analysis of the
WD/CP/PF composites, an erosion test was carried out
based on eight different impingement angles (20, 30, 40,
50, 60, 70, 80 and 90)°, keeping all the other process
parameters constant (the initial level values).
The effects of the impingement angles on the erosion
rate of the WD/CP/PF composites are presented in Fig-
ure 3. From this figure, it can be observed that the
erosion rate is high at the impingement angle of 60° for
all the composite specimens, irrespective of the particle
content. However, a more brittle erosive behavior was
identified for the 40 % mass-fraction (20WD/20CP)
composite specimen. However, in the 20 % and 30 %
mass-fraction composite specimens, a semi-brittle
erosive behavior was identified. This may be due to the
addition of WD and CP particles to the PF composites.
When the higher amounts of particles are added to the
polymer material, it behaves as a typical brittle material.
Therefore, the brittleness of the composites with 20 %
mass fraction and 30 % mass fraction of particles is
lower than the composite with 40 % mass fraction of
particles. Due to this, a semi-brittle erosive situation
exists during the erosive process. It is also clear from
Figure 3 that the erosion rate increased with the increase
in the particle content. This may be due to the increased
hardness of the PF composite material caused by the
addition of WD and CP particles.
3.3 Analysis of the erosion rate
The erosion rates for 27 combinations of the erosive
experiments conducted on the WD/CP/PF composites
are given in Table 3. The erosion analysis was made
with popular software, namely, MINITAB 17. From
Table 3, it can be concluded that the parameter combina-
tion of particle loading-A (level II = 30 % mass fraction),
impact velocity-B (level I = 41 m/s), impingement
angle-C (level I = 30°), erodent size-D (level II = 500
um) and standoff distance-E (160 mm) gives the mini-
mum erosion rate (189. 8 mg/kg). Moreover, another
parameter combination (experiment number 4) allows
the next level of the minimum erosion rate (194.1
mg/kg). The difference between these two erosion rates
is small, as seen from Table 3. Anyway, the first para-
meter combination mentioned above is recognized as the
better combination of parameters to obtain the minimum
erosion rate. The overall mean of the signal-to-noise
ratio for the erosion rate is found to be 49.75 dB.
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
808 Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Steady-state erosive behavior of WD/CP/PF composites for
eight different impingement angles
Figure 2: Variation of: a) tensile property, b) flexural property, and c)
impact strength based on the particle content
The effects of five erosive-process parameters on the
erosion rate are graphically presented in Figure 4a.
From this figure, it can be clearly concluded that para-
meter A (the particle content), parameter B (the impact
velocity) and parameter C (the impingement angle) are
the most significant parameters. Parameter E (the stan-
doff distance) shows a moderately significant influence,
while parameter D (the erodent size) has a relatively less
significant influence. Figure 4b shows the interaction
between the erosive parameters. From this figure, it is
observed that a moderate interaction exists between para-
meters A and B, and between A and C. The interaction
between parameters B and C is below the moderate level.
Figures 5a to 5c show 3D surface plots of the ero-
sion rate with significant process parameters. The
observation is similar to the one made of the interaction
plots of the erosion rate. From the erosion test analysis
of the WD/CP/PF composites, it can be concluded that
the erodent size is most insignificant for the erosion rate.
The standoff distance shows relatively less significance
when compared to the other three process parameters
(particle content, impact velocity and impingement
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811 809
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: 3D surface plots of erosion rate vs process parameters:
a) A × B, b) A × C, and c) B × C
Figure 4: a) Effects of erosive-process parameters and b) effects of
interactions of erosive-process parameters on the erosion rate
Table 3: The erosion rates and their S/N ratio of WD/CP/PF compo-
sites for 27 combinations
Experi
ment
No.
A B C D E
Erosion
rate
mg/kg
S/N
ratio
dB
1 20 41 30 300 80 209.5 -46.42
2 20 41 60 500 120 273.8 -48.75
3 20 41 90 700 160 220.7 -46.87
4 20 52 30 500 120 194.1 -45.76
5 20 52 60 700 160 267.3 -48.54
6 20 52 90 300 80 246.2 -47.82
7 20 63 30 700 160 211.9 -46.52
8 20 63 60 300 80 301.3 -49.58
9 20 63 90 500 120 277.1 -48.85
10 30 41 30 500 160 189.8 -45.65
11 30 41 60 700 80 343.5 -50.72
12 30 41 90 300 120 312.7 -49.90
13 30 52 30 700 80 298.9 -49.51
14 30 52 60 300 120 367.2 -51.29
15 30 52 90 500 160 351.3 -50.91
16 30 63 30 300 120 300.8 -49.56
17 30 63 60 500 160 378.5 -51.56
18 30 63 90 700 80 361.9 -51.17
19 40 41 30 700 120 332.8 -50.44
20 40 41 60 300 160 387.5 -51.76
21 40 41 90 500 80 370.6 -51.37
22 40 52 30 300 160 359.1 -51.10
23 40 52 60 500 80 398.3 -52.00
24 40 52 90 700 120 381.7 -51.63
25 40 63 30 500 80 379.2 -51.58
26 40 63 60 700 120 427.6 -52.62
27 40 63 90 300 160 369.8 -51.36
was increased by about 68.99 % due to the incorporation
of 20 % mass fraction of CP particles.
The impact-strength values of the WD/CP/PF com-
posites after the impact tests are presented in Figure 2c.
It can be seen that the impact strength of the WD/PF
composite is slightly lower than the value of the neat-
resin sample. It is also shown that the impact strength of
the WD/PF composite increases with the addition of 10
% mass fraction of CP particles, but it decreases with the
incorporation of 20 % mass fraction of CP particles. This
may be due to a poor adhesion between the particles and
the matrix. It may also be due to the stress concentration
of the resin matrix. It is also observed that the incor-
poration of 10 % mass fraction and 20 % mass fraction
of CP particles shows 15.57 % and 8.19 % higher impact
values compared to the WD/PF composite.
3.2 Steady-state erosion: effects of the impingement
angle
Generally, the erosive-wear behavior of polymer
composite materials can be categorized as brittle and
ductile. A ductile erosive situation is created in thermo-
plastic polymer composites, whereas the a brittle erosive
situation may be created in thermosetting polymer
composites. For the steady-state-erosion analysis of the
WD/CP/PF composites, an erosion test was carried out
based on eight different impingement angles (20, 30, 40,
50, 60, 70, 80 and 90)°, keeping all the other process
parameters constant (the initial level values).
The effects of the impingement angles on the erosion
rate of the WD/CP/PF composites are presented in Fig-
ure 3. From this figure, it can be observed that the
erosion rate is high at the impingement angle of 60° for
all the composite specimens, irrespective of the particle
content. However, a more brittle erosive behavior was
identified for the 40 % mass-fraction (20WD/20CP)
composite specimen. However, in the 20 % and 30 %
mass-fraction composite specimens, a semi-brittle
erosive behavior was identified. This may be due to the
addition of WD and CP particles to the PF composites.
When the higher amounts of particles are added to the
polymer material, it behaves as a typical brittle material.
Therefore, the brittleness of the composites with 20 %
mass fraction and 30 % mass fraction of particles is
lower than the composite with 40 % mass fraction of
particles. Due to this, a semi-brittle erosive situation
exists during the erosive process. It is also clear from
Figure 3 that the erosion rate increased with the increase
in the particle content. This may be due to the increased
hardness of the PF composite material caused by the
addition of WD and CP particles.
3.3 Analysis of the erosion rate
The erosion rates for 27 combinations of the erosive
experiments conducted on the WD/CP/PF composites
are given in Table 3. The erosion analysis was made
with popular software, namely, MINITAB 17. From
Table 3, it can be concluded that the parameter combina-
tion of particle loading-A (level II = 30 % mass fraction),
impact velocity-B (level I = 41 m/s), impingement
angle-C (level I = 30°), erodent size-D (level II = 500
um) and standoff distance-E (160 mm) gives the mini-
mum erosion rate (189. 8 mg/kg). Moreover, another
parameter combination (experiment number 4) allows
the next level of the minimum erosion rate (194.1
mg/kg). The difference between these two erosion rates
is small, as seen from Table 3. Anyway, the first para-
meter combination mentioned above is recognized as the
better combination of parameters to obtain the minimum
erosion rate. The overall mean of the signal-to-noise
ratio for the erosion rate is found to be 49.75 dB.
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
808 Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Steady-state erosive behavior of WD/CP/PF composites for
eight different impingement angles
Figure 2: Variation of: a) tensile property, b) flexural property, and c)
impact strength based on the particle content
The effects of five erosive-process parameters on the
erosion rate are graphically presented in Figure 4a.
From this figure, it can be clearly concluded that para-
meter A (the particle content), parameter B (the impact
velocity) and parameter C (the impingement angle) are
the most significant parameters. Parameter E (the stan-
doff distance) shows a moderately significant influence,
while parameter D (the erodent size) has a relatively less
significant influence. Figure 4b shows the interaction
between the erosive parameters. From this figure, it is
observed that a moderate interaction exists between para-
meters A and B, and between A and C. The interaction
between parameters B and C is below the moderate level.
Figures 5a to 5c show 3D surface plots of the ero-
sion rate with significant process parameters. The
observation is similar to the one made of the interaction
plots of the erosion rate. From the erosion test analysis
of the WD/CP/PF composites, it can be concluded that
the erodent size is most insignificant for the erosion rate.
The standoff distance shows relatively less significance
when compared to the other three process parameters
(particle content, impact velocity and impingement
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811 809
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: 3D surface plots of erosion rate vs process parameters:
a) A × B, b) A × C, and c) B × C
Figure 4: a) Effects of erosive-process parameters and b) effects of
interactions of erosive-process parameters on the erosion rate
Table 3: The erosion rates and their S/N ratio of WD/CP/PF compo-
sites for 27 combinations
Experi
ment
No.
A B C D E
Erosion
rate
mg/kg
S/N
ratio
dB
1 20 41 30 300 80 209.5 -46.42
2 20 41 60 500 120 273.8 -48.75
3 20 41 90 700 160 220.7 -46.87
4 20 52 30 500 120 194.1 -45.76
5 20 52 60 700 160 267.3 -48.54
6 20 52 90 300 80 246.2 -47.82
7 20 63 30 700 160 211.9 -46.52
8 20 63 60 300 80 301.3 -49.58
9 20 63 90 500 120 277.1 -48.85
10 30 41 30 500 160 189.8 -45.65
11 30 41 60 700 80 343.5 -50.72
12 30 41 90 300 120 312.7 -49.90
13 30 52 30 700 80 298.9 -49.51
14 30 52 60 300 120 367.2 -51.29
15 30 52 90 500 160 351.3 -50.91
16 30 63 30 300 120 300.8 -49.56
17 30 63 60 500 160 378.5 -51.56
18 30 63 90 700 80 361.9 -51.17
19 40 41 30 700 120 332.8 -50.44
20 40 41 60 300 160 387.5 -51.76
21 40 41 90 500 80 370.6 -51.37
22 40 52 30 300 160 359.1 -51.10
23 40 52 60 500 80 398.3 -52.00
24 40 52 90 700 120 381.7 -51.63
25 40 63 30 500 80 379.2 -51.58
26 40 63 60 700 120 427.6 -52.62
27 40 63 90 300 160 369.8 -51.36
angle). It can be concluded that the combination of the
parameters – the particle content at level II, the impact
velocity at level I and the impingement angle at level I –
gives the minimum erosion rate. Therefore, this com-
bination is recognized as the best combination of the
erosive-process parameters to get the minimum erosion
rate within the selected parameter range.
3.4 Analysis of the variance for the erosion rate
Taguchi’s analysis of variance can be used to find the
set of significant parameters as well as their interactions
in any system. In this study, ANOVA is used to under-
stand the contribution of the parameters to the erosion
rate and the effects of their interactions on the erosion
rate of the WD/CP/PF composite. The ANOVA results
for the erosion rate of the WD/CP/PF composite are
given in Table 4. The last column (p-value) in this table
indicates the highly significant parameters and their
main effects depend upon the value of p. The p values
for the particle content, impact velocity, impingement
angle and standoff distance are 0.000, 0.023, 0.003 and
0.186, indicating their great influence on the erosion rate
of the composite. The p values of the interactions of the
process parameters show a moderate significance and a
significance below the moderate level. Table 5 shows the
responses for the S/N ratio (the small-the-better charac-
teristic). The order of the erosive-process parameters
based on their contributions to obtain the minimum ero-
sion rate is the particle content, impingement angle,
impact velocity, standoff distance and erodent size.
Table 4: Analysis of variance for Means of erosion rate
Source DF Seq SS Adj SS Adj MS F P
A 2 81404 81403.9 40701.9 120.99 0.000
B 2 7520 7519.9 3759.9 11.18 0.023
C 2 25188 25188.1 12594.0 37.44 0.003
D 2 97 96.7 48.4 0.14 0.870
E 2 1778 1778.4 889.2 2.64 0.186
A*B 4 2698 2698.2 674.6 2.01 0.258
A*C 4 3305 3305.3 826.3 2.46 0.203
B*C 4 791 791.1 197.8 0.59 0.690
Residual
Error 4 1346 1345.6 336.4
Total 26 124127
Table 5: Response table for signal to noise ratios (smaller is better)
Level A B C D E
1 -47.68 -49.10 -48.51 -49.87 -50.02
2 -50.03 -49.84 -50.76 -49.61 -49.87
3 -51.54 -50.31 -49.99 -49.78 -49.37
Delta 3.86 1.21 2.25 0.26 0.65
Rank 1 3 2 5 4
3.5 Confirmation experiment
At the end of this study, we carried out a confirma-
tory experiment as the final test to validate the estimated
results obtained during the erosive analysis of the
WD/CP/PF composite. Therefore, the experimental
results were verified with the estimated results using the
confirmation test. This test was conducted to predict the
erosion rate caused by a new set of erosive-process-
parameter levels (A2B1C1E3). To predict the S/N ratio,
the following equation can be used:
S/Np = (A2-T) + (B1-T) + (C1-T) + (E3-T) + T (2)
where T is the overall experimental average of the S/N
ratio and S/Np is the value of the predicted S/N ratio.
The comparison results for the predicted and experimen-
tal S/N ratio of the optimum process parameters are
given in Table 6. The difference between the predicted
and experimental S/N ratio is 0.51, i.e., an error of
1.1 %. It proves that the model can predict the erosion
rate with a reasonable accuracy.
Table 6: Comparison results of predicted and experimental signal-
to-noise ratio of optimal process parameters
Optimal erosive process parameters
Predicted Experimental Difference
Parameter
level A2B1C1E3 A2B1C1E3
Predicted-
Experimental
Erosion rate
mg/kg 191.5 189.8 1.7
Signal-to-noise
ratio dB -46.17 -45.66
4 CONCLUSIONS
Mechanical properties of a WD/CP/PF composite
were analyzed based on the particle content. The results
show that the tensile strength of the WD/PF composite
increased with the addition of 10 % mass fraction of CP
particles, but decreased with the addition of 20 % mass
fraction of CP particles. The flexural properties of the
WD/PF composite increased with the increase in CP
particles. The impact strength of the WD/PF composite
also increased with the addition of 10 % mass fraction of
CP particles and decreased with further addition of CP
particles. The steady-state erosion analysis was carried
out for eight different impingement angles on the
WD/CP/PF composite. The composite with the lower
particle content shows a semi-brittle erosive behavior
with a higher erosion wear at the 60° impingement angle.
On the other hand, the composite with the higher particle
content showed a fully brittle nature of the erosive beha-
vior with a higher erosion wear at the 60° impingement
angle. From the erosive analysis of the WD/CP/PF com-
posite, the process parameters like the particle content,
impingement angle and impact velocity are found to be
the most significant parameters influencing the erosion
rate. The standoff distance shows a moderate influence
on the erosion rate, while the erodent size shows a less
significant influence on the erosion rate.
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
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MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
5 REFERENCES
1 G. K. Mani, J. B. B. Rayappan, D. K. Bisoyi, Synthesis and cha-
racterization of kapok fibers and its composites, J. of Appl. Sci., 12
(2012), 1661–1665, doi:10.3923/jas.2012.1661.1665
2 J. A. Khan, M. A. Khan, R. Islam, Mechanical, thermal and degra-
dation properties of jute fabric – reinforced polypropylene
composites: Effect of potassium permanganate as oxidizing agent,
Polym. Compos., 34 (2013), 671–680, doi:10.1002/pc.22470
3 S. Singh, D. Deepak, L. Aggarwal, V. K. Gupta, Tensile and flexural
behavior of hemp fiber reinforced virgin recycled HDPE matrix com-
posites, Procedia Mater. Sci., 6 (2014), 1696–1702, doi:10.1016/
j.mspro.2014.07.155
4 I. V. Surendra, K. V. Rao, K. V. P. P. Chandu, Fabrication and investi-
gation of mechanical properties of sisal, jute & okra natural fiber
reinforced hybrid polymer composites, Int. J. Eng. Trends Technol.,
19 (2015), 116–120, doi: 10.14445/22315381/IJETT-V19P220
5 D. Kurniawan, B. S. Kim, H. Y. Lee, J. Y. Lim, Effects of repetitive
processing, wood content, and coupling agent on the mechanical,
thermal, and water absorption properties of wood/polypropylene
green composites, J. Adhes. Sci. Technol., 27 (2013), 1301–1312,
doi:10.1080/01694243.2012.695948
6 E. Muñoz, J. A. García-Manrique, Water absorption behaviour and
its effect on the mechanical properties of flax fibre reinforced
bioepoxy composites, Int. J. Polym. Sci., (2015), 1–10, doi:10.1155/
2015/390275
7 M. G. A. Selvan, A. Athijayamani, Mechanical properties of fragrant
screwpine fiber reinforced unsaturated polyester composite: Effect of
fiber length, fiber treatment and water absorption, Fibers Polym., 17
(2016), 104–116, doi:10.1007/s12221-016-5593-x
8 S. I. Durowaye, G. I. Lawal, M. A. Akande, V. O. Durowaye, Me-
chanical properties of particulate coconut shell and palm fruit poly-
ester composites, Int. J. Mater. Eng., 4 (2014), 141–147,
doi:10.5923/j.ijme.20140404.04
9 L. Netra, S. Thomas, C. K. Das, R. Adhikari, Analysis of morpho-
logical and mechanical behaviours of bamboo flour reinforced
polypropylene composites, Nepal J. Sci. Technol., 13 (2012),
95–100, doi:10.3126/njst.v13i1.7447
10 M. A. M. M. Idrus, S. Hamdan, M. R. Rahman, M. S. Islam, Treated
tropical wood sawdust-polypropylene polymer composite: mecha-
nical and morphological study, J. Biomater. Nano-Biotechnol., 2
(2011), 435–444, doi:10.4236/jbnb.2011.24053
11 R. Panneerdhass, A. Gnanavelbabu, K. Rajkumar, Mechanical pro-
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j.proeng.2014.12.447
12 ASTM D 638-10, Standard test method for tensile properties of
plastics, Annual Book of ASTM Standards, 08.01 (2010), 1–16,
ASTM International, West Conshohocken
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un-reinforced and reinforced plastics and electrical insulating
materials, Annual Book of ASTM Standards, 08.01 (2010), 1–11,
ASTM International, West Conshohocken
14 ISO 180:2000, Plastics – determination of Izod impact strength, third
edition, ISO Central Secretariat, Switzerland, 2000
15 U. S. Rao, L. L. R. Rodrigues, Influence of machining parameters on
tool wear in drilling of GFRP composites –Taguchi analysis and
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doi:10.15224/978-1-63248-002-6-84
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811 811
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
angle). It can be concluded that the combination of the
parameters – the particle content at level II, the impact
velocity at level I and the impingement angle at level I –
gives the minimum erosion rate. Therefore, this com-
bination is recognized as the best combination of the
erosive-process parameters to get the minimum erosion
rate within the selected parameter range.
3.4 Analysis of the variance for the erosion rate
Taguchi’s analysis of variance can be used to find the
set of significant parameters as well as their interactions
in any system. In this study, ANOVA is used to under-
stand the contribution of the parameters to the erosion
rate and the effects of their interactions on the erosion
rate of the WD/CP/PF composite. The ANOVA results
for the erosion rate of the WD/CP/PF composite are
given in Table 4. The last column (p-value) in this table
indicates the highly significant parameters and their
main effects depend upon the value of p. The p values
for the particle content, impact velocity, impingement
angle and standoff distance are 0.000, 0.023, 0.003 and
0.186, indicating their great influence on the erosion rate
of the composite. The p values of the interactions of the
process parameters show a moderate significance and a
significance below the moderate level. Table 5 shows the
responses for the S/N ratio (the small-the-better charac-
teristic). The order of the erosive-process parameters
based on their contributions to obtain the minimum ero-
sion rate is the particle content, impingement angle,
impact velocity, standoff distance and erodent size.
Table 4: Analysis of variance for Means of erosion rate
Source DF Seq SS Adj SS Adj MS F P
A 2 81404 81403.9 40701.9 120.99 0.000
B 2 7520 7519.9 3759.9 11.18 0.023
C 2 25188 25188.1 12594.0 37.44 0.003
D 2 97 96.7 48.4 0.14 0.870
E 2 1778 1778.4 889.2 2.64 0.186
A*B 4 2698 2698.2 674.6 2.01 0.258
A*C 4 3305 3305.3 826.3 2.46 0.203
B*C 4 791 791.1 197.8 0.59 0.690
Residual
Error 4 1346 1345.6 336.4
Total 26 124127
Table 5: Response table for signal to noise ratios (smaller is better)
Level A B C D E
1 -47.68 -49.10 -48.51 -49.87 -50.02
2 -50.03 -49.84 -50.76 -49.61 -49.87
3 -51.54 -50.31 -49.99 -49.78 -49.37
Delta 3.86 1.21 2.25 0.26 0.65
Rank 1 3 2 5 4
3.5 Confirmation experiment
At the end of this study, we carried out a confirma-
tory experiment as the final test to validate the estimated
results obtained during the erosive analysis of the
WD/CP/PF composite. Therefore, the experimental
results were verified with the estimated results using the
confirmation test. This test was conducted to predict the
erosion rate caused by a new set of erosive-process-
parameter levels (A2B1C1E3). To predict the S/N ratio,
the following equation can be used:
S/Np = (A2-T) + (B1-T) + (C1-T) + (E3-T) + T (2)
where T is the overall experimental average of the S/N
ratio and S/Np is the value of the predicted S/N ratio.
The comparison results for the predicted and experimen-
tal S/N ratio of the optimum process parameters are
given in Table 6. The difference between the predicted
and experimental S/N ratio is 0.51, i.e., an error of
1.1 %. It proves that the model can predict the erosion
rate with a reasonable accuracy.
Table 6: Comparison results of predicted and experimental signal-
to-noise ratio of optimal process parameters
Optimal erosive process parameters
Predicted Experimental Difference
Parameter
level A2B1C1E3 A2B1C1E3
Predicted-
Experimental
Erosion rate
mg/kg 191.5 189.8 1.7
Signal-to-noise
ratio dB -46.17 -45.66
4 CONCLUSIONS
Mechanical properties of a WD/CP/PF composite
were analyzed based on the particle content. The results
show that the tensile strength of the WD/PF composite
increased with the addition of 10 % mass fraction of CP
particles, but decreased with the addition of 20 % mass
fraction of CP particles. The flexural properties of the
WD/PF composite increased with the increase in CP
particles. The impact strength of the WD/PF composite
also increased with the addition of 10 % mass fraction of
CP particles and decreased with further addition of CP
particles. The steady-state erosion analysis was carried
out for eight different impingement angles on the
WD/CP/PF composite. The composite with the lower
particle content shows a semi-brittle erosive behavior
with a higher erosion wear at the 60° impingement angle.
On the other hand, the composite with the higher particle
content showed a fully brittle nature of the erosive beha-
vior with a higher erosion wear at the 60° impingement
angle. From the erosive analysis of the WD/CP/PF com-
posite, the process parameters like the particle content,
impingement angle and impact velocity are found to be
the most significant parameters influencing the erosion
rate. The standoff distance shows a moderate influence
on the erosion rate, while the erodent size shows a less
significant influence on the erosion rate.
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
810 Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
5 REFERENCES
1 G. K. Mani, J. B. B. Rayappan, D. K. Bisoyi, Synthesis and cha-
racterization of kapok fibers and its composites, J. of Appl. Sci., 12
(2012), 1661–1665, doi:10.3923/jas.2012.1661.1665
2 J. A. Khan, M. A. Khan, R. Islam, Mechanical, thermal and degra-
dation properties of jute fabric – reinforced polypropylene
composites: Effect of potassium permanganate as oxidizing agent,
Polym. Compos., 34 (2013), 671–680, doi:10.1002/pc.22470
3 S. Singh, D. Deepak, L. Aggarwal, V. K. Gupta, Tensile and flexural
behavior of hemp fiber reinforced virgin recycled HDPE matrix com-
posites, Procedia Mater. Sci., 6 (2014), 1696–1702, doi:10.1016/
j.mspro.2014.07.155
4 I. V. Surendra, K. V. Rao, K. V. P. P. Chandu, Fabrication and investi-
gation of mechanical properties of sisal, jute & okra natural fiber
reinforced hybrid polymer composites, Int. J. Eng. Trends Technol.,
19 (2015), 116–120, doi: 10.14445/22315381/IJETT-V19P220
5 D. Kurniawan, B. S. Kim, H. Y. Lee, J. Y. Lim, Effects of repetitive
processing, wood content, and coupling agent on the mechanical,
thermal, and water absorption properties of wood/polypropylene
green composites, J. Adhes. Sci. Technol., 27 (2013), 1301–1312,
doi:10.1080/01694243.2012.695948
6 E. Muñoz, J. A. García-Manrique, Water absorption behaviour and
its effect on the mechanical properties of flax fibre reinforced
bioepoxy composites, Int. J. Polym. Sci., (2015), 1–10, doi:10.1155/
2015/390275
7 M. G. A. Selvan, A. Athijayamani, Mechanical properties of fragrant
screwpine fiber reinforced unsaturated polyester composite: Effect of
fiber length, fiber treatment and water absorption, Fibers Polym., 17
(2016), 104–116, doi:10.1007/s12221-016-5593-x
8 S. I. Durowaye, G. I. Lawal, M. A. Akande, V. O. Durowaye, Me-
chanical properties of particulate coconut shell and palm fruit poly-
ester composites, Int. J. Mater. Eng., 4 (2014), 141–147,
doi:10.5923/j.ijme.20140404.04
9 L. Netra, S. Thomas, C. K. Das, R. Adhikari, Analysis of morpho-
logical and mechanical behaviours of bamboo flour reinforced
polypropylene composites, Nepal J. Sci. Technol., 13 (2012),
95–100, doi:10.3126/njst.v13i1.7447
10 M. A. M. M. Idrus, S. Hamdan, M. R. Rahman, M. S. Islam, Treated
tropical wood sawdust-polypropylene polymer composite: mecha-
nical and morphological study, J. Biomater. Nano-Biotechnol., 2
(2011), 435–444, doi:10.4236/jbnb.2011.24053
11 R. Panneerdhass, A. Gnanavelbabu, K. Rajkumar, Mechanical pro-
perties of luffa fiber and ground nut reinforced epoxy polymer hybrid
composites, Procedia Eng., 97 (2014), 2042–2051, doi:10.1016/
j.proeng.2014.12.447
12 ASTM D 638-10, Standard test method for tensile properties of
plastics, Annual Book of ASTM Standards, 08.01 (2010), 1–16,
ASTM International, West Conshohocken
13 ASTM D 790–10, Standard test methods for flexural properties of
un-reinforced and reinforced plastics and electrical insulating
materials, Annual Book of ASTM Standards, 08.01 (2010), 1–11,
ASTM International, West Conshohocken
14 ISO 180:2000, Plastics – determination of Izod impact strength, third
edition, ISO Central Secretariat, Switzerland, 2000
15 U. S. Rao, L. L. R. Rodrigues, Influence of machining parameters on
tool wear in drilling of GFRP composites –Taguchi analysis and
ANOVA methodology, Proc. of the Inter. Conf. on Advances in
Mechanical and Robotics Engineering – MRE 2014, 25–29,
doi:10.15224/978-1-63248-002-6-84
A. SUJIN JOSE: EFFECTS OF AN ADDITION OF COIR-PITH PARTICLES ON THE MECHANICAL AND ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 805–811 811
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
813–822
OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF
GALVANIZED STEEL USING ARTIFICIAL NEURAL NETWORK
COUPLED WITH GENETIC ALGORITHM
OPTIMALNA DOL@INA PODPORE ([ABLONE, VODILA) PRI
TERMI^NEM VRTANJU GALVANIZIRANEGA JEKLA Z UPORABO
UMETNE NEVRONSKE MRE@E IN GENETSKEGA ALGORITMA
Navasingh Rajesh Jesudoss Hynes1, Ramar Kumar1, Jebaraj Angela Jennifa Sujana2
1Mepco Schlenk Engineering College, Department of Mechanical Engineering, Sivakasi, Vindhunagan, 626005 Tamil Nadu, India
2Mepco Schlenk Engineering College, Department of Information Technology, Sivakasi, Vindhunagan, 626005 Tamil Nadu, India
findhynes@yahoo.co.in, mepcokumar@gmail.com, ang_jenefa@mepcoeng.ac.in
Prejem rokopisa – received: 2016-09-27; sprejem za objavo – accepted for publication: 2017-01-24
doi:10.17222/mit.2016.290
Thermal drilling is a novel sheet-metal-hole-making technique that utilizes the heat produced at the interface of the rotating
conical tool and workpiece in order to soften the workpiece and pierce a hole into it. In this work, experiments with thermal
drilling of galvanized steel were conducted based on the Taguchi L27 orthogonal array. Significant process parameters such as
rotational speed, tool angle and workpiece thickness were varied during the experimentation. In thermal drilling, the
thermal-drill tool pushes aside a large amount of workpiece material to form a sleeve, which is often referred to as the bushing
length. A predictive model for the bushing length was developed using a feed-forward artificial neural network based on experi-
mental data. As the bushing length is closely associated with the tapping process, the influences of the input process parameters
play a vital role in fastening galvanized steel with threaded fasteners in diverse engineering applications. The optimization
problem was solved by implementing a genetic algorithm under constraint limits to maximize the bushing length. Further, a
confirmation test was conducted with the intention to compare the optimum value and its corresponding bushing length
predicted by the genetic algorithm. Good agreement was observed between the predicted and the experimental values.
Keywords: thermal drilling, artificial neural network, genetic algorithm, galvanized steel, bushing length
Termi~no vrtanje je nova, za vrtanje lukenj v plo~evino, uporabljena tehnika, ki izkori{~a toploto, proizvedeno na povr{ini
vrte~e se konice orodja na obdelovancu z namenom, da ga zmeh~a in vanj naredi luknjo. V delu so bili izvedeni preizkusi na
osnovi metode Taguchi L27 z ortogonalno matriko s termi~nim vrtanjem galvaniziranega jekla. Pomembni parametri postopka,
kot so: hitrost vrtenja, kot orodja in debelina obdelovanca, so se med eksperimentiranjem spreminjali. Pri toplotnem vrtanju,
vrtanje orodja potisne stran ve~ materiala obdelovanca tako, da se tvori navarek (rokav) okoli luknje, ki se pogosto omenja kot
dol`ina {ablone. Napovedni model za dol`ino {ablone, je razvit z uporabo umetne nevronske mre`e, ki temelji na znanstvenih
podatkih. Ker je dol`ina {ablone precej povezana s procesom izdelave navoja, vplivi teh vhodnih procesnih parametrov igrajo
klju~no vlogo pri pritrditvi galvaniziranega jekla z navojem pritrdilnih elementov v razli~nih in`enirskih aplikacijah. Problem
optimizacije je bil re{en z implementacijo genetskega algoritma na podlagi omejitev za pove~anje dol`ine {ablone. Ugotovljeno
je bilo dobro ujemanje med napovedano in eksperimentalno vrednostjo.
Klju~ne besede: termi~no vrtanje, umetna nevronska mre`a, genetski algoritem, galvanizirano jeklo, dol`ina {ablone
1 INTRODUCTION
Thermal drilling is an emerging hole-making process
with significant breakthroughs in many drilling situ-
ations in both automobile and aerospace applications.
This process uses the heat energy produced at the inter-
face in order to soften and then pierce the workpiece.1
Moreover, this heat energy enhances the flow ability of
the workpiece material, which is extruded onto both the
front and back sides of a drilled hole. Finally, the
extruded or deformed material forms a bushing shape,
which surrounds the drilled hole.2 The length of the
bushing formed on the workpiece after a thermal-drilling
operation is called the bushing length. Achieving a
sufficient bushing length is very important since the
bushing length can increase the threading depth and the
clamp-load-bearing capability in various engineering
applications.
Contrary to the conventional drilling process, in
thermal drilling, there is no chip and wastage of the ma-
terial, since all the deformed material contributes to
developing a bushing. As it abolishes chip generation, it
can be seen as a clean, eco-friendly and chipless hole-
making technique. When joining metal sheets or thin-
walled structures by drilling holes using the conven-
tional-drilling process, just one or two threads can be
made in them; however, reliability is a matter of great
concern in such a situation. Alternatively, weld nuts and
threaded rivet nuts are used, but due to thermal
distortion, they get jammed and twisted during the
fastening of the structures.3 On the other hand, during the
thermal-drilling process, the bushing formed is about 3
to 4 times thicker than the workpiece. The bushing thus
allows a greater contact area and high structural rigidity
of fasteners in joining situations.4
Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822 813
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 519.6:620.1:67.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)813(2017)
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
813–822
OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF
GALVANIZED STEEL USING ARTIFICIAL NEURAL NETWORK
COUPLED WITH GENETIC ALGORITHM
OPTIMALNA DOL@INA PODPORE ([ABLONE, VODILA) PRI
TERMI^NEM VRTANJU GALVANIZIRANEGA JEKLA Z UPORABO
UMETNE NEVRONSKE MRE@E IN GENETSKEGA ALGORITMA
Navasingh Rajesh Jesudoss Hynes1, Ramar Kumar1, Jebaraj Angela Jennifa Sujana2
1Mepco Schlenk Engineering College, Department of Mechanical Engineering, Sivakasi, Vindhunagan, 626005 Tamil Nadu, India
2Mepco Schlenk Engineering College, Department of Information Technology, Sivakasi, Vindhunagan, 626005 Tamil Nadu, India
findhynes@yahoo.co.in, mepcokumar@gmail.com, ang_jenefa@mepcoeng.ac.in
Prejem rokopisa – received: 2016-09-27; sprejem za objavo – accepted for publication: 2017-01-24
doi:10.17222/mit.2016.290
Thermal drilling is a novel sheet-metal-hole-making technique that utilizes the heat produced at the interface of the rotating
conical tool and workpiece in order to soften the workpiece and pierce a hole into it. In this work, experiments with thermal
drilling of galvanized steel were conducted based on the Taguchi L27 orthogonal array. Significant process parameters such as
rotational speed, tool angle and workpiece thickness were varied during the experimentation. In thermal drilling, the
thermal-drill tool pushes aside a large amount of workpiece material to form a sleeve, which is often referred to as the bushing
length. A predictive model for the bushing length was developed using a feed-forward artificial neural network based on experi-
mental data. As the bushing length is closely associated with the tapping process, the influences of the input process parameters
play a vital role in fastening galvanized steel with threaded fasteners in diverse engineering applications. The optimization
problem was solved by implementing a genetic algorithm under constraint limits to maximize the bushing length. Further, a
confirmation test was conducted with the intention to compare the optimum value and its corresponding bushing length
predicted by the genetic algorithm. Good agreement was observed between the predicted and the experimental values.
Keywords: thermal drilling, artificial neural network, genetic algorithm, galvanized steel, bushing length
Termi~no vrtanje je nova, za vrtanje lukenj v plo~evino, uporabljena tehnika, ki izkori{~a toploto, proizvedeno na povr{ini
vrte~e se konice orodja na obdelovancu z namenom, da ga zmeh~a in vanj naredi luknjo. V delu so bili izvedeni preizkusi na
osnovi metode Taguchi L27 z ortogonalno matriko s termi~nim vrtanjem galvaniziranega jekla. Pomembni parametri postopka,
kot so: hitrost vrtenja, kot orodja in debelina obdelovanca, so se med eksperimentiranjem spreminjali. Pri toplotnem vrtanju,
vrtanje orodja potisne stran ve~ materiala obdelovanca tako, da se tvori navarek (rokav) okoli luknje, ki se pogosto omenja kot
dol`ina {ablone. Napovedni model za dol`ino {ablone, je razvit z uporabo umetne nevronske mre`e, ki temelji na znanstvenih
podatkih. Ker je dol`ina {ablone precej povezana s procesom izdelave navoja, vplivi teh vhodnih procesnih parametrov igrajo
klju~no vlogo pri pritrditvi galvaniziranega jekla z navojem pritrdilnih elementov v razli~nih in`enirskih aplikacijah. Problem
optimizacije je bil re{en z implementacijo genetskega algoritma na podlagi omejitev za pove~anje dol`ine {ablone. Ugotovljeno
je bilo dobro ujemanje med napovedano in eksperimentalno vrednostjo.
Klju~ne besede: termi~no vrtanje, umetna nevronska mre`a, genetski algoritem, galvanizirano jeklo, dol`ina {ablone
1 INTRODUCTION
Thermal drilling is an emerging hole-making process
with significant breakthroughs in many drilling situ-
ations in both automobile and aerospace applications.
This process uses the heat energy produced at the inter-
face in order to soften and then pierce the workpiece.1
Moreover, this heat energy enhances the flow ability of
the workpiece material, which is extruded onto both the
front and back sides of a drilled hole. Finally, the
extruded or deformed material forms a bushing shape,
which surrounds the drilled hole.2 The length of the
bushing formed on the workpiece after a thermal-drilling
operation is called the bushing length. Achieving a
sufficient bushing length is very important since the
bushing length can increase the threading depth and the
clamp-load-bearing capability in various engineering
applications.
Contrary to the conventional drilling process, in
thermal drilling, there is no chip and wastage of the ma-
terial, since all the deformed material contributes to
developing a bushing. As it abolishes chip generation, it
can be seen as a clean, eco-friendly and chipless hole-
making technique. When joining metal sheets or thin-
walled structures by drilling holes using the conven-
tional-drilling process, just one or two threads can be
made in them; however, reliability is a matter of great
concern in such a situation. Alternatively, weld nuts and
threaded rivet nuts are used, but due to thermal
distortion, they get jammed and twisted during the
fastening of the structures.3 On the other hand, during the
thermal-drilling process, the bushing formed is about 3
to 4 times thicker than the workpiece. The bushing thus
allows a greater contact area and high structural rigidity
of fasteners in joining situations.4
Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822 813
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 519.6:620.1:67.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)813(2017)
A. H. Streppel and H. J. J Kals5 suggested that the
thermal-drilling process could be applied to different
materials like carbon steels, stainless steel, copper, brass
and aluminium. They carried out experimentation on the
thermal drilling of aluminium. Due to strong pre-strain
hardening, the process resulted in bad quality of the
bushing. M. Kerkhofs and M. Van Stappen6 compared
the performance of (Ti, Al) N coated tungsten carbide
thermal drills with uncoated drills. They reported that the
coated thermal drills allowed a longer tool life than the
uncoated flow drills. S. F. Miller et al.7 conducted
experiments with the thermal-drilling process on steel,
aluminum and titanium. They studied the properties such
as the hardness and microstructure of drilled holes. They
concluded that in thermal drilling large deformation and
high frictional heat are generated, resulting in the
changes in the material properties and microstructure of
a workpiece. They reported that the thermal conductivity
of the material drilled affects the quality and integrity of
the hole.
S. F. Miller et al.8 conducted an experimental and
numerical analysis of a thermal-drilling process on the
AISI 1020 cold-rolled carbon steel under a constant feed
rate. Moreover, they developed two classic models for
the thermal-drilling process. They predicted the tempe-
rature distribution using a finite-element-based thermal
model and in another attempt, they predicted the thrust
force and torque using the basic principles of physics. S.
F. Miller et al.9 thermally drilled into aluminium and
magnesium alloys using tungsten carbide and titanium
carbide in a cobalt matrix under different rotational
speeds and feed rates. They analyzed the thrust force,
torque, energy, power and peak power required for
drilling. They also evaluated the formation of the
bushing shape of cast metals. The bushings produced
during the thermal drilling of brittle cast metals demon-
strated a radial fracture. In their conclusions, they
suggested to pre-heat a cast metal workpiece and create a
high rotational speed for obtaining a cylindrically shaped
bushing without a significant radial fracture.
S. F. Miller et al.10 developed a three-dimensional
finite-element model for the thermal-drilling process in
order to evaluate the plastic strain and deformation of a
workpiece. S. F. Miller et al.11 investigated thermal-drill
tool-wear characteristics during the thermal drilling of an
AISI 1015 carbon steel workpiece. Furthermore, the
thrust force and torque were analyzed. They concluded
that the carbide tool is durable and it demonstrates the
least tool wear after drilling 11.000 holes. However,
progressively severe abrasive grooving on the tool tip
was observed. S. M. Lee et al.12 carried out thermal
drilling into the IN-713LC super alloy under different
rotation speeds and feed rates. They investigated the
material properties such as hardness, roundness, and
surface roughness of the holes drilled into the IN-713LC
super alloy. They reported that the hardness is higher
near the wall of a drilled hole and it decreases with the
increasing distance from the edge of the hole. In addition
to that, higher rotational speeds and feed rates demon-
strate better roundness and lower surface roughness.
H. M. Chow et al.13 optimized the process parameters
of a thermally drilled austenite stainless-steel workpiece
using the Taguchi method.14 They considered the input
parameters such as friction angle, friction/contact-area
ratio, feed rate and drilling speed, and studied their in-
fluence on the response parameters like surface rough-
ness. Moreover, the hardness and microstructural aspects
of drilled holes were studied. It was observed that the
surrounding area of a drilled hole acquired a fine grain
size and a compact structure with a higher microhardness
than that of the area away from the drilled region. S. M.
Lee et al.15 employed thermal drilling for the AISI 304
stainless steel using tungsten carbide drills with and
without coating. Their results illustrated that at the same
rotational speed and for the same number of holes
drilled, the coated drills experienced less tool wear than
the uncoated drills. Furthermore, they investigated the
changes in the relationship between the drilled-surface
temperature, tool wear and axial thrust force.
W. L. Ku et al.16 optimized the parameters of thermal
drilling into austenite stainless steel using the Taguchi
method. They studied the effects of the friction angle,
friction/contact-area ratio, feed rate and rotational speed
on the response-quality characteristics such as the
surface roughness and bushing length. They revealed that
at optimized drilling conditions, the bushing length of a
drilled hole was nearly three times longer than the plate
thickness, and a mirror-like quality wall surface of the
drilled hole could be obtained. M. Folea et al.17 investi-
gated the thermal drilling of a maraging-steel workpiece.
The authors reported that the temperature was the most
important factor in the thermal-drilling process. They
also revealed that a greater quality of the holes drilled
into maraging steel could be achieved with higher rota-
tional speeds. T. K. Mehmet et al.18 studied the effects of
the thermal-drilling parameters such as friction angle,
friction/contact-area ratio, feed rate and rotational speed
on workpiece temperature, thrust force and torque in
thermal drilling of the ST12 material. They revealed that
the thrust force and torque reduce with the increasing
rotational speed and increase with the friction angle, feed
rate and friction/contact-area ratio.
D. Biermann and Y. Liu19 demonstrated the feasibi-
lity of the thermal drilling of a magnesium wrought alloy
and analyzed the thrust forces and torque. They
measured the process temperature online and examined
the strength of the joint through tapping and thread
forming. B. B. Mehmet20 experimentally and numerically
investigated the thermal drilling of the AISI 1020 steel.
In their study, an analytical model for the torque, axial
power and heat-transfer coefficient was developed. Good
agreement was observed between experimental and
numerical values.
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
814 Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Similar work was performed by P. Krasauskas et al.21
on the thermal drilling into the AISI 304 steel. P. D.
Pantawane and B. B. Ahuja22 carried out regression
modelling of the thermal drilling of the AISI 1015 steel
using the Taguchi method. They studied the effects of the
tool-diameter ratio and rotational speed on the material
thickness and the effect of the feed on the response
parameters including the thrust force, torque and surface
roughness of drilled holes. G. Somasundaram et al.23
carried thermal drilling on the Al-SiC composite
material and studied the roundness errors on drilled
holes. They considered the input process parameters
such as the composition of workpiece, workpiece thick-
ness, rotational speed and feed rate. They developed an
empirical relation between the process parameters and
the roundness error using the response surface methodo-
logy.
Artificial neural network (ANN) is a biologically in-
spired computer program designed to simulate the way,
in which the human brain processes information.24 ANN
is widely applied in modeling many machining ope-
rations like turning, milling and drilling. S. R. Karnik
and V. N. Gaitone25 used a multilayer feed-forward ANN
to predict the influence of the process parameters such as
the cutting speed, feed, drill diameter, point angle and
lip-clearance angle on the burr height and burr thickness
in drilling the AISI 316L stainless steel. R. S. Mamilla et
al.26 applied a multilayer ANN for modeling an abrasive
flow-finishing process on the AISI 1040 and AISI 4340
steel. S. Assarzadeh and M. Ghoreishi27 successfully
used a multilayer ANN for modeling the metal-removal
rate and surface roughness in an electrical discharge
machining process involving the BD3 steel. O. Babur et
al.28 developed an ANN model according to experimental
measurement data for the end milling of Inconel 718.
They coupled a genetic algorithm (GA) with an ANN to
forecast the surface roughness. S. Sarkar et al.29 pro-
duced a multilayer feed-forward-ANN model to predict
the process parameters of the machining of  titanium
aluminide with a wire-electrical-discharge machine. A.
K. Singh et al.30 used a multilayer feed-forward ANN to
predict the flank wear of high-speed steel drill bits for
drilling holes into a copper workpiece.
The literature review on thermal drilling reveals that
experimental investigations and numerical simulations of
the effects of mechanical and physical properties on the
thermal-drilling process was mainly carried out on
copper, brass, aluminium, magnesium and stainless-steel
alloys, maraging steel, ST12 steel and IN-713LC super
alloy. Galvanized steel is being widely used for car body
structures and it is a good candidate for the implemen-
tation of the thermal-drilling process. The application of
galvanized steel is widely used in the areas of roofing
material, doors, ship’s ducts and panels, electrical-
appliance automobile parts, etc.31
In the present work, thermal drilling was carried out
on galvanized steel. Based on experimental measurement
data, an ANN model was developed to predict the for-
mation of the bushing length in thermal drilling of
galvanized steel. This model considers rotational speed,
tool angle and workpiece thickness as significant ther-
mal-drilling process parameters. Then the optimum
values of these drilling parameters were computed, with
the aim of achieving the maximum bushing length, by
solving the optimization problem implementing a GA.
Thus, the present work allows us to achieve the maxi-
mum bushing length, which is highly desirable for the
subsequent tapping and joining in automotive applica-
tions. A schematic diagram of the combined ANN-GA
optimization24 is shown in Figure 1. It indicates the
stages of the GA process and its relation with the ANN
process.
2 THERMAL DRILLING
2.1 Physical description of the process
Figure 2 shows a schematic representation of the
thermal-drilling process. Based on the thermal-drill
geometry, there are four steps involved in thermal drill-
ing. Initially, the center point zone of a drill approaches
and pierces the workpiece. Secondly, the produced heat
softens the workpiece as a result of the friction between
the thermal drill and the workpiece. Thirdly, the softened
material is pushed downward and the drill moves for-
ward to form the bushing using the cylindrical zone of
the tool. Fourthly, the extruded burr on the workpiece
surface is pressed to form a boss with the shoulder zone
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822 815
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Flow chart of the integrated ANN-GA optimization process
A. H. Streppel and H. J. J Kals5 suggested that the
thermal-drilling process could be applied to different
materials like carbon steels, stainless steel, copper, brass
and aluminium. They carried out experimentation on the
thermal drilling of aluminium. Due to strong pre-strain
hardening, the process resulted in bad quality of the
bushing. M. Kerkhofs and M. Van Stappen6 compared
the performance of (Ti, Al) N coated tungsten carbide
thermal drills with uncoated drills. They reported that the
coated thermal drills allowed a longer tool life than the
uncoated flow drills. S. F. Miller et al.7 conducted
experiments with the thermal-drilling process on steel,
aluminum and titanium. They studied the properties such
as the hardness and microstructure of drilled holes. They
concluded that in thermal drilling large deformation and
high frictional heat are generated, resulting in the
changes in the material properties and microstructure of
a workpiece. They reported that the thermal conductivity
of the material drilled affects the quality and integrity of
the hole.
S. F. Miller et al.8 conducted an experimental and
numerical analysis of a thermal-drilling process on the
AISI 1020 cold-rolled carbon steel under a constant feed
rate. Moreover, they developed two classic models for
the thermal-drilling process. They predicted the tempe-
rature distribution using a finite-element-based thermal
model and in another attempt, they predicted the thrust
force and torque using the basic principles of physics. S.
F. Miller et al.9 thermally drilled into aluminium and
magnesium alloys using tungsten carbide and titanium
carbide in a cobalt matrix under different rotational
speeds and feed rates. They analyzed the thrust force,
torque, energy, power and peak power required for
drilling. They also evaluated the formation of the
bushing shape of cast metals. The bushings produced
during the thermal drilling of brittle cast metals demon-
strated a radial fracture. In their conclusions, they
suggested to pre-heat a cast metal workpiece and create a
high rotational speed for obtaining a cylindrically shaped
bushing without a significant radial fracture.
S. F. Miller et al.10 developed a three-dimensional
finite-element model for the thermal-drilling process in
order to evaluate the plastic strain and deformation of a
workpiece. S. F. Miller et al.11 investigated thermal-drill
tool-wear characteristics during the thermal drilling of an
AISI 1015 carbon steel workpiece. Furthermore, the
thrust force and torque were analyzed. They concluded
that the carbide tool is durable and it demonstrates the
least tool wear after drilling 11.000 holes. However,
progressively severe abrasive grooving on the tool tip
was observed. S. M. Lee et al.12 carried out thermal
drilling into the IN-713LC super alloy under different
rotation speeds and feed rates. They investigated the
material properties such as hardness, roundness, and
surface roughness of the holes drilled into the IN-713LC
super alloy. They reported that the hardness is higher
near the wall of a drilled hole and it decreases with the
increasing distance from the edge of the hole. In addition
to that, higher rotational speeds and feed rates demon-
strate better roundness and lower surface roughness.
H. M. Chow et al.13 optimized the process parameters
of a thermally drilled austenite stainless-steel workpiece
using the Taguchi method.14 They considered the input
parameters such as friction angle, friction/contact-area
ratio, feed rate and drilling speed, and studied their in-
fluence on the response parameters like surface rough-
ness. Moreover, the hardness and microstructural aspects
of drilled holes were studied. It was observed that the
surrounding area of a drilled hole acquired a fine grain
size and a compact structure with a higher microhardness
than that of the area away from the drilled region. S. M.
Lee et al.15 employed thermal drilling for the AISI 304
stainless steel using tungsten carbide drills with and
without coating. Their results illustrated that at the same
rotational speed and for the same number of holes
drilled, the coated drills experienced less tool wear than
the uncoated drills. Furthermore, they investigated the
changes in the relationship between the drilled-surface
temperature, tool wear and axial thrust force.
W. L. Ku et al.16 optimized the parameters of thermal
drilling into austenite stainless steel using the Taguchi
method. They studied the effects of the friction angle,
friction/contact-area ratio, feed rate and rotational speed
on the response-quality characteristics such as the
surface roughness and bushing length. They revealed that
at optimized drilling conditions, the bushing length of a
drilled hole was nearly three times longer than the plate
thickness, and a mirror-like quality wall surface of the
drilled hole could be obtained. M. Folea et al.17 investi-
gated the thermal drilling of a maraging-steel workpiece.
The authors reported that the temperature was the most
important factor in the thermal-drilling process. They
also revealed that a greater quality of the holes drilled
into maraging steel could be achieved with higher rota-
tional speeds. T. K. Mehmet et al.18 studied the effects of
the thermal-drilling parameters such as friction angle,
friction/contact-area ratio, feed rate and rotational speed
on workpiece temperature, thrust force and torque in
thermal drilling of the ST12 material. They revealed that
the thrust force and torque reduce with the increasing
rotational speed and increase with the friction angle, feed
rate and friction/contact-area ratio.
D. Biermann and Y. Liu19 demonstrated the feasibi-
lity of the thermal drilling of a magnesium wrought alloy
and analyzed the thrust forces and torque. They
measured the process temperature online and examined
the strength of the joint through tapping and thread
forming. B. B. Mehmet20 experimentally and numerically
investigated the thermal drilling of the AISI 1020 steel.
In their study, an analytical model for the torque, axial
power and heat-transfer coefficient was developed. Good
agreement was observed between experimental and
numerical values.
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
814 Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Similar work was performed by P. Krasauskas et al.21
on the thermal drilling into the AISI 304 steel. P. D.
Pantawane and B. B. Ahuja22 carried out regression
modelling of the thermal drilling of the AISI 1015 steel
using the Taguchi method. They studied the effects of the
tool-diameter ratio and rotational speed on the material
thickness and the effect of the feed on the response
parameters including the thrust force, torque and surface
roughness of drilled holes. G. Somasundaram et al.23
carried thermal drilling on the Al-SiC composite
material and studied the roundness errors on drilled
holes. They considered the input process parameters
such as the composition of workpiece, workpiece thick-
ness, rotational speed and feed rate. They developed an
empirical relation between the process parameters and
the roundness error using the response surface methodo-
logy.
Artificial neural network (ANN) is a biologically in-
spired computer program designed to simulate the way,
in which the human brain processes information.24 ANN
is widely applied in modeling many machining ope-
rations like turning, milling and drilling. S. R. Karnik
and V. N. Gaitone25 used a multilayer feed-forward ANN
to predict the influence of the process parameters such as
the cutting speed, feed, drill diameter, point angle and
lip-clearance angle on the burr height and burr thickness
in drilling the AISI 316L stainless steel. R. S. Mamilla et
al.26 applied a multilayer ANN for modeling an abrasive
flow-finishing process on the AISI 1040 and AISI 4340
steel. S. Assarzadeh and M. Ghoreishi27 successfully
used a multilayer ANN for modeling the metal-removal
rate and surface roughness in an electrical discharge
machining process involving the BD3 steel. O. Babur et
al.28 developed an ANN model according to experimental
measurement data for the end milling of Inconel 718.
They coupled a genetic algorithm (GA) with an ANN to
forecast the surface roughness. S. Sarkar et al.29 pro-
duced a multilayer feed-forward-ANN model to predict
the process parameters of the machining of  titanium
aluminide with a wire-electrical-discharge machine. A.
K. Singh et al.30 used a multilayer feed-forward ANN to
predict the flank wear of high-speed steel drill bits for
drilling holes into a copper workpiece.
The literature review on thermal drilling reveals that
experimental investigations and numerical simulations of
the effects of mechanical and physical properties on the
thermal-drilling process was mainly carried out on
copper, brass, aluminium, magnesium and stainless-steel
alloys, maraging steel, ST12 steel and IN-713LC super
alloy. Galvanized steel is being widely used for car body
structures and it is a good candidate for the implemen-
tation of the thermal-drilling process. The application of
galvanized steel is widely used in the areas of roofing
material, doors, ship’s ducts and panels, electrical-
appliance automobile parts, etc.31
In the present work, thermal drilling was carried out
on galvanized steel. Based on experimental measurement
data, an ANN model was developed to predict the for-
mation of the bushing length in thermal drilling of
galvanized steel. This model considers rotational speed,
tool angle and workpiece thickness as significant ther-
mal-drilling process parameters. Then the optimum
values of these drilling parameters were computed, with
the aim of achieving the maximum bushing length, by
solving the optimization problem implementing a GA.
Thus, the present work allows us to achieve the maxi-
mum bushing length, which is highly desirable for the
subsequent tapping and joining in automotive applica-
tions. A schematic diagram of the combined ANN-GA
optimization24 is shown in Figure 1. It indicates the
stages of the GA process and its relation with the ANN
process.
2 THERMAL DRILLING
2.1 Physical description of the process
Figure 2 shows a schematic representation of the
thermal-drilling process. Based on the thermal-drill
geometry, there are four steps involved in thermal drill-
ing. Initially, the center point zone of a drill approaches
and pierces the workpiece. Secondly, the produced heat
softens the workpiece as a result of the friction between
the thermal drill and the workpiece. Thirdly, the softened
material is pushed downward and the drill moves for-
ward to form the bushing using the cylindrical zone of
the tool. Fourthly, the extruded burr on the workpiece
surface is pressed to form a boss with the shoulder zone
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822 815
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Flow chart of the integrated ANN-GA optimization process
of the thermal drill. Finally, the thermal drill retracts
leaving a hole with a bushing length.37
2.2 Formation of the bushing length
The initial volume of the material (Vi) available to
produce a hole by thermal drilling is given in Equation
(1).33 As shown in Figure 2, D1 and Pt represent the
inner-hole diameter in mm and the thickness of the
workpiece in mm, respectively.
V D Pi =
π
4 1
2
t (1)
During thermal drilling, the volume of the material to
be displaced (Vf) in order to produce a bushing is given
in Equation (2).33 As shown in Figure 2, D2 and L
represent the outer-bushing diameter in mm and the
bushing length in mm, respectively.
[ ]V D D Lf = −
π
4 2
2
1
2 (2)
It is known that the initial approximation of the
volume of the material for producing a hole is the same
as the volume of the material to be displaced during the
formation of the bushing. Therefore, Equations (1) and
(2) are assumed to be equal. Equating Equations (1) and
(2), Equation (3) is obtained:
[ ]π π
4 41
2
2
2
1
2D P D D Lt = − (3)
While rearranging the terms in Equation (3), the
bushing length (L) can be determined as shown in
Equation (4):
[ ]L
D
D D
P=
−
1
2
2
2
1
2 t
(4)
3 EXPERIMENTAL PART
In the present work, the experiments were carried out
at in a vertical drilling machine. A hot-dip-galvanized
steel workpiece with thickness values of 1, 1.5, 2 mm
and with dimensions of 50 mm × 150 mm was used.
6-mm-diameter thermal drills with different tool angles
were used in the experimentation. They were machined
out of high-speed steel rods, having the required
dimensions as shown in Figure 3. The feed rate of the
thermal-drilling process is kept constant at 240 mm/min.
The chemical composition of the galvanized steel is
given in Table 1. Axial forces were measured using a
digital drilling-tool dynamometer. Temperature measure-
ments were done on-line using a non-contact-type
infra-red thermometer and the temperature profile of the
process was obtained using the DATA TEMP MX
software with a time increment of 0.016 s. The micro-
structure of the holes drilled into galvanized steel was
measured using a scanning electron microscope (Carl
Zeiss EVO18, Germany). Three different levels of the
input process parameters were used, as shown in Table
2.
Based on the proper selection of the input process
parameters, the desirable output parameters can be
obtained during the experimentation. The selected input
process parameters such as the rotational speed, tool
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
816 Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Obtained thermal drills with different tool angles: a) 35 de-
gree tool angle b) 37.5 degree tool angle, c) 45 degree tool angle
Table 1: Chemical composition of galvanized steel (w/%)
C Si S P Mn Ni Cr
0.003 0.006 0.005 0.018 0.173 0.011 0.031
Mo V Cu Al Nb Zn Ti Fe
0.001 0.002 0.017 0.035 0.001 0.004 0.05 Balance
Table 2: Input process parameters and their levels
Factors Levels
1 2 3
Rotational speed (min–1) 1944 2772 3600
Tool angle (°) 30 37.5 45
Workpiece thickness (mm) 1 1.5 2Figure 2: Schematic representation of the thermal-drilling process
angle and workpiece thickness, and the output parameter,
the bushing length of the thermally drilled galvanized
steel sheet, were analyzed in this work. Based on the
Taguchi’s L27 orthogonal array, experiments were con-
ducted at three levels of the drilling-process parameters
and the corresponding bushing length was measured
using a digital vernier caliper (Model No CD-12C, Mitu-
toya Corporation, Japan). All the experimental values of
the bushing length are displayed in Table 3. Samples of
the thermally drilled holes are shown in Figure 4.
During the thermal drilling, the burr is extruded onto the
top surface, and it is leveled when it comes in contact
with the shoulder of the thermal-drill tool as shown in
Figure 5a. The formation of the bushing at the rear side
of the galvanized steel plate is shown in Figure 5b. It
can be used to act as a cylindrical sleeve bearing and it
can be taped to fabricate an internal screw. Additionally,
it widely reduces the complexity of joining components
in the aerospace and automotive industries.
Table 3: Experimental results of the thermal-drilling process
Expt No.
Rotational
speed
(min–1)
Tool angle
(°)
Workpiece
thickness
(mm)
Bushing
length
(mm)
1 1944 30 1 2.54
2 1944 30 1.5 3.42
3 1944 30 2 4.38
4 1944 37.5 1 3.24
5 1944 37.5 1.5 3.21
6 1944 37.5 2 4.41
7 1944 45 1 2.97
8 1944 45 1.5 4.87
9 1944 45 2 5.26
10 2772 30 1 2.67
11 2772 30 1.5 3.84
12 2772 30 2 4.02
13 2772 37.5 1 2.91
14 2772 37.5 1.5 3.83
15 2772 37.5 2 4.42
16 2772 45 1 2.94
17 2772 45 1.5 4.31
18 2772 45 2 5.57
19 3600 30 1 3.55
20 3600 30 1.5 4.64
21 3600 30 2 5.84
22 3600 37.5 1 2.78
23 3600 37.5 1.5 3.85
24 3600 37.5 2 5.92
25 3600 45 1 2.42
26 3600 45 1.5 4.42
27 3600 45 2 5.61
4 ARTIFICIAL-NEURAL-NETWORK
MODELING
Recently, the use of the artificial intelligence perfor-
mance has been remarkable in the entire engineering
field.24 To understand and manage any industrial course
of action, modeling and optimization of the process are
essential. An accurate control is a requirement necessary
to accomplish better quality and productivity. ANN plays
a significant role in forecasting the solutions for non-
linear problems in all engineering fields. Using statistical
techniques, numerous researchers endeavored to develop
a model based on experimental data. In general, a neural
network represents a network of many processors ope-
rating in parallel. Each processor contains a small size of
the local memory. Then, these units are coupled through
communication channels, which typically carry numeric
data. An excellent instance of a biological neural net-
work is the brain of a human. It comprises the most
demanding and controlling arrangement, in which
learning along with training controls the behavior of a
human to take action to solve any problem met in
day-by-day life.
ANN can be successfully employed to predict the
output parameters for any given input parameters, based
on the training set for a given complex problem.28 In the
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822 817
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Appearance of a hole made by thermal drilling: a) top view
and b) rear view
Figure 4: Holes made by thermal drilling in galvanized steel plates
of the thermal drill. Finally, the thermal drill retracts
leaving a hole with a bushing length.37
2.2 Formation of the bushing length
The initial volume of the material (Vi) available to
produce a hole by thermal drilling is given in Equation
(1).33 As shown in Figure 2, D1 and Pt represent the
inner-hole diameter in mm and the thickness of the
workpiece in mm, respectively.
V D Pi =
π
4 1
2
t (1)
During thermal drilling, the volume of the material to
be displaced (Vf) in order to produce a bushing is given
in Equation (2).33 As shown in Figure 2, D2 and L
represent the outer-bushing diameter in mm and the
bushing length in mm, respectively.
[ ]V D D Lf = −
π
4 2
2
1
2 (2)
It is known that the initial approximation of the
volume of the material for producing a hole is the same
as the volume of the material to be displaced during the
formation of the bushing. Therefore, Equations (1) and
(2) are assumed to be equal. Equating Equations (1) and
(2), Equation (3) is obtained:
[ ]π π
4 41
2
2
2
1
2D P D D Lt = − (3)
While rearranging the terms in Equation (3), the
bushing length (L) can be determined as shown in
Equation (4):
[ ]L
D
D D
P=
−
1
2
2
2
1
2 t
(4)
3 EXPERIMENTAL PART
In the present work, the experiments were carried out
at in a vertical drilling machine. A hot-dip-galvanized
steel workpiece with thickness values of 1, 1.5, 2 mm
and with dimensions of 50 mm × 150 mm was used.
6-mm-diameter thermal drills with different tool angles
were used in the experimentation. They were machined
out of high-speed steel rods, having the required
dimensions as shown in Figure 3. The feed rate of the
thermal-drilling process is kept constant at 240 mm/min.
The chemical composition of the galvanized steel is
given in Table 1. Axial forces were measured using a
digital drilling-tool dynamometer. Temperature measure-
ments were done on-line using a non-contact-type
infra-red thermometer and the temperature profile of the
process was obtained using the DATA TEMP MX
software with a time increment of 0.016 s. The micro-
structure of the holes drilled into galvanized steel was
measured using a scanning electron microscope (Carl
Zeiss EVO18, Germany). Three different levels of the
input process parameters were used, as shown in Table
2.
Based on the proper selection of the input process
parameters, the desirable output parameters can be
obtained during the experimentation. The selected input
process parameters such as the rotational speed, tool
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
816 Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Obtained thermal drills with different tool angles: a) 35 de-
gree tool angle b) 37.5 degree tool angle, c) 45 degree tool angle
Table 1: Chemical composition of galvanized steel (w/%)
C Si S P Mn Ni Cr
0.003 0.006 0.005 0.018 0.173 0.011 0.031
Mo V Cu Al Nb Zn Ti Fe
0.001 0.002 0.017 0.035 0.001 0.004 0.05 Balance
Table 2: Input process parameters and their levels
Factors Levels
1 2 3
Rotational speed (min–1) 1944 2772 3600
Tool angle (°) 30 37.5 45
Workpiece thickness (mm) 1 1.5 2Figure 2: Schematic representation of the thermal-drilling process
angle and workpiece thickness, and the output parameter,
the bushing length of the thermally drilled galvanized
steel sheet, were analyzed in this work. Based on the
Taguchi’s L27 orthogonal array, experiments were con-
ducted at three levels of the drilling-process parameters
and the corresponding bushing length was measured
using a digital vernier caliper (Model No CD-12C, Mitu-
toya Corporation, Japan). All the experimental values of
the bushing length are displayed in Table 3. Samples of
the thermally drilled holes are shown in Figure 4.
During the thermal drilling, the burr is extruded onto the
top surface, and it is leveled when it comes in contact
with the shoulder of the thermal-drill tool as shown in
Figure 5a. The formation of the bushing at the rear side
of the galvanized steel plate is shown in Figure 5b. It
can be used to act as a cylindrical sleeve bearing and it
can be taped to fabricate an internal screw. Additionally,
it widely reduces the complexity of joining components
in the aerospace and automotive industries.
Table 3: Experimental results of the thermal-drilling process
Expt No.
Rotational
speed
(min–1)
Tool angle
(°)
Workpiece
thickness
(mm)
Bushing
length
(mm)
1 1944 30 1 2.54
2 1944 30 1.5 3.42
3 1944 30 2 4.38
4 1944 37.5 1 3.24
5 1944 37.5 1.5 3.21
6 1944 37.5 2 4.41
7 1944 45 1 2.97
8 1944 45 1.5 4.87
9 1944 45 2 5.26
10 2772 30 1 2.67
11 2772 30 1.5 3.84
12 2772 30 2 4.02
13 2772 37.5 1 2.91
14 2772 37.5 1.5 3.83
15 2772 37.5 2 4.42
16 2772 45 1 2.94
17 2772 45 1.5 4.31
18 2772 45 2 5.57
19 3600 30 1 3.55
20 3600 30 1.5 4.64
21 3600 30 2 5.84
22 3600 37.5 1 2.78
23 3600 37.5 1.5 3.85
24 3600 37.5 2 5.92
25 3600 45 1 2.42
26 3600 45 1.5 4.42
27 3600 45 2 5.61
4 ARTIFICIAL-NEURAL-NETWORK
MODELING
Recently, the use of the artificial intelligence perfor-
mance has been remarkable in the entire engineering
field.24 To understand and manage any industrial course
of action, modeling and optimization of the process are
essential. An accurate control is a requirement necessary
to accomplish better quality and productivity. ANN plays
a significant role in forecasting the solutions for non-
linear problems in all engineering fields. Using statistical
techniques, numerous researchers endeavored to develop
a model based on experimental data. In general, a neural
network represents a network of many processors ope-
rating in parallel. Each processor contains a small size of
the local memory. Then, these units are coupled through
communication channels, which typically carry numeric
data. An excellent instance of a biological neural net-
work is the brain of a human. It comprises the most
demanding and controlling arrangement, in which
learning along with training controls the behavior of a
human to take action to solve any problem met in
day-by-day life.
ANN can be successfully employed to predict the
output parameters for any given input parameters, based
on the training set for a given complex problem.28 In the
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822 817
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Appearance of a hole made by thermal drilling: a) top view
and b) rear view
Figure 4: Holes made by thermal drilling in galvanized steel plates
present work, an ANN is developed using MATLAB
version 7.10, neural network toolbox, with the intention
of predicting the bushing length as a function of three
input process parameters including the rotational speed,
tool angle and workpiece thickness. Figure 6 shows the
neural-network architecture of the bushing length ob-
tained with thermal drilling. The input data is provided
through input layer neurons, and it is fed forward
through the hidden layer. Then the response is obtained
at the output layer neurons. The neurons are linked by
weights that take part in the learning process. The
weights help in attaining the optimum solution and
escape from the local minima. To establish the optimum
structural design, we apply the trial-and-error method
with the appropriate activation function and the most
excellent training algorithm. A number of models were
created and tested. A hyperbolic tangent sigmoid transfer
function is considered as the hidden-layer activation
function, and it is given in Equation (5):
f x x
e x
( ) ( )= =
−
−−tansig
2
1
12 (5)
The most important selection criteria used for the
suitable model are the mean squared error and the
regression-coefficient value. The mean sum of the
squared error (MSE) is the average squared difference
between the predicted value and the actual experimental
value as shown in Equation (6). A lower value of the
MSE indicates a better model. The regression coefficient
(R) was measured by correlating the predicted and the
actual values as given in Equation (7). An R value of 1
indicates a close relationship between the predicted and
actual experimental value, i.e., a very small error.
MSE
n
A Pi i
i
= −∑1 2 (6)
R
A P
P
i i
i
n
i
i
n= −
−
⎛
⎝
⎜
⎜
⎜
⎞
⎠
⎟
⎟
⎟
=
=
∑
∑
1
2
1
2
1
(7)
where Ai is the actual experimental value, Pi is the
predicted value and n is the number of patterns. With
the intention of measuring the accuracy of the prediction
model, the percentage of error is calculated with
Equation (8):
z
E P
E
=
−
×100 (8)
where z = % of prediction error, E = experimental value,
P = predicted value
5 RESULTS AND DISCUSSION
The results of the ANN coupled with the GA34 are
utilized to predict and optimize the bushing length based
on the input process parameters such as the rotational
speed, tool angle and workpiece thickness in the ther-
mal-drilling process, and they are discussed below.
5.1 Prediction of the bushing length by the ANN
Table 4 displays the parameters used with this
modeling technique. The developed ANN model of the
bushing length was trained by means of the selected
parameters. During the training process, a 19 set input
for training trials were presented each time and the
resultant output was acquired. The number of the
hidden-layer neurons varied from 1 to 20. It was deter-
mined based on trial and error, through step-by-step
increasing the number of neurons and examining their
result against the forecast value. The final neuron
number of the hidden layer obtained was 10. The ulti-
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
818 Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7: Correlation between the predicted values of the ANN model
and the experimental data for the prediction of the bushing length
using: a) training data, b) validation data, c) test data and d) entire data
Figure 6: ANN modeling of the bushing length in thermal drilling
mate weights between the input layer and hidden layer
(w1, w2, w3), the bias (b) and weights between the
hidden layer and the output layer (w4) are shown in
Table 5. Finally, the lowest MSE value achieved was
0.00415 when the configuration of the ANN was 3-10-1.
There are three input neurons, each indicating one input
variable form the input layer, 10 neurons in the hidden
layer and the output layer with one neuron correspond to
the bushing length.
Table 4: Parameters used in the ANN modeling
Input variables Rotational speed, tool angle,workpiece thickness
Output variable Bushing length
Network type Feed-forward backpropagation
Algorithm
Levenberg-Marquardt
back-propagation algorithm
(trainlm)
Activation function Hyperbolic tangent sigmoidtransfer function
Data division 70 % (training); 15 %(validation); 15 % (testing)
Number of neurons in the
hidden layer 10
Table 5: Final weights in between the layers
No. of
neurons
Weights
Input to hidden layer Hidden tooutput layer
w1 w2 w3 b w4
1 -0.682 -1.572 1.929 3.529 0.680
2 -0.865 1.620 -2.115 2.483 -0.576
3 3.045 0.934 -0.191 -1.179 0.111
4 1.722 -2.097 1.179 -1.184 -0.253
5 -1.826 -1.470 1.863 0.273 -0.183
6 -0.027 -0.448 3.025 -0.131 1.113
7 -0.619 2.955 0.332 -0.694 0.268
8 0.855 -1.318 -2.373 1.895 0.036
9 1.645 0.306 -2.767 2.418 0.939
10 1.822 -1.597 -1.613 3.188 -0.638
The output performance of the established ANN is
examined on the basis of the regression-correlation
coefficient (R value) between the predicted values and
the experimental values. The predicted responses of the
ANN model are in excellent conformity with the expe-
rimental values, i.e., the correlation regression coeffi-
cients of 100 % for the training set, 98.41 % for the
validation set and 94.80 % for the testing set are
achieved as shown in Figure 7. It was concluded that the
feed-forward back-propagation algorithm of the configu-
ration of 3-10-1 gives the most excellent results for the
prediction of the bushing length as shown in Table 6.
The graphical representation of the experimental and
predicted values of the bushing length by the ANN is
shown in Figure 8. It shows that the ANN prediction
values are very close to the experimental values of the
bushing length. The average prediction error between the
experimental and predicted values of the bushing length
is 1.843 %. Therefore, it is demonstrated that the
developed ANN model is suitable for forecasting the
bushing length of drilled holes in thermal drilling of
galvanized steel, having the highest accuracy of
98.157 %.
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822 819
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: Correlation between experimental and predicted values
Table 6: Comparison of experimental and predicted values of the
bushing length
Expt No. Bushing length (mm)
Prediction
error (%)Experimental
value
Predicted ANN
value
1 2.54 2.485 2.165
2 3.42 3.427 0.204
3 4.38 4.252 2.922
4 3.24 2.618 19.197
5 3.21 3.336 3.925
6 4.41 4.425 0.340
7 2.97 2.843 4.276
8 4.87 4.677 3.963
9 5.26 5.517 4.886
10 2.67 2.385 10.674
11 3.84 3.714 3.281
12 4.02 4.275 6.343
13 2.91 2.905 0.172
14 3.83 3.814 0.418
15 4.42 4.615 4.412
16 2.94 2.964 0.816
17 4.31 4.284 0.603
18 5.57 5.599 0.521
19 3.55 3.738 5.296
20 4.64 4.585 1.185
21 5.84 5.893 0.907
22 2.78 2.760 0.719
23 3.85 3.875 0.649
24 5.92 5.836 1.419
25 2.42 2.755 13.843
26 4.42 4.450 0.679
27 5.61 6.007 7.077
present work, an ANN is developed using MATLAB
version 7.10, neural network toolbox, with the intention
of predicting the bushing length as a function of three
input process parameters including the rotational speed,
tool angle and workpiece thickness. Figure 6 shows the
neural-network architecture of the bushing length ob-
tained with thermal drilling. The input data is provided
through input layer neurons, and it is fed forward
through the hidden layer. Then the response is obtained
at the output layer neurons. The neurons are linked by
weights that take part in the learning process. The
weights help in attaining the optimum solution and
escape from the local minima. To establish the optimum
structural design, we apply the trial-and-error method
with the appropriate activation function and the most
excellent training algorithm. A number of models were
created and tested. A hyperbolic tangent sigmoid transfer
function is considered as the hidden-layer activation
function, and it is given in Equation (5):
f x x
e x
( ) ( )= =
−
−−tansig
2
1
12 (5)
The most important selection criteria used for the
suitable model are the mean squared error and the
regression-coefficient value. The mean sum of the
squared error (MSE) is the average squared difference
between the predicted value and the actual experimental
value as shown in Equation (6). A lower value of the
MSE indicates a better model. The regression coefficient
(R) was measured by correlating the predicted and the
actual values as given in Equation (7). An R value of 1
indicates a close relationship between the predicted and
actual experimental value, i.e., a very small error.
MSE
n
A Pi i
i
= −∑1 2 (6)
R
A P
P
i i
i
n
i
i
n= −
−
⎛
⎝
⎜
⎜
⎜
⎞
⎠
⎟
⎟
⎟
=
=
∑
∑
1
2
1
2
1
(7)
where Ai is the actual experimental value, Pi is the
predicted value and n is the number of patterns. With
the intention of measuring the accuracy of the prediction
model, the percentage of error is calculated with
Equation (8):
z
E P
E
=
−
×100 (8)
where z = % of prediction error, E = experimental value,
P = predicted value
5 RESULTS AND DISCUSSION
The results of the ANN coupled with the GA34 are
utilized to predict and optimize the bushing length based
on the input process parameters such as the rotational
speed, tool angle and workpiece thickness in the ther-
mal-drilling process, and they are discussed below.
5.1 Prediction of the bushing length by the ANN
Table 4 displays the parameters used with this
modeling technique. The developed ANN model of the
bushing length was trained by means of the selected
parameters. During the training process, a 19 set input
for training trials were presented each time and the
resultant output was acquired. The number of the
hidden-layer neurons varied from 1 to 20. It was deter-
mined based on trial and error, through step-by-step
increasing the number of neurons and examining their
result against the forecast value. The final neuron
number of the hidden layer obtained was 10. The ulti-
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
818 Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7: Correlation between the predicted values of the ANN model
and the experimental data for the prediction of the bushing length
using: a) training data, b) validation data, c) test data and d) entire data
Figure 6: ANN modeling of the bushing length in thermal drilling
mate weights between the input layer and hidden layer
(w1, w2, w3), the bias (b) and weights between the
hidden layer and the output layer (w4) are shown in
Table 5. Finally, the lowest MSE value achieved was
0.00415 when the configuration of the ANN was 3-10-1.
There are three input neurons, each indicating one input
variable form the input layer, 10 neurons in the hidden
layer and the output layer with one neuron correspond to
the bushing length.
Table 4: Parameters used in the ANN modeling
Input variables Rotational speed, tool angle,workpiece thickness
Output variable Bushing length
Network type Feed-forward backpropagation
Algorithm
Levenberg-Marquardt
back-propagation algorithm
(trainlm)
Activation function Hyperbolic tangent sigmoidtransfer function
Data division 70 % (training); 15 %(validation); 15 % (testing)
Number of neurons in the
hidden layer 10
Table 5: Final weights in between the layers
No. of
neurons
Weights
Input to hidden layer Hidden tooutput layer
w1 w2 w3 b w4
1 -0.682 -1.572 1.929 3.529 0.680
2 -0.865 1.620 -2.115 2.483 -0.576
3 3.045 0.934 -0.191 -1.179 0.111
4 1.722 -2.097 1.179 -1.184 -0.253
5 -1.826 -1.470 1.863 0.273 -0.183
6 -0.027 -0.448 3.025 -0.131 1.113
7 -0.619 2.955 0.332 -0.694 0.268
8 0.855 -1.318 -2.373 1.895 0.036
9 1.645 0.306 -2.767 2.418 0.939
10 1.822 -1.597 -1.613 3.188 -0.638
The output performance of the established ANN is
examined on the basis of the regression-correlation
coefficient (R value) between the predicted values and
the experimental values. The predicted responses of the
ANN model are in excellent conformity with the expe-
rimental values, i.e., the correlation regression coeffi-
cients of 100 % for the training set, 98.41 % for the
validation set and 94.80 % for the testing set are
achieved as shown in Figure 7. It was concluded that the
feed-forward back-propagation algorithm of the configu-
ration of 3-10-1 gives the most excellent results for the
prediction of the bushing length as shown in Table 6.
The graphical representation of the experimental and
predicted values of the bushing length by the ANN is
shown in Figure 8. It shows that the ANN prediction
values are very close to the experimental values of the
bushing length. The average prediction error between the
experimental and predicted values of the bushing length
is 1.843 %. Therefore, it is demonstrated that the
developed ANN model is suitable for forecasting the
bushing length of drilled holes in thermal drilling of
galvanized steel, having the highest accuracy of
98.157 %.
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822 819
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: Correlation between experimental and predicted values
Table 6: Comparison of experimental and predicted values of the
bushing length
Expt No. Bushing length (mm)
Prediction
error (%)Experimental
value
Predicted ANN
value
1 2.54 2.485 2.165
2 3.42 3.427 0.204
3 4.38 4.252 2.922
4 3.24 2.618 19.197
5 3.21 3.336 3.925
6 4.41 4.425 0.340
7 2.97 2.843 4.276
8 4.87 4.677 3.963
9 5.26 5.517 4.886
10 2.67 2.385 10.674
11 3.84 3.714 3.281
12 4.02 4.275 6.343
13 2.91 2.905 0.172
14 3.83 3.814 0.418
15 4.42 4.615 4.412
16 2.94 2.964 0.816
17 4.31 4.284 0.603
18 5.57 5.599 0.521
19 3.55 3.738 5.296
20 4.64 4.585 1.185
21 5.84 5.893 0.907
22 2.78 2.760 0.719
23 3.85 3.875 0.649
24 5.92 5.836 1.419
25 2.42 2.755 13.843
26 4.42 4.450 0.679
27 5.61 6.007 7.077
5.2 Optimization of the bushing length using a GA
The optimization of the thermal drilling process is
performed using the GA in MATLAB toolbox, with the
intention of enhancing the effectiveness of the drilling
process in order to achieve high structural rigidity for
fasteners in joining situations.35,36 This algorithm makes
a binary coding system to characterize the variables such
as rotational speed (RS), tool angle (TA) and workpiece
thickness (WT). All of the process variables are symbo-
lized by a ten-bit binary equivalent. In the chromosome,
the process variables are represented as a substring. The
GA employs different types of crossover and mutation
operators to predict the maximum values of the bushing
length. At this time, the second-order mathematical
model is considered as an objective function with the
aim of maximizing the output bushing length (BL). In
this model, the rotational speed (RS), tool angle (TA)
and workpiece thickness (WT) are considered as the
input parameters. The aim of the optimization is to
maximize the bushing length; however, usually, a GA is
used to achieve the minimum function value for a mini-
mization problem. Hence, in the present situation, the
objective function was converted into a minimization
problem. A unity negative factor is multiplied to the
objective function in the larger-the-best-type bushing
length (L) of the responses characteristic of thermal
drilling to make them minimize the type objective. The
parameters used in the GA technique are shown in Table
7. Equation (9) is a developed regression model of the
bushing length and it is used as the fitness function or
objective function for this problem. It is developed with
the knowledge gained form ANN values.37,38
Table 7: GA parameters
Population type Double vector
Population size 100
Scaling function Rank
Selection function Roulette wheel
Elite count 2
Crossover fraction 0.8
Mutation function Adaptive feasible
Crossover function Heuristic
Stall generations 50
Maximum bushing length (L) = 6.56203 – 4.64305 ×
10–4 × RS – 0.210894 × TA – 1.04080 × WT +
3.54118 × 10–7 × RS2 – 0.00413827 – TA2 – 0.282222
× WT2 – 5.82394 × 10–5 × RS × TA + 0.000668277 ×
RS × WT + 0.0584444 × TA × WT (9)
is subjected to constrained variables such as:
1944 min–1 = RS = 3600 min–1 (RS – rotational speed)
30° = TA = 45° (TA – tool angle)
1 mm = WT = 2 mm (WT – workpiece thickness).
Figure 9 shows the bushing-length fitness-function
plot obtained with the genetic algorithm. The negative
sign from the final result can be eliminated to get the
maximum value of the bushing length. The optimum
result is achieved at the 52nd iteration of the genetic
algorithm. It is found that the mean fitness value is
5.7472 mm, whereas the best fitness value is 5.7584 mm.
Table 8 presents the optimized and experimental value
of the bushing length. Very close agreement between the
optimized and experimental values of the bushing length
is obtained. It confirms the potential applicability of
these GA techniques for the industry-related problems.
Figure 10 shows the cross-section of the bushing length
of a hole drilled in galvanized steel under the condition
of the optimal process parameter. The bushing length can
offer a longer contact-surface area, which can uphold a
shaft securely, and the drilled-hole inner surface has a
thermally affected zone.
Table 8: Comparison of predicted and experimental values of the
bushing length
Optimum drilling input
process parameters
Output
para-
meters
Rota-
tional
speed
(min–1)
Tool
angle
(°)
Work-
piece
thickness
(mm)
Bushing
length
(mm)
Parameters optimized
with genetic algorithm 3552.461 45 2 5.758
Experimentally used
parameters 3600 45 2 5.6
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
820 Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 9: Fitness-function plot of the bushing length from the genetic
algorithm
Figure 10: Cross-section of a thermally drilled hole
5.3 Microstructural investigation
Figure 11 shows a macrograph obtained with
scanning electron microscopy at a magnification of 17×.
The scanning electron microscope was utilized to
observe and investigate the inside surface of a hole
drilled in galvanized steel. During the thermal-drilling
operation, the rubbing action between the interface of the
thermal drill and the galvanized-steel workpiece led to a
rise in the temperature of up to 798 °C as recorded
on-line during the experimentation. Due to such fric-
tional heating and application of the axial tensile force,
Luders bands were formed to relieve the high amount of
the internal stresses experienced in the galvanized
steel.39,40 Luders bands started to form at the tail end of
the bushing at a distance of 914 μm, after the piercing of
a hole with the center zone of the thermal drill. Until a
distance of 1725 μm, stretcher-strain markings such as
Luders bands are observed in the micrograph shown in
Figure 11. The width of each band thus formed is
different because of non-uniform yielding of the galva-
nized steel during the contact with the cylindrical zone of
the thermal drill.
6 CONCLUSIONS
The present work demonstrates the possibility of
thermal drilling of galvanized steel that has tremendous
applications in the domain of automobile and aerospace
engineering. Owing to its importance, the mechanism
and formation of the bushing length is studied.
Due to frictional heating and applied axial force,
internal stresses tend to increase in that region and
subsequently lead to the formation of Luders bands at
the tail end of the bushing. The formation of these
stretcher-strain marks is due to discontinuous non-uni-
form yielding of the galvanized steel.
The relationship between the input parameters such
as the rotational speed, tool angle and workpiece thick-
ness, and the output parameters like the bushing length is
modeled through an ANN technique. The developed
ANN model is appropriately incorporated with the GA to
optimize the thermal-drilling process parameters.
A good correlation was observed between the
experimental measurements and the predicted optimum
values. This shows that the ANN model combined with
the GA can be successfully applied to find the optimum
conditions for achieving the maximum bushing length in
the thermal drilling of galvanized steel.
The modeling and optimization are valid for one
material and coating only. For different materials, the
building of a new ANN model is required, and the gene-
tic optimization is to be performed again.
Acknowledgements
The authors are grateful for the financial grant by the
Mepco Schlenk Engineering College (Autonomous),
Sivakasi, under the Students Project Scheme (Letter No.
OF/EDC/2840/2015-2016 dated 24.10.2015) for establi-
shing an experimental set-up. The authors acknowledge
the support and encouragement by Dr. S. Arivazhagan,
Principal and Dr. P. Nagaraj, Head of Mechanical
Engineering, towards this work.
7 REFERENCES
1 N. R. J. Hynes, M. V. Maheshwaran, Numerical analysis on thermal
drilling of aluminum metal matrix composite, AIP. Conf. Proc., 1728
(2016), 1–5, doi:10.1063/1.4946597
2 O. Cebeli, D. Zulkuf, Investigate the friction drilling of aluminium
alloys according to the thermal conductivity, Tem. J., 2 (2013) 1,
93–101
3 S. F. Miller, Experimental analysis and numerical modeling of the
friction drilling process, Thesis, University of Michigan, 2006
4 N. R. J. Hynes, M. Muthukumaran, N. Rakesh, C. K. Gurubaran,
Numerical analysis in friction drilling of AISI1020 steel and AA
6061 T6 alloy, Recent Advances in Environmental and Earth
Sciences and Economics, 39 (2015), 145–149
5 A. H. Streppel, H. J. J Kals, Flow drilling: a preliminary analysis of a
new bush-making operation, CIRP Ann. Manuf. Techn., 32 (1983) 1,
167–171, doi:10.1016/S0007-8506(07)63383-6
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822 821
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 11: Microstructure of a thermally drilled hole
5.2 Optimization of the bushing length using a GA
The optimization of the thermal drilling process is
performed using the GA in MATLAB toolbox, with the
intention of enhancing the effectiveness of the drilling
process in order to achieve high structural rigidity for
fasteners in joining situations.35,36 This algorithm makes
a binary coding system to characterize the variables such
as rotational speed (RS), tool angle (TA) and workpiece
thickness (WT). All of the process variables are symbo-
lized by a ten-bit binary equivalent. In the chromosome,
the process variables are represented as a substring. The
GA employs different types of crossover and mutation
operators to predict the maximum values of the bushing
length. At this time, the second-order mathematical
model is considered as an objective function with the
aim of maximizing the output bushing length (BL). In
this model, the rotational speed (RS), tool angle (TA)
and workpiece thickness (WT) are considered as the
input parameters. The aim of the optimization is to
maximize the bushing length; however, usually, a GA is
used to achieve the minimum function value for a mini-
mization problem. Hence, in the present situation, the
objective function was converted into a minimization
problem. A unity negative factor is multiplied to the
objective function in the larger-the-best-type bushing
length (L) of the responses characteristic of thermal
drilling to make them minimize the type objective. The
parameters used in the GA technique are shown in Table
7. Equation (9) is a developed regression model of the
bushing length and it is used as the fitness function or
objective function for this problem. It is developed with
the knowledge gained form ANN values.37,38
Table 7: GA parameters
Population type Double vector
Population size 100
Scaling function Rank
Selection function Roulette wheel
Elite count 2
Crossover fraction 0.8
Mutation function Adaptive feasible
Crossover function Heuristic
Stall generations 50
Maximum bushing length (L) = 6.56203 – 4.64305 ×
10–4 × RS – 0.210894 × TA – 1.04080 × WT +
3.54118 × 10–7 × RS2 – 0.00413827 – TA2 – 0.282222
× WT2 – 5.82394 × 10–5 × RS × TA + 0.000668277 ×
RS × WT + 0.0584444 × TA × WT (9)
is subjected to constrained variables such as:
1944 min–1 = RS = 3600 min–1 (RS – rotational speed)
30° = TA = 45° (TA – tool angle)
1 mm = WT = 2 mm (WT – workpiece thickness).
Figure 9 shows the bushing-length fitness-function
plot obtained with the genetic algorithm. The negative
sign from the final result can be eliminated to get the
maximum value of the bushing length. The optimum
result is achieved at the 52nd iteration of the genetic
algorithm. It is found that the mean fitness value is
5.7472 mm, whereas the best fitness value is 5.7584 mm.
Table 8 presents the optimized and experimental value
of the bushing length. Very close agreement between the
optimized and experimental values of the bushing length
is obtained. It confirms the potential applicability of
these GA techniques for the industry-related problems.
Figure 10 shows the cross-section of the bushing length
of a hole drilled in galvanized steel under the condition
of the optimal process parameter. The bushing length can
offer a longer contact-surface area, which can uphold a
shaft securely, and the drilled-hole inner surface has a
thermally affected zone.
Table 8: Comparison of predicted and experimental values of the
bushing length
Optimum drilling input
process parameters
Output
para-
meters
Rota-
tional
speed
(min–1)
Tool
angle
(°)
Work-
piece
thickness
(mm)
Bushing
length
(mm)
Parameters optimized
with genetic algorithm 3552.461 45 2 5.758
Experimentally used
parameters 3600 45 2 5.6
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
820 Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 9: Fitness-function plot of the bushing length from the genetic
algorithm
Figure 10: Cross-section of a thermally drilled hole
5.3 Microstructural investigation
Figure 11 shows a macrograph obtained with
scanning electron microscopy at a magnification of 17×.
The scanning electron microscope was utilized to
observe and investigate the inside surface of a hole
drilled in galvanized steel. During the thermal-drilling
operation, the rubbing action between the interface of the
thermal drill and the galvanized-steel workpiece led to a
rise in the temperature of up to 798 °C as recorded
on-line during the experimentation. Due to such fric-
tional heating and application of the axial tensile force,
Luders bands were formed to relieve the high amount of
the internal stresses experienced in the galvanized
steel.39,40 Luders bands started to form at the tail end of
the bushing at a distance of 914 μm, after the piercing of
a hole with the center zone of the thermal drill. Until a
distance of 1725 μm, stretcher-strain markings such as
Luders bands are observed in the micrograph shown in
Figure 11. The width of each band thus formed is
different because of non-uniform yielding of the galva-
nized steel during the contact with the cylindrical zone of
the thermal drill.
6 CONCLUSIONS
The present work demonstrates the possibility of
thermal drilling of galvanized steel that has tremendous
applications in the domain of automobile and aerospace
engineering. Owing to its importance, the mechanism
and formation of the bushing length is studied.
Due to frictional heating and applied axial force,
internal stresses tend to increase in that region and
subsequently lead to the formation of Luders bands at
the tail end of the bushing. The formation of these
stretcher-strain marks is due to discontinuous non-uni-
form yielding of the galvanized steel.
The relationship between the input parameters such
as the rotational speed, tool angle and workpiece thick-
ness, and the output parameters like the bushing length is
modeled through an ANN technique. The developed
ANN model is appropriately incorporated with the GA to
optimize the thermal-drilling process parameters.
A good correlation was observed between the
experimental measurements and the predicted optimum
values. This shows that the ANN model combined with
the GA can be successfully applied to find the optimum
conditions for achieving the maximum bushing length in
the thermal drilling of galvanized steel.
The modeling and optimization are valid for one
material and coating only. For different materials, the
building of a new ANN model is required, and the gene-
tic optimization is to be performed again.
Acknowledgements
The authors are grateful for the financial grant by the
Mepco Schlenk Engineering College (Autonomous),
Sivakasi, under the Students Project Scheme (Letter No.
OF/EDC/2840/2015-2016 dated 24.10.2015) for establi-
shing an experimental set-up. The authors acknowledge
the support and encouragement by Dr. S. Arivazhagan,
Principal and Dr. P. Nagaraj, Head of Mechanical
Engineering, towards this work.
7 REFERENCES
1 N. R. J. Hynes, M. V. Maheshwaran, Numerical analysis on thermal
drilling of aluminum metal matrix composite, AIP. Conf. Proc., 1728
(2016), 1–5, doi:10.1063/1.4946597
2 O. Cebeli, D. Zulkuf, Investigate the friction drilling of aluminium
alloys according to the thermal conductivity, Tem. J., 2 (2013) 1,
93–101
3 S. F. Miller, Experimental analysis and numerical modeling of the
friction drilling process, Thesis, University of Michigan, 2006
4 N. R. J. Hynes, M. Muthukumaran, N. Rakesh, C. K. Gurubaran,
Numerical analysis in friction drilling of AISI1020 steel and AA
6061 T6 alloy, Recent Advances in Environmental and Earth
Sciences and Economics, 39 (2015), 145–149
5 A. H. Streppel, H. J. J Kals, Flow drilling: a preliminary analysis of a
new bush-making operation, CIRP Ann. Manuf. Techn., 32 (1983) 1,
167–171, doi:10.1016/S0007-8506(07)63383-6
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822 821
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 11: Microstructure of a thermally drilled hole
6 M. Kerkhofs, M. Van Stappen, The performance of (Ti, Al) N-coated
flow drills, Surf. Coat. Technol., 68 (1994) 69, 741–746
7 S. F. Miller, P. J. Blau, A. J. Shih, Microstructural alterations
associated with friction drilling of steel, aluminum, and titanium, J.
Mater. Eng. Perform., 14 (2005), 647–653, doi:10.1361/
105994905X64558
8 S. F. Miller, R. Li, H. Wang, A. J. Shih, Experimental and numerical
analysis of the friction drilling process, J. Manuf. Sci. Eng., 128
(2006), 802–810, doi:10.1115/1.2193554
9 S. F. Miller, J. Tao, A. J. Shih, Friction drilling of cast metals, Int. J.
Mach. Tools Manuf., 46 (2006), 1526–1535, doi:10.1016/
j.ijmachtools.2005.09.003
10 S. F. Miller, A. J. Shih, Thermo-mechanical finite element modeling
of the friction drilling process, J. Manuf. Sci. Eng., 129 (2007),
531–538
11 S. F. Miller, P. J. Blau, A. J. Shih, Tool wear in friction drilling, Int.
J. Mach. Tools Manuf., 47 (2007), 1636–1645, doi:10.1016/
j.ijmachtools.2006.10.009
12 S. M. Lee, H. M. Chow, B. H. Yan, Friction drilling of IN-713LC
cast superalloy, Mater. Manuf. Proces., 22 (2007), 893–897,
doi:10.1080/10426910701451697
13 H. M. Chow, S. M. Lee, L. D. Yang, Machining characteristic study
of friction drilling on AISI 304 stainless steel, J. Mater. Process.
Technol., 207 (2008), 180–186, doi:10.1016/j.jmatprotec.2007.
12.064
14 B. Mehmet, C. Ibrahim, G. Mustafa, O. Feridun, Application of the
taguchi method to optimize the cutting conditions in hard turning of
a ring bore, Mater. Tech., 49 (2015) 5, 765–772, doi:10.17222/mit.
2014.246
15 S. M. Lee, H. M. Chow, F. Y. Huang, B. H. Yan, Friction drilling of
austenitic stainless steel by uncoated and PVD AlCrN- and
TiAlN-coated tungsten carbide tools, Int. J. Mach. Tools Manuf., 49
(2009), 81–88,doi:10.1016/j.ijmachtools.2008.07.012
16 W. L. Ku, C. L. Hung, S. M. Lee, H. M. Chow, Optimization in ther-
mal friction drilling for SUS 304 stainless steel, Int. J. Adv. Manuf.
Technol., 53 (2011), 935–944, doi:10.1007/s00170-010-2899-5
17 M. Folea, D. Schlegel, E. Gete, C. Langlade, A. Roman, Preliminary
tests on flow drilling of maraging steels, Acad. J. Manuf. Eng. 10
(2012) 4, 42–47
18 T. K. Mehmet, A. Alaattin, B. Bertan, K. A. Hamza, An experimental
study on friction drilling of ST12 steel, T. Can. Soc. Mech. Eng., 38
(2014), 319–329
19 D. Biermann, Y. Liu, Innovative flow drilling on magnesium wrought
alloy AZ31, Proc. CIRP., 18 (2014), 209–214, doi:10.1016/j.procir.
2014.06.133
20 B. B. Mehmet, G. Kadir, G. Arif, Three-dimensional finite element
model of friction drilling process in hot forming processes, J. Proc.
Mech. Eng., (2015), 1–7, doi:10.1177/0954408915614300
21 P. Krasauskas, S. Kilikevi~ius, R. ^esnavi~ius, D. Pa~enga, Expe-
rimental analysis and numerical simulation of the stainless AISI 304
steel friction drilling process, Mecha., 20 (2014) 6, 590–595,
doi:10.5755/j01.mech.20.6.8664
22 P. D. Pantawane, B. B. Ahuja, Parametric analysis and modelling of
friction drilling process on AISI 1015, Int. J. Mecha. Manuf. Sys., 7
(2014) 1, 60–79, doi:10.1504/IJMMS.2014.062771
23 G. Somasundaram, B. S. Rajendra, K. Palanikumar, Modeling and
analysis of roundness error in friction drilling of aluminum silicon
carbide metal matrix composite, J. Compos. Mater., 46 (2011) 2,
169–181, doi:10.1177/0021998311410493
24 S. Changyu, W. Lixia, C. Wei, W. Jinxing, Optimization for injection
molding process conditions of the refrigeratory top cover using com-
bination method of artificial neural network and genetic algorithms,
Poly. Plast. Tech. Eng., 46 (2007), 105–112, doi:10.1080/
03602550601152853
25 S. R. Karnik, V. N. Gaitone, Development of artificial neural network
models to study the effect of process parameters on burr size in
drilling, Int. J. Adv. Manuf. Technol., 39 (2008), 439–453,
doi:10.1007/s00170-007-1231-5
26 R. S. Mamilla, S. Mondal, J. Ramkumar, V. K. Jain, Experimental
investigations and modeling of drill bit-guided abrasive flow
finishing (DBG-AFF) process, Int. J. Adv. Manuf. Technol., 42
(2009), 678–688, doi:10.1007/s00170-008-1642-y
27 S. Assarzadeh, M. Ghoreishi, Neural-network-based modeling and
optimization of the electro-discharge machining process, Int. J. Adv.
Manuf. Technol., 39 (2008), 488–500, doi:10.1007/s00170-007-
1235-1
28 O. Babur, O. Hasan, K. Hasan, Optimum surface roughness in end
milling Inconel 718 by coupling neural network model and genetic
algorithm, Int. J. Adv. Manuf. Technol., 27 (2005), 234–241,
doi:10.1007/s00170-004-2175-7
29 S. Sarkar, S. Mitra, B. Bhattacharyya, Parametric optimisation of
wire electrical discharge machining of ã titanium aluminide alloy
through an artificial neural network model, Int. J. Adv. Manuf.
Technol., 27 (2006), 501–508, doi:10.1007/s00170-004-2203-7
30 A. K. Singh, S. S. Panda, S. K. Pal, D. Chakraborty, Predicting drill
wear using an artificial neural network, Int. J. Adv. Manuf. Technol.,
28 (2006), 456–462, doi:10.1007/s00170-004-2376-0
31 S. M. Hamidinejad, F. Kolahan, A. H. Kokabi, The modeling and
process analysis of resistance spot welding on galvanized steel sheets
used in car body manufacturing, Mater. Des., 34 (2012), 759–767,
doi:10.1016/j.matdes.2011.06.064
32 Y. Lieh-Dai, K. Wei-Liang, C. Han-Ming, W. Der-An, L. Yan-
Cherng, Mar-M247, Haynes-230 & Inconel-718 study of machining
characteristics for Ni-based super alloys on friction drilling, Adv.
Mat. Res., 459 (2012), 632–637, doi:10.4028/www.scientific.net/
AMR.459.632
33 J. B. Peter, C. J. Brian, Q. Jun, Feasibility of thermally drilling
automotive alloy sheet, castings, and hydro formed shapes, (2007),
http://web.ornl.gov/info/reports/2007/3445605662084.pdf, 08/2016
34 P. Sathiya, K. Panneerselvam, M. Y. Abdul Jaleel, Optimization of
laser welding process parameters for super austenitic stainless steel
using artificial neural networks and genetic algorithm, Mater. Des.,
36 (2012), 490–498, doi:10.1016/j.matdes.2011.11.028
35 K. Girish, S. S. Kuldip, Predictive modelling and optimization of
machining parameters to minimize surface roughness using artificial
neural network coupled with genetic algorithm, Proc. CIRP., 31
(2015), 453–458, doi:10.1016/j.procir.2015.03.043
36 S. S. Kuldip, S. Sachin, K. Girish, Optimization of machining
parameters to minimize surface roughness using integrated ANN-GA
approach, Proc. CIRP., 29 (2015), 305–310, doi:10.1016/j.procir.
2015.02.002
37 D. S. Nagesh, G. L. Datta, Genetic algorithm for optimization of
welding variables for height to width ratio and application of ANN
for prediction of bead geometry for TIG welding process, App. Soft.
Comp., 10 (2010), 897–907, doi:10.1016/j.asoc.2009.10.007
38 M. Z. Azlan, H. Habibollah, S. Safian, Genetic algorithm and
simulated annealing to estimate optimal process parameters of the
abrasive water jet machining, Eng. Comp., 27 (2011), 251–259,
doi:10.1007/s00366-010-0195-5
39 V. S. Ananthan, E. O. Hall, Macroscopic aspects of Luders band
deformation in mild steel, Actametall. Mater., 39 (1991) 12,
3153–3160, doi:10.1016/0956-7151(91)90049-7
40 F. H. Julian, K. Stelios, On the effect of Luders bands on the bending
of steel tubes. Part II: Analysis, Int. J. Solids. Struct., 48 (2011),
3285–3298, doi:10.1016/j.ijsolstr.2011.07.012
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
822 Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
823–829
GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX
IN A DYE-SENSITIZED SOLAR CELL
@ELIRNI POLISAHARID KOT ELEKTROLITNA OSNOVA V
SOLARNIH CELICAH, OB^UTLJIVIH NA BARVILA
Jose Paolo Bantang1, Drexel Camacho1,2
1De La Salle University, Chemistry Department, 2401 Taft Avenue, 1004 Manila, Philippines
2De La Salle University, Organic Materials and Interfaces Unit, CENSER, 2401 Taft Avenue, 1004 Manila, Philippines
drexel.camacho@dlsu.edu.ph
Prejem rokopisa – received: 2016-09-30; sprejem za objavo – accepted for publication: 2017-03-16
doi:10.17222/mit.2016.294
Hydrophilic polysaccharide, -carrageenan, was utilized as the polymer matrix in gel-electrolyte systems for dye-sensitized
solar-cell (DSSC) applications. The influence of the solvent system was investigated to optimize the solubility of -carrageenan
and tetrabutylammonium-iodide (TBAI)/I2 electrolytes by minimizing the water content because of its unfavorable effect on
DSSCs. We report herein that two solvent systems, a water/acetonitrile mixed solvent and DMSO, were found to effectively
dissolve the components. The composite natures of the -carrageenan-electrolyte systems in these solvents were confirmed with
an FTIR analysis. The presence of -carrageenan did not impede the electrochemical properties of the electrolytes, as confirmed
with cyclic voltammetry, electrochemical impedance spectroscopy and linear sweep voltammetry. The incorporation of the gel
electrolytes in DSSCs showed that the DMSO system exhibited better solar-cell efficiency compared to the mixed-solvent
system.
Keywords: dye-sensitized solar cell, -carrageenan, gel electrolyte, electrochemical impedance spectroscopy, ionic conductivity
Hidrofilni polisaharid, -karagen, je bil uporabljen kot polimerna osnova v `elirno elektrolitskih sistemih za uporabo v barvno
ob~utljivih son~nih celicah (angl. DSSC). Da bi izbolj{ali topnost -karagena in elektrolitov tetrabutilamonijevega iodida
(TBAI)/I2 z zmanj{anjem vsebnosti vode, zaradi njenega ne`elenega u~inka na DSSC, je bila preiskovana ob~utljivost sistema
topil. ^lanek poro~a, da sta bila najdena dva nova sistema topil, me{anica topila voda/acetonitril in DMSO, ki u~inkovito
raztopita komponente. Lastnosti oz. obna{anje kompozitov elektrolitskega sistema -karagen v teh raztopinah, so bile potrjene z
FTIR-analizo. Prisotnost -karagenskega elektrolitnega sistem ni predstavljala ovire za imepdan~no spektroskopijo in lienarno
"sweep" voltametrijo. Vklju~itev gel-elektrolitov v DSSC je pokazala, da DMSO-sistem ka`e bolj{o solarno celi~no
u~inkovitost kot sistem me{anih topil.
Klju~ne besede: barvno ob~utljive son~ne celice, rastlinska `elatina -karagen, gel-elektrolit, elektrokemi~na impedan~na
spekroskopija, ionska prevodnost
1 INTRODUCTION
Gel-polymer electrolytes are promising materials for
a potential incorporation in electrochemical devices1
such as dye-sensitized solar cells (DSC). This is because
their mechanical behavior is that of solids, yet their inter-
nal structure is flexible and their conductivity behavior
resembles that of a liquid state allowing a good elec-
trode/electrolyte contact.2,3 Moreover, its ease of fabrica-
tion allows more tunability in designing systems for
specific applications.
Natural polysaccharide is a potential matrix for
polymer electrolytes owing to its hydrophilic and
gel-forming capacity that traps the solvent together with
the redox couple inside the polymer matrix. The high
water-retention property of polysaccharide-based poly-
mer-electrolyte systems was shown to promote good
ionic conductivities and thermal stability.4,5 Agarose,6,7
cellulose and its derivatives8,9 and -carrageenan10 are the
gel-forming natural polysaccharides that have been used
as polymer-electrolyte (PE) systems for dye-sensi-
tized-solar-cell (DSSC) applications. To maximize the
water-retention property of these polysaccharide-based
gel systems, an aqueous medium is necessary. However,
we can see that water disturbs the interfacial attachment
of dyes to TiO2. Thus, it is highly desirable to develop a
polysaccharide-based electrolyte system in a smaller
amount of water or in a non-aqueous medium.
-Carrageenan, which is composed mainly of
disaccharide units of -D-galactopyranose with either an
-D-galactopyranose or 3,6-anhydrogalactose is a pro-
mising polysaccharide matrix for electrolytes due to its
hydrophilic, linear and sulfated properties. The electro-
static interactions of ions with hydroxyl groups and
sulfate groups in the main chain are essential to the
conduction mechanism of a polysaccharide electrolyte
system.11,12 The promising features of -carrageenan as a
polymer-electrolyte system can be potentially applied to
DSSCs. Solid-state DSSCs using carrageenan-gel elec-
trolyte systems were fabricated by soaking the aqueous
polymer gel with the redox electrolytes11 and forming
thin polymer membranes13 leading to decent efficiencies.
To date, studies on the preparation and characterization
of carrageenan-electrolyte systems have been limited
Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829 823
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:621.3.035.22:544.623 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)823(2017)
6 M. Kerkhofs, M. Van Stappen, The performance of (Ti, Al) N-coated
flow drills, Surf. Coat. Technol., 68 (1994) 69, 741–746
7 S. F. Miller, P. J. Blau, A. J. Shih, Microstructural alterations
associated with friction drilling of steel, aluminum, and titanium, J.
Mater. Eng. Perform., 14 (2005), 647–653, doi:10.1361/
105994905X64558
8 S. F. Miller, R. Li, H. Wang, A. J. Shih, Experimental and numerical
analysis of the friction drilling process, J. Manuf. Sci. Eng., 128
(2006), 802–810, doi:10.1115/1.2193554
9 S. F. Miller, J. Tao, A. J. Shih, Friction drilling of cast metals, Int. J.
Mach. Tools Manuf., 46 (2006), 1526–1535, doi:10.1016/
j.ijmachtools.2005.09.003
10 S. F. Miller, A. J. Shih, Thermo-mechanical finite element modeling
of the friction drilling process, J. Manuf. Sci. Eng., 129 (2007),
531–538
11 S. F. Miller, P. J. Blau, A. J. Shih, Tool wear in friction drilling, Int.
J. Mach. Tools Manuf., 47 (2007), 1636–1645, doi:10.1016/
j.ijmachtools.2006.10.009
12 S. M. Lee, H. M. Chow, B. H. Yan, Friction drilling of IN-713LC
cast superalloy, Mater. Manuf. Proces., 22 (2007), 893–897,
doi:10.1080/10426910701451697
13 H. M. Chow, S. M. Lee, L. D. Yang, Machining characteristic study
of friction drilling on AISI 304 stainless steel, J. Mater. Process.
Technol., 207 (2008), 180–186, doi:10.1016/j.jmatprotec.2007.
12.064
14 B. Mehmet, C. Ibrahim, G. Mustafa, O. Feridun, Application of the
taguchi method to optimize the cutting conditions in hard turning of
a ring bore, Mater. Tech., 49 (2015) 5, 765–772, doi:10.17222/mit.
2014.246
15 S. M. Lee, H. M. Chow, F. Y. Huang, B. H. Yan, Friction drilling of
austenitic stainless steel by uncoated and PVD AlCrN- and
TiAlN-coated tungsten carbide tools, Int. J. Mach. Tools Manuf., 49
(2009), 81–88,doi:10.1016/j.ijmachtools.2008.07.012
16 W. L. Ku, C. L. Hung, S. M. Lee, H. M. Chow, Optimization in ther-
mal friction drilling for SUS 304 stainless steel, Int. J. Adv. Manuf.
Technol., 53 (2011), 935–944, doi:10.1007/s00170-010-2899-5
17 M. Folea, D. Schlegel, E. Gete, C. Langlade, A. Roman, Preliminary
tests on flow drilling of maraging steels, Acad. J. Manuf. Eng. 10
(2012) 4, 42–47
18 T. K. Mehmet, A. Alaattin, B. Bertan, K. A. Hamza, An experimental
study on friction drilling of ST12 steel, T. Can. Soc. Mech. Eng., 38
(2014), 319–329
19 D. Biermann, Y. Liu, Innovative flow drilling on magnesium wrought
alloy AZ31, Proc. CIRP., 18 (2014), 209–214, doi:10.1016/j.procir.
2014.06.133
20 B. B. Mehmet, G. Kadir, G. Arif, Three-dimensional finite element
model of friction drilling process in hot forming processes, J. Proc.
Mech. Eng., (2015), 1–7, doi:10.1177/0954408915614300
21 P. Krasauskas, S. Kilikevi~ius, R. ^esnavi~ius, D. Pa~enga, Expe-
rimental analysis and numerical simulation of the stainless AISI 304
steel friction drilling process, Mecha., 20 (2014) 6, 590–595,
doi:10.5755/j01.mech.20.6.8664
22 P. D. Pantawane, B. B. Ahuja, Parametric analysis and modelling of
friction drilling process on AISI 1015, Int. J. Mecha. Manuf. Sys., 7
(2014) 1, 60–79, doi:10.1504/IJMMS.2014.062771
23 G. Somasundaram, B. S. Rajendra, K. Palanikumar, Modeling and
analysis of roundness error in friction drilling of aluminum silicon
carbide metal matrix composite, J. Compos. Mater., 46 (2011) 2,
169–181, doi:10.1177/0021998311410493
24 S. Changyu, W. Lixia, C. Wei, W. Jinxing, Optimization for injection
molding process conditions of the refrigeratory top cover using com-
bination method of artificial neural network and genetic algorithms,
Poly. Plast. Tech. Eng., 46 (2007), 105–112, doi:10.1080/
03602550601152853
25 S. R. Karnik, V. N. Gaitone, Development of artificial neural network
models to study the effect of process parameters on burr size in
drilling, Int. J. Adv. Manuf. Technol., 39 (2008), 439–453,
doi:10.1007/s00170-007-1231-5
26 R. S. Mamilla, S. Mondal, J. Ramkumar, V. K. Jain, Experimental
investigations and modeling of drill bit-guided abrasive flow
finishing (DBG-AFF) process, Int. J. Adv. Manuf. Technol., 42
(2009), 678–688, doi:10.1007/s00170-008-1642-y
27 S. Assarzadeh, M. Ghoreishi, Neural-network-based modeling and
optimization of the electro-discharge machining process, Int. J. Adv.
Manuf. Technol., 39 (2008), 488–500, doi:10.1007/s00170-007-
1235-1
28 O. Babur, O. Hasan, K. Hasan, Optimum surface roughness in end
milling Inconel 718 by coupling neural network model and genetic
algorithm, Int. J. Adv. Manuf. Technol., 27 (2005), 234–241,
doi:10.1007/s00170-004-2175-7
29 S. Sarkar, S. Mitra, B. Bhattacharyya, Parametric optimisation of
wire electrical discharge machining of ã titanium aluminide alloy
through an artificial neural network model, Int. J. Adv. Manuf.
Technol., 27 (2006), 501–508, doi:10.1007/s00170-004-2203-7
30 A. K. Singh, S. S. Panda, S. K. Pal, D. Chakraborty, Predicting drill
wear using an artificial neural network, Int. J. Adv. Manuf. Technol.,
28 (2006), 456–462, doi:10.1007/s00170-004-2376-0
31 S. M. Hamidinejad, F. Kolahan, A. H. Kokabi, The modeling and
process analysis of resistance spot welding on galvanized steel sheets
used in car body manufacturing, Mater. Des., 34 (2012), 759–767,
doi:10.1016/j.matdes.2011.06.064
32 Y. Lieh-Dai, K. Wei-Liang, C. Han-Ming, W. Der-An, L. Yan-
Cherng, Mar-M247, Haynes-230 & Inconel-718 study of machining
characteristics for Ni-based super alloys on friction drilling, Adv.
Mat. Res., 459 (2012), 632–637, doi:10.4028/www.scientific.net/
AMR.459.632
33 J. B. Peter, C. J. Brian, Q. Jun, Feasibility of thermally drilling
automotive alloy sheet, castings, and hydro formed shapes, (2007),
http://web.ornl.gov/info/reports/2007/3445605662084.pdf, 08/2016
34 P. Sathiya, K. Panneerselvam, M. Y. Abdul Jaleel, Optimization of
laser welding process parameters for super austenitic stainless steel
using artificial neural networks and genetic algorithm, Mater. Des.,
36 (2012), 490–498, doi:10.1016/j.matdes.2011.11.028
35 K. Girish, S. S. Kuldip, Predictive modelling and optimization of
machining parameters to minimize surface roughness using artificial
neural network coupled with genetic algorithm, Proc. CIRP., 31
(2015), 453–458, doi:10.1016/j.procir.2015.03.043
36 S. S. Kuldip, S. Sachin, K. Girish, Optimization of machining
parameters to minimize surface roughness using integrated ANN-GA
approach, Proc. CIRP., 29 (2015), 305–310, doi:10.1016/j.procir.
2015.02.002
37 D. S. Nagesh, G. L. Datta, Genetic algorithm for optimization of
welding variables for height to width ratio and application of ANN
for prediction of bead geometry for TIG welding process, App. Soft.
Comp., 10 (2010), 897–907, doi:10.1016/j.asoc.2009.10.007
38 M. Z. Azlan, H. Habibollah, S. Safian, Genetic algorithm and
simulated annealing to estimate optimal process parameters of the
abrasive water jet machining, Eng. Comp., 27 (2011), 251–259,
doi:10.1007/s00366-010-0195-5
39 V. S. Ananthan, E. O. Hall, Macroscopic aspects of Luders band
deformation in mild steel, Actametall. Mater., 39 (1991) 12,
3153–3160, doi:10.1016/0956-7151(91)90049-7
40 F. H. Julian, K. Stelios, On the effect of Luders bands on the bending
of steel tubes. Part II: Analysis, Int. J. Solids. Struct., 48 (2011),
3285–3298, doi:10.1016/j.ijsolstr.2011.07.012
N. R. J. HYNES et al.: OPTIMUM BUSHING LENGTH IN THERMAL DRILLING OF GALVANIZED STEEL ...
822 Materiali in tehnologije / Materials and technology 51 (2017) 5, 813–822
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
823–829
GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX
IN A DYE-SENSITIZED SOLAR CELL
@ELIRNI POLISAHARID KOT ELEKTROLITNA OSNOVA V
SOLARNIH CELICAH, OB^UTLJIVIH NA BARVILA
Jose Paolo Bantang1, Drexel Camacho1,2
1De La Salle University, Chemistry Department, 2401 Taft Avenue, 1004 Manila, Philippines
2De La Salle University, Organic Materials and Interfaces Unit, CENSER, 2401 Taft Avenue, 1004 Manila, Philippines
drexel.camacho@dlsu.edu.ph
Prejem rokopisa – received: 2016-09-30; sprejem za objavo – accepted for publication: 2017-03-16
doi:10.17222/mit.2016.294
Hydrophilic polysaccharide, -carrageenan, was utilized as the polymer matrix in gel-electrolyte systems for dye-sensitized
solar-cell (DSSC) applications. The influence of the solvent system was investigated to optimize the solubility of -carrageenan
and tetrabutylammonium-iodide (TBAI)/I2 electrolytes by minimizing the water content because of its unfavorable effect on
DSSCs. We report herein that two solvent systems, a water/acetonitrile mixed solvent and DMSO, were found to effectively
dissolve the components. The composite natures of the -carrageenan-electrolyte systems in these solvents were confirmed with
an FTIR analysis. The presence of -carrageenan did not impede the electrochemical properties of the electrolytes, as confirmed
with cyclic voltammetry, electrochemical impedance spectroscopy and linear sweep voltammetry. The incorporation of the gel
electrolytes in DSSCs showed that the DMSO system exhibited better solar-cell efficiency compared to the mixed-solvent
system.
Keywords: dye-sensitized solar cell, -carrageenan, gel electrolyte, electrochemical impedance spectroscopy, ionic conductivity
Hidrofilni polisaharid, -karagen, je bil uporabljen kot polimerna osnova v `elirno elektrolitskih sistemih za uporabo v barvno
ob~utljivih son~nih celicah (angl. DSSC). Da bi izbolj{ali topnost -karagena in elektrolitov tetrabutilamonijevega iodida
(TBAI)/I2 z zmanj{anjem vsebnosti vode, zaradi njenega ne`elenega u~inka na DSSC, je bila preiskovana ob~utljivost sistema
topil. ^lanek poro~a, da sta bila najdena dva nova sistema topil, me{anica topila voda/acetonitril in DMSO, ki u~inkovito
raztopita komponente. Lastnosti oz. obna{anje kompozitov elektrolitskega sistema -karagen v teh raztopinah, so bile potrjene z
FTIR-analizo. Prisotnost -karagenskega elektrolitnega sistem ni predstavljala ovire za imepdan~no spektroskopijo in lienarno
"sweep" voltametrijo. Vklju~itev gel-elektrolitov v DSSC je pokazala, da DMSO-sistem ka`e bolj{o solarno celi~no
u~inkovitost kot sistem me{anih topil.
Klju~ne besede: barvno ob~utljive son~ne celice, rastlinska `elatina -karagen, gel-elektrolit, elektrokemi~na impedan~na
spekroskopija, ionska prevodnost
1 INTRODUCTION
Gel-polymer electrolytes are promising materials for
a potential incorporation in electrochemical devices1
such as dye-sensitized solar cells (DSC). This is because
their mechanical behavior is that of solids, yet their inter-
nal structure is flexible and their conductivity behavior
resembles that of a liquid state allowing a good elec-
trode/electrolyte contact.2,3 Moreover, its ease of fabrica-
tion allows more tunability in designing systems for
specific applications.
Natural polysaccharide is a potential matrix for
polymer electrolytes owing to its hydrophilic and
gel-forming capacity that traps the solvent together with
the redox couple inside the polymer matrix. The high
water-retention property of polysaccharide-based poly-
mer-electrolyte systems was shown to promote good
ionic conductivities and thermal stability.4,5 Agarose,6,7
cellulose and its derivatives8,9 and -carrageenan10 are the
gel-forming natural polysaccharides that have been used
as polymer-electrolyte (PE) systems for dye-sensi-
tized-solar-cell (DSSC) applications. To maximize the
water-retention property of these polysaccharide-based
gel systems, an aqueous medium is necessary. However,
we can see that water disturbs the interfacial attachment
of dyes to TiO2. Thus, it is highly desirable to develop a
polysaccharide-based electrolyte system in a smaller
amount of water or in a non-aqueous medium.
-Carrageenan, which is composed mainly of
disaccharide units of -D-galactopyranose with either an
-D-galactopyranose or 3,6-anhydrogalactose is a pro-
mising polysaccharide matrix for electrolytes due to its
hydrophilic, linear and sulfated properties. The electro-
static interactions of ions with hydroxyl groups and
sulfate groups in the main chain are essential to the
conduction mechanism of a polysaccharide electrolyte
system.11,12 The promising features of -carrageenan as a
polymer-electrolyte system can be potentially applied to
DSSCs. Solid-state DSSCs using carrageenan-gel elec-
trolyte systems were fabricated by soaking the aqueous
polymer gel with the redox electrolytes11 and forming
thin polymer membranes13 leading to decent efficiencies.
To date, studies on the preparation and characterization
of carrageenan-electrolyte systems have been limited
Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829 823
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:621.3.035.22:544.623 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)823(2017)
only to pure aqueous systems. A significant amount of
water present in the electrolyte systems applied to
DSSCs causes a decrease in the photocurrent of the
device associated with dye desorption, iodate formation
and a decrease in the electron lifetime.14 Hence, the
amount of water must be controlled, if not eliminated. It
is the objective of this paper to investigate the influence
of the solvent system on carrageenan electrolyte and its
performance in a DSSC. This paper reports, for the first
time, on a novel carrageenan-based electrolyte system
with an iodide/tri-iodide redox couple prepared using a
mixed solvent system and a non-aqueous solvent.
2 MATERIALS AND METHODS
2.1 Optimization of the -carrageenan electrolyte sys-
tem
Different -carrageenan (Shemberg, Philippines)
biopolymer gel electrolytes were optimized by varying
the ratio of water/acetonitrile and the amount of polymer
with a constant amount of tetrabutylammonium iodide
(TBAI, Sigma Aldrich, a reagent grade of 98 %; 20 %
w/w of TBAI with respect to -carrageenan) and iodine
(I2, Aldrich, = 99.99 % metal basis). The molar ratio of
TBAI/I2 was 4:1. The ratio of water/ACN was optimized
by dissolving -carrageenan (0.2 g) with TBAI/I2 in
different solvent systems consisting of 100 % water, 3:1
water/ACN, 1:1 water/ACN, 1:3 water/ACN, and 100 %
ACN at 80 °C until a homogeneous solution was formed.
Thin films were formed by casting solutions on plastic
petri dishes and slowly drying in air. Different biopoly-
mer electrolytes were prepared using the same procedure
by varying the concentration of -carrageenan (0, 0.5,
1.0, 1.5 and 2 % w/v). Using the optimized parameters,
-carrageenan gel electrolytes containing 20 % w/w
(with respect to the amount of -carrageenan) of other
salts such as potassium iodide (Sigma Aldrich) and
trimethylsulfonium iodide (TMSI, Sigma Aldrich) were
also prepared.
A liquid-state polymer electrolyte and a gel-state
polymer electrolyte based on -carrageenan were pre-
pared by dissolving 1 % w/v or 2 % w/v of the poly-
saccharide in 2 mL of dimethylsulfoxide at 70 °C.
Tetrabutylammonium iodide (0.5 M), I2 (0.05 M) and LiI
(0.1 M) were added to the -carrageenan solution. As the
control, the same composition was used without adding
the biopolymer.
2.2 Fabrication of dye-sensitized solar cells (DSSCs)
The photoanode was prepared spreading a TiO2 paste
(Solaronix, Ti-Nanoxide T/SP) on a fluorine-doped
tin-oxide (FTO) conducting glass (TCO22-7) with the
doctor-blading method. The active area of the photo-
anode was fixed to 1 cm × 1 cm with adhesive tapes. The
deposited paste was sintered at 450 °C for 30 min and
cooled down slowly to room temperature. The deposited
TiO2 film was soaked in a dye solution containing a
ruthenium dye (Ruthenizer 535-bisTBA, Solaronix) in
ethanol for 24 h. The counter electrode was prepared by
sintering an FTO conducting glass coated with a pla-
tinum solution (Platisol 41121, Solaronix) at 450 °C for
30 mins. The photoanode and counter electrode were
sealed with a hot-melt thermoplastic material called
Surlyn (DuPont, thickness of 60 μm), also acting as a
spacer. A dye-sensitized solar cell was fabricated by
inserting a hot gel-electrolyte solution in a pre-drilled
hole on the counter electrode. Three different solar cells
were fabricated containing different iodide salts (KI,
TMSI, and TBAI) in a water/ACN -carrageenan matrix.
A control DSSC was prepared using an acetonitrile-
based liquid electrolyte (Iodolyte AN50, Solaronix). The
same process was used for the fabrication of a DSSC
incorporating gel electrolytes in DMSO.
2.3 Characterization
The surface morphology and thickness of the thin
films were characterized using scanning electron micro-
scopy (JEOL JSM-5310). The electrical conductivity
was measured using the 4-point probe Van der Pauw
method. In brief, four wires were attached to the corners
of a 1 cm × 1 cm thin film using a silver paste. The
voltage at a constant current was measured. The
resistivity and electrical conductivity were calculated
using the following equations:
Resistivity, s
d V
l
V
l
f
R
R
= +
⎛
⎝
⎜
⎞
⎠
⎟
⎛
⎝
⎜
⎞
⎠
⎟
π
ln 2
1
2
2
2
1
2
and Electrical
conductivity,  = 1/s, where d, V, I and f R /R( )1 2 refer to
the thickness, potential, current and van der Pauw
function, respectively.
The redox properties of the -carrageenan gel
electrolyte were characterized using 3-electrode cyclic
voltammetry (Powerlab/4SP potentiostat). The electrodes
composed of the Ag/AgCl reference electrode, plati-
num-wire counter electrode and glassy-carbon working
electrode were pierced through the gel. Electrochemical
impedance spectroscopy (Metrohm Autolab potentiostat
PGSTAT128N) was obtained by sandwiching the gel
electrolyte in between two platinum-coated FTO
conducting-glass electrodes. The active area was fixed to
1 × 1 cm2 and the thickness of the spacer was about
50 μm. The frequency was set from 106 Hz to 10–1 Hz
with an AC applied voltage of 10 mV. The ionic
conductivity () of the gel electrolyte was calculated
from the solution resistance (Rs) using Equation (3):
 =
L
ARs
(3)
The solution resistance (Rs) is determined as the
intercept of the Nyquist plot against the real part of the
impedance or the horizontal axis. The variables L and A
correspond to the thickness of the spacer and the active
area, respectively. The solar-cell parameters such as the
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
824 Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
open-circuit voltage, short-circuit current, fill factor and
efficiency of the fabricated dye-sensitized solar cells
with -carrageenan gel electrolyte were assessed using a
100 mW/cm2 AM 1.5 solar simulator (Abet Technolo-
gies) coupled with a potentiostat to obtain a current-
voltage (IV) curve.
3 RESULTS AND DISCUSSIONS
3.1 Carrageenan-electrolyte system in a mixed solvent
system
Acetonitrile (ACN), which is a common organic
solvent used for preparing liquid electrolytes for DSSC
applications, does not dissolve -carrageenan. Thus, to
form -carrageenan electrolytes containing a redox
couple I–/I3– (from TBAI and I2), the amount of water is
crucial to dissolve the hydrophilic polysaccharide.15
Using water only as the solvent, a black film was formed
with observable colorless crystals (Figure 1a) indicating
the unsuccessful dissolution of both TBAI salt and
iodine. Likewise, crystallization of the salt was formed
on the surface of the film prepared using a 3:1 v/v
water/ACN system (Figure 1b). The salt crystallization
is attributed to the relative insolubility of the salt in a
water-rich solvent. Moreover, the presence of the salt and
iodine crystals on the film surface indicates that it fails to
dissociate in the polymer matrix, leading to an un-
successful incorporation of the charge-carrier redox
couple I–/I3–. For a 1:1 v/v water/ACN mixed-solvent
system, a uniform orange polymer-electrolyte film with-
out visible crystals was formed (Figure 1c), indicating
good dissolution and incorporation of the I–/I3– redox
couple in the polymer matrix. No films were cast for the
1:3 v/v water/acetonitrile solvent systems and in 100 %
acetonitrile because of the relative insolubility of
carrageenan in an acetonitrile-rich solvent.
The concentration of carrageenan in the polymer
electrolyte containing 20 % w/w TBAI (w.r.t. -carra-
geenan) in the 1:1 H2O/acetonitrile mixed solvent affects
the property and the ionic conductivity of the electrolyte.
A liquid polymer electrolyte is formed when the con-
centrations of carrageenan are in a range of 0.1–0.5 %
w/v. Increasing the carrageenan concentration to  1.0 %
w/v gave a gel polymer electrolyte. Consequently, the
ionic conductivity of the polymer electrolytes changes in
response to its composite form. The ionic conductivity
(Figure 2) of -carrageenan-electrolyte systems in-
creases as the carrageenan concentration increases from
0.1 to 1.0 % w/v. The increase in the ionic conductivity at
low -carrageenan concentrations is attributed to high
free-ion concentrations due to their interaction with the
functional group present in -carrageenan.16. The ionic
conductivity decreases at carrageenan concentrations of
1.0–2.0 % w/v. The increase in the polymer concentra-
tion causes chain entanglements, impeding the move-
ment of the ions, thus, decreasing its ionic conductivity.16
Different iodide salts were incorporated in
-carrageenan to prepare the gel electrolytes. Films were
formed and the IR spectra of the electrolyte films con-
tained typical functional-group vibrations of -carra-
geenan.17 The bands at 1272 cm–1 were identified as an
O=S=O symmetric vibration that would respond in the
presence of a cation. A shift in the frequency was
observed in the O=S=O symmetric vibration compared
with the pure k-carrageenan reference (Table 1, entry 1).
This shift is attributed to the electrostatic interaction
between the negatively charged sulfate ions in the
carrageenan chains and the positively charged cations
from the iodide salt.4 The effect of the addition of iodide
salts generally increased the electrical conductivity
(Table 1, entry 2) of the films formed, as compared with
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829 825
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Effect of -carrageenan concentration on ionic conductivity
of carrageenan/TBAI/I2 systems
Figure 1: Optical images of 1 % w/v -carrageenan films mixed with
40 % w/w TBAI (w.r.t. -carrageenan) in 4:1 mole ratio TBAI /I2
electrolyte in varying solvents a) in 100 % water, b) in 3:1
H2O/acetonitrile, c) in 1:1 H2O/acetonitrile, d) polymer gel
electrolytes containing -carrageenan and TBAI/LiI/I2 electrolyte:
left: 1 % w/v -carrageenan in DMSO (liquid state), right: 2 % w/v
-carrageenan in DMSO (gel-state)
only to pure aqueous systems. A significant amount of
water present in the electrolyte systems applied to
DSSCs causes a decrease in the photocurrent of the
device associated with dye desorption, iodate formation
and a decrease in the electron lifetime.14 Hence, the
amount of water must be controlled, if not eliminated. It
is the objective of this paper to investigate the influence
of the solvent system on carrageenan electrolyte and its
performance in a DSSC. This paper reports, for the first
time, on a novel carrageenan-based electrolyte system
with an iodide/tri-iodide redox couple prepared using a
mixed solvent system and a non-aqueous solvent.
2 MATERIALS AND METHODS
2.1 Optimization of the -carrageenan electrolyte sys-
tem
Different -carrageenan (Shemberg, Philippines)
biopolymer gel electrolytes were optimized by varying
the ratio of water/acetonitrile and the amount of polymer
with a constant amount of tetrabutylammonium iodide
(TBAI, Sigma Aldrich, a reagent grade of 98 %; 20 %
w/w of TBAI with respect to -carrageenan) and iodine
(I2, Aldrich, = 99.99 % metal basis). The molar ratio of
TBAI/I2 was 4:1. The ratio of water/ACN was optimized
by dissolving -carrageenan (0.2 g) with TBAI/I2 in
different solvent systems consisting of 100 % water, 3:1
water/ACN, 1:1 water/ACN, 1:3 water/ACN, and 100 %
ACN at 80 °C until a homogeneous solution was formed.
Thin films were formed by casting solutions on plastic
petri dishes and slowly drying in air. Different biopoly-
mer electrolytes were prepared using the same procedure
by varying the concentration of -carrageenan (0, 0.5,
1.0, 1.5 and 2 % w/v). Using the optimized parameters,
-carrageenan gel electrolytes containing 20 % w/w
(with respect to the amount of -carrageenan) of other
salts such as potassium iodide (Sigma Aldrich) and
trimethylsulfonium iodide (TMSI, Sigma Aldrich) were
also prepared.
A liquid-state polymer electrolyte and a gel-state
polymer electrolyte based on -carrageenan were pre-
pared by dissolving 1 % w/v or 2 % w/v of the poly-
saccharide in 2 mL of dimethylsulfoxide at 70 °C.
Tetrabutylammonium iodide (0.5 M), I2 (0.05 M) and LiI
(0.1 M) were added to the -carrageenan solution. As the
control, the same composition was used without adding
the biopolymer.
2.2 Fabrication of dye-sensitized solar cells (DSSCs)
The photoanode was prepared spreading a TiO2 paste
(Solaronix, Ti-Nanoxide T/SP) on a fluorine-doped
tin-oxide (FTO) conducting glass (TCO22-7) with the
doctor-blading method. The active area of the photo-
anode was fixed to 1 cm × 1 cm with adhesive tapes. The
deposited paste was sintered at 450 °C for 30 min and
cooled down slowly to room temperature. The deposited
TiO2 film was soaked in a dye solution containing a
ruthenium dye (Ruthenizer 535-bisTBA, Solaronix) in
ethanol for 24 h. The counter electrode was prepared by
sintering an FTO conducting glass coated with a pla-
tinum solution (Platisol 41121, Solaronix) at 450 °C for
30 mins. The photoanode and counter electrode were
sealed with a hot-melt thermoplastic material called
Surlyn (DuPont, thickness of 60 μm), also acting as a
spacer. A dye-sensitized solar cell was fabricated by
inserting a hot gel-electrolyte solution in a pre-drilled
hole on the counter electrode. Three different solar cells
were fabricated containing different iodide salts (KI,
TMSI, and TBAI) in a water/ACN -carrageenan matrix.
A control DSSC was prepared using an acetonitrile-
based liquid electrolyte (Iodolyte AN50, Solaronix). The
same process was used for the fabrication of a DSSC
incorporating gel electrolytes in DMSO.
2.3 Characterization
The surface morphology and thickness of the thin
films were characterized using scanning electron micro-
scopy (JEOL JSM-5310). The electrical conductivity
was measured using the 4-point probe Van der Pauw
method. In brief, four wires were attached to the corners
of a 1 cm × 1 cm thin film using a silver paste. The
voltage at a constant current was measured. The
resistivity and electrical conductivity were calculated
using the following equations:
Resistivity, s
d V
l
V
l
f
R
R
= +
⎛
⎝
⎜
⎞
⎠
⎟
⎛
⎝
⎜
⎞
⎠
⎟
π
ln 2
1
2
2
2
1
2
and Electrical
conductivity,  = 1/s, where d, V, I and f R /R( )1 2 refer to
the thickness, potential, current and van der Pauw
function, respectively.
The redox properties of the -carrageenan gel
electrolyte were characterized using 3-electrode cyclic
voltammetry (Powerlab/4SP potentiostat). The electrodes
composed of the Ag/AgCl reference electrode, plati-
num-wire counter electrode and glassy-carbon working
electrode were pierced through the gel. Electrochemical
impedance spectroscopy (Metrohm Autolab potentiostat
PGSTAT128N) was obtained by sandwiching the gel
electrolyte in between two platinum-coated FTO
conducting-glass electrodes. The active area was fixed to
1 × 1 cm2 and the thickness of the spacer was about
50 μm. The frequency was set from 106 Hz to 10–1 Hz
with an AC applied voltage of 10 mV. The ionic
conductivity () of the gel electrolyte was calculated
from the solution resistance (Rs) using Equation (3):
 =
L
ARs
(3)
The solution resistance (Rs) is determined as the
intercept of the Nyquist plot against the real part of the
impedance or the horizontal axis. The variables L and A
correspond to the thickness of the spacer and the active
area, respectively. The solar-cell parameters such as the
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
824 Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
open-circuit voltage, short-circuit current, fill factor and
efficiency of the fabricated dye-sensitized solar cells
with -carrageenan gel electrolyte were assessed using a
100 mW/cm2 AM 1.5 solar simulator (Abet Technolo-
gies) coupled with a potentiostat to obtain a current-
voltage (IV) curve.
3 RESULTS AND DISCUSSIONS
3.1 Carrageenan-electrolyte system in a mixed solvent
system
Acetonitrile (ACN), which is a common organic
solvent used for preparing liquid electrolytes for DSSC
applications, does not dissolve -carrageenan. Thus, to
form -carrageenan electrolytes containing a redox
couple I–/I3– (from TBAI and I2), the amount of water is
crucial to dissolve the hydrophilic polysaccharide.15
Using water only as the solvent, a black film was formed
with observable colorless crystals (Figure 1a) indicating
the unsuccessful dissolution of both TBAI salt and
iodine. Likewise, crystallization of the salt was formed
on the surface of the film prepared using a 3:1 v/v
water/ACN system (Figure 1b). The salt crystallization
is attributed to the relative insolubility of the salt in a
water-rich solvent. Moreover, the presence of the salt and
iodine crystals on the film surface indicates that it fails to
dissociate in the polymer matrix, leading to an un-
successful incorporation of the charge-carrier redox
couple I–/I3–. For a 1:1 v/v water/ACN mixed-solvent
system, a uniform orange polymer-electrolyte film with-
out visible crystals was formed (Figure 1c), indicating
good dissolution and incorporation of the I–/I3– redox
couple in the polymer matrix. No films were cast for the
1:3 v/v water/acetonitrile solvent systems and in 100 %
acetonitrile because of the relative insolubility of
carrageenan in an acetonitrile-rich solvent.
The concentration of carrageenan in the polymer
electrolyte containing 20 % w/w TBAI (w.r.t. -carra-
geenan) in the 1:1 H2O/acetonitrile mixed solvent affects
the property and the ionic conductivity of the electrolyte.
A liquid polymer electrolyte is formed when the con-
centrations of carrageenan are in a range of 0.1–0.5 %
w/v. Increasing the carrageenan concentration to  1.0 %
w/v gave a gel polymer electrolyte. Consequently, the
ionic conductivity of the polymer electrolytes changes in
response to its composite form. The ionic conductivity
(Figure 2) of -carrageenan-electrolyte systems in-
creases as the carrageenan concentration increases from
0.1 to 1.0 % w/v. The increase in the ionic conductivity at
low -carrageenan concentrations is attributed to high
free-ion concentrations due to their interaction with the
functional group present in -carrageenan.16. The ionic
conductivity decreases at carrageenan concentrations of
1.0–2.0 % w/v. The increase in the polymer concentra-
tion causes chain entanglements, impeding the move-
ment of the ions, thus, decreasing its ionic conductivity.16
Different iodide salts were incorporated in
-carrageenan to prepare the gel electrolytes. Films were
formed and the IR spectra of the electrolyte films con-
tained typical functional-group vibrations of -carra-
geenan.17 The bands at 1272 cm–1 were identified as an
O=S=O symmetric vibration that would respond in the
presence of a cation. A shift in the frequency was
observed in the O=S=O symmetric vibration compared
with the pure k-carrageenan reference (Table 1, entry 1).
This shift is attributed to the electrostatic interaction
between the negatively charged sulfate ions in the
carrageenan chains and the positively charged cations
from the iodide salt.4 The effect of the addition of iodide
salts generally increased the electrical conductivity
(Table 1, entry 2) of the films formed, as compared with
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829 825
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Effect of -carrageenan concentration on ionic conductivity
of carrageenan/TBAI/I2 systems
Figure 1: Optical images of 1 % w/v -carrageenan films mixed with
40 % w/w TBAI (w.r.t. -carrageenan) in 4:1 mole ratio TBAI /I2
electrolyte in varying solvents a) in 100 % water, b) in 3:1
H2O/acetonitrile, c) in 1:1 H2O/acetonitrile, d) polymer gel
electrolytes containing -carrageenan and TBAI/LiI/I2 electrolyte:
left: 1 % w/v -carrageenan in DMSO (liquid state), right: 2 % w/v
-carrageenan in DMSO (gel-state)
the reference pure -carrageenan due to the addition and
incorporation of the redox species. The ionic conduc-
tivities (Table 1, entry 3) of the gel electrolytes are
generally higher than that of the acetonitrile-based liquid
electrolyte, indicating a facile transport of the redox
couple in the carrageenan matrix. Cyclic voltammetry of
the gel electrolytes with and without -carrageenan
(Figure 3) showed that the addition of the polymer does
not impede the electrochemical behavior of the I–/I3–
redox couple in the mixed water/acetonitrile system. Two
redox processes were observed, indicating the migration
of I–/I3– in the 3D matrix of -carrageenan.18
Solar cells were fabricated using the conventional
set-up composed of a TiO2/ruthenium dye as the photo-
anode and a platinum counter electrode sandwiching the
-carrageenan-gel electrolyte systems in water/ACN. I-V
characteristics of the DSSCs (Table 2 and Figure 4a)
containing the -carrageenan-gel electrolytes showed
poor efficiency (<0.10 %) as compared to the liquid
electrolyte (2.33 %), attributed to a much lower short-
circuit current, which can be associated with the pre-
sence of a significant amount of water in the electrolyte
system. The presence of water causes dye detachment
due to the weakening of the dye-TiO2 linkage, resulting
in a limited amount of electrons injected from the excited
state of the dye and consequently leading to the lowering
of the short-circuit current.16 The obtained efficiencies of
different -carrageenan-gel electrolytes with a mixture of
acetonitrile and water suggest that imparting a large
amount of water (50 %) in the electrolyte system for
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
826 Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: I-V curves of the DSSCs: a) containing different iodide salts
(KI, TMSI and TBAI) in the water/ACN -carrageenan matrix and
b) containing -carrageenan-based electrolytes in DMSO under a 100
mW/cm2 illumination
Table 2: Characteristics of the fabricated DSSCs under a 100
mW/cm2 illumination
Electrolyte Voc (V) Isc(mA/cm2) FF  ( %)
Iodolyte
(control) 0.695 6.444 0.521 2.333
KI:I2 0.746 0.100 0.819 0.061
TMSI:I2 0.776 0.0414 0.152 0.005
TBAI:I2 0.778 0.0817 0.765 0.049
Figure 3: Electrochemical-behavior (cyclic voltametry) analysis of
the -carrageenan-gel electrolytes containing different salts; control:
LiI/I2 in 1:1 v/v water/ACN without carrageenan
Table 1: FTIR and conductivity of the -carrageenan electrolytes containing different salts
Entry Pure - carrageenan(reference)
Iodolyte AN50
(control)
KI/I2 TMSI/I2 TBAI/I2
in -carrageenan
1 FTIR O=S=O vibration (cm–1) 1272 – 1261 1274 1263
2
Electrical conductivity of
polymer-electrolyte filmsa
(S cm–1)
0.127 – 0.281 0.189 0.326
3
Ionic conductivity of
polymer-electrolyte gelsb
(x10–4 S cm–1)
– 1.66 2.52 3.05 2.17
aMeasured with 4-point probe Van der Pauw technique
bMeasured with electrochemical impedance spectroscopy (EIS)
DSSCs is still detrimental to the performance of a solar-
cell device.
3.2 Carrageenan-electrolyte system in a non-aqueous
solvent
Carrageenan is not known to dissolve in purely non-
aqueous solvents. However, in the course of our
investigation of electrolyte systems, we observed that it
dissolves in dimethylsulfoxide (DMSO).19,20 We investi-
gated the solubility properties of -carrageenan in
DMSO, dissolving different amounts of the biopolymer
at 60 °C. A liquid solution is formed with 1 % w/v
-carrageenan (a solubility of 0.01 g/mL DMSO at room
temperature and at 60 °C). Increasing the -carrageenan
concentration from 1.5 to 3.5 % w/v allows the dissolu-
tion of the biopolymer in DMSO, forming a transparent
solution at 60 °C. However, upon cooling down at room
temperature, it forms into a stable gel (Figure 1d). A
further increase in the concentration of -carrageenan of
4–5 % w/v leads to the formation of a stable gel even at
60 °C. Other organic solvents such as ethanol, methanol,
acetonitrile, tetrahydrofuran, dimethylformamide and
methoxyethanol did not show good solubility. The
advantages of using DMSO are its relatively low vapor
pressure and relatively non-volatile nature, which are
useful for maintaining its gel state. Moreover, this orga-
nic solvent shows a relatively low toxicity and is con-
sidered environmentally benign.6 The addition of electro-
lyte systems TBAI and I2 turned the -carrageenan
DMSO solution into a dark-brown one indicating the
incorporation of the electrolytes. The IR spectrum of the
polymer-electrolyte system showed functional groups
typical of -carrageenan. A shift in the O=S=O symme-
tric vibration from typical 1272 for pure -carrageenan
to 1226 cm–1 for carrageenan gel electrolyte is attributed
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829 827
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: EIS spectrum of a DSSC containing -carrageenan-based
electrolytes in DMSO; inset: equivalent circuit model (R1 = series
resistance, R2 = TiO2/dye/electrolyte interface resistance, CPE =
constant phase element)
Figure 5: a) EIS spectra of electrolytes measured using a symme-
trical-cell set-up between two platinized FTO conducting glasses in a
frequency range of 1 MHz to 0.1 Hz, b) linear sweep voltammogram
of the electrolytes measured between two platinized FTO conducting
glasses swept between –1 V to 1 V. Scan rate: 10 mV/s
Table 3: I-V electrochemical properties of the carrageenan electrolyte in DMSO measured using two platinized FTO conducting glasses and I-V
characteristics of DSSCs under a 100 mW/cm2 illumination*
Electrochemical properties of electrolyte I-V characteristics of DSSC
Electrolytes in
DMSO Rb
** (Ù) RPt
(Ù)
Ionic con-
ductivity,
×10–4 (S/cm)
Diffusivity of
I3–, ×10–9
(cm2 s–1)
Short-circuit
current
(mA/cm2)
Open-circuit
voltage (V) Fill factor
Efficiency
(%)
Control (without
-carrageenan) 21.28 218.13 2.82 1.48 0.690 0.781 0.410 0.221
1% w/v
-carrageenan 25.33 296.41 2.37 1.41 0.960 0.763 0.483 0.354
2 % w/v
-carrageenan 25.25 338.46 2.38 1.11 0.983 0.751 0.488 0.360
*The electrolyte used was composed of tetrabutylammonium iodide (0.5 M), I2, (0.05 M) and LiI (0.1 M)
**Rb corresponds to the solution resistance, which corresponds to the intercept at the real part of the EIS spectrum
the reference pure -carrageenan due to the addition and
incorporation of the redox species. The ionic conduc-
tivities (Table 1, entry 3) of the gel electrolytes are
generally higher than that of the acetonitrile-based liquid
electrolyte, indicating a facile transport of the redox
couple in the carrageenan matrix. Cyclic voltammetry of
the gel electrolytes with and without -carrageenan
(Figure 3) showed that the addition of the polymer does
not impede the electrochemical behavior of the I–/I3–
redox couple in the mixed water/acetonitrile system. Two
redox processes were observed, indicating the migration
of I–/I3– in the 3D matrix of -carrageenan.18
Solar cells were fabricated using the conventional
set-up composed of a TiO2/ruthenium dye as the photo-
anode and a platinum counter electrode sandwiching the
-carrageenan-gel electrolyte systems in water/ACN. I-V
characteristics of the DSSCs (Table 2 and Figure 4a)
containing the -carrageenan-gel electrolytes showed
poor efficiency (<0.10 %) as compared to the liquid
electrolyte (2.33 %), attributed to a much lower short-
circuit current, which can be associated with the pre-
sence of a significant amount of water in the electrolyte
system. The presence of water causes dye detachment
due to the weakening of the dye-TiO2 linkage, resulting
in a limited amount of electrons injected from the excited
state of the dye and consequently leading to the lowering
of the short-circuit current.16 The obtained efficiencies of
different -carrageenan-gel electrolytes with a mixture of
acetonitrile and water suggest that imparting a large
amount of water (50 %) in the electrolyte system for
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
826 Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: I-V curves of the DSSCs: a) containing different iodide salts
(KI, TMSI and TBAI) in the water/ACN -carrageenan matrix and
b) containing -carrageenan-based electrolytes in DMSO under a 100
mW/cm2 illumination
Table 2: Characteristics of the fabricated DSSCs under a 100
mW/cm2 illumination
Electrolyte Voc (V) Isc(mA/cm2) FF  ( %)
Iodolyte
(control) 0.695 6.444 0.521 2.333
KI:I2 0.746 0.100 0.819 0.061
TMSI:I2 0.776 0.0414 0.152 0.005
TBAI:I2 0.778 0.0817 0.765 0.049
Figure 3: Electrochemical-behavior (cyclic voltametry) analysis of
the -carrageenan-gel electrolytes containing different salts; control:
LiI/I2 in 1:1 v/v water/ACN without carrageenan
Table 1: FTIR and conductivity of the -carrageenan electrolytes containing different salts
Entry Pure - carrageenan(reference)
Iodolyte AN50
(control)
KI/I2 TMSI/I2 TBAI/I2
in -carrageenan
1 FTIR O=S=O vibration (cm–1) 1272 – 1261 1274 1263
2
Electrical conductivity of
polymer-electrolyte filmsa
(S cm–1)
0.127 – 0.281 0.189 0.326
3
Ionic conductivity of
polymer-electrolyte gelsb
(x10–4 S cm–1)
– 1.66 2.52 3.05 2.17
aMeasured with 4-point probe Van der Pauw technique
bMeasured with electrochemical impedance spectroscopy (EIS)
DSSCs is still detrimental to the performance of a solar-
cell device.
3.2 Carrageenan-electrolyte system in a non-aqueous
solvent
Carrageenan is not known to dissolve in purely non-
aqueous solvents. However, in the course of our
investigation of electrolyte systems, we observed that it
dissolves in dimethylsulfoxide (DMSO).19,20 We investi-
gated the solubility properties of -carrageenan in
DMSO, dissolving different amounts of the biopolymer
at 60 °C. A liquid solution is formed with 1 % w/v
-carrageenan (a solubility of 0.01 g/mL DMSO at room
temperature and at 60 °C). Increasing the -carrageenan
concentration from 1.5 to 3.5 % w/v allows the dissolu-
tion of the biopolymer in DMSO, forming a transparent
solution at 60 °C. However, upon cooling down at room
temperature, it forms into a stable gel (Figure 1d). A
further increase in the concentration of -carrageenan of
4–5 % w/v leads to the formation of a stable gel even at
60 °C. Other organic solvents such as ethanol, methanol,
acetonitrile, tetrahydrofuran, dimethylformamide and
methoxyethanol did not show good solubility. The
advantages of using DMSO are its relatively low vapor
pressure and relatively non-volatile nature, which are
useful for maintaining its gel state. Moreover, this orga-
nic solvent shows a relatively low toxicity and is con-
sidered environmentally benign.6 The addition of electro-
lyte systems TBAI and I2 turned the -carrageenan
DMSO solution into a dark-brown one indicating the
incorporation of the electrolytes. The IR spectrum of the
polymer-electrolyte system showed functional groups
typical of -carrageenan. A shift in the O=S=O symme-
tric vibration from typical 1272 for pure -carrageenan
to 1226 cm–1 for carrageenan gel electrolyte is attributed
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829 827
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: EIS spectrum of a DSSC containing -carrageenan-based
electrolytes in DMSO; inset: equivalent circuit model (R1 = series
resistance, R2 = TiO2/dye/electrolyte interface resistance, CPE =
constant phase element)
Figure 5: a) EIS spectra of electrolytes measured using a symme-
trical-cell set-up between two platinized FTO conducting glasses in a
frequency range of 1 MHz to 0.1 Hz, b) linear sweep voltammogram
of the electrolytes measured between two platinized FTO conducting
glasses swept between –1 V to 1 V. Scan rate: 10 mV/s
Table 3: I-V electrochemical properties of the carrageenan electrolyte in DMSO measured using two platinized FTO conducting glasses and I-V
characteristics of DSSCs under a 100 mW/cm2 illumination*
Electrochemical properties of electrolyte I-V characteristics of DSSC
Electrolytes in
DMSO Rb
** (Ù) RPt
(Ù)
Ionic con-
ductivity,
×10–4 (S/cm)
Diffusivity of
I3–, ×10–9
(cm2 s–1)
Short-circuit
current
(mA/cm2)
Open-circuit
voltage (V) Fill factor
Efficiency
(%)
Control (without
-carrageenan) 21.28 218.13 2.82 1.48 0.690 0.781 0.410 0.221
1% w/v
-carrageenan 25.33 296.41 2.37 1.41 0.960 0.763 0.483 0.354
2 % w/v
-carrageenan 25.25 338.46 2.38 1.11 0.983 0.751 0.488 0.360
*The electrolyte used was composed of tetrabutylammonium iodide (0.5 M), I2, (0.05 M) and LiI (0.1 M)
**Rb corresponds to the solution resistance, which corresponds to the intercept at the real part of the EIS spectrum
to the interaction of the cation with the sulfate functional
groups.5
The ionic conductivity and diffusivity of the
tri-iodide ion (Table 3) in DMSO showed that the
addition of -carrageenan decreases the ionic conduc-
tivity of the electrolyte system in comparison to the
DMSO-based liquid-electrolyte system. EIS spectra
(Figure 5a) of the electrolytes in DMSO show a broader
semi-circle arc upon the addition of -carrageenan,
which is associated with an increase in the
charge-transfer resistance at the platinum/electrolyte
interface (RPt). A higher RPt indicates a less efficient
electro-activity at the Pt/electrolyte interface attributed to
a slower ionic conductivity.21 The concentration of iodide
ions is in excess in comparison with the amount of
tri-iodide ions. Hence, the limiting current density of the
gel electrolytes is associated with the diffusion coeffi-
cient of the tri-iodide ions. The limiting current densities
of the electrolytes (Figure 5b) show negative and
positive values due to the transport of charge carriers
from one electrode to the other through the whole elec-
trolyte system.8 A slow diffusion of the tri-iodide ions is
consistent with the slow ionic conductivity of the
carrageenan-electrolyte system. The slow tri-iodide ion
diffusion in the -carrageenan electrolytes resulted in a
higher recombination.22 These results were consistent
with the decreasing trend in the open-circuit voltage of
the electrolytes with -carrageenan. The photovoltaic
performance of the -carrageenan-based dye-sensitized
solar cells (Figure 4b and Table 3) showed that the
addition of -carrageenan to the I–/I3– redox electrolyte
system in dimethylsulfoxide decreases the open-circuit
voltage and increases the short-circuit current of a solar
cell. The improvement in the short-circuit current of the
DSSCs with a -carrageenan/DMSO electrolyte system
is attributed to the ability of the hydroxyl functional
group and negatively charged sulfate ions of -carra-
geenan to allow dissociation of the ions in the polymer
matrix through the electrostatic interaction with the
cations. The enhancement in the short-circuit current of
the -carrageenan/DMSO electrolytes is the major
contribution to the improvement of the overall efficiency
of the dye-sensitized solar cells. Good dissociation of
ions in the polymer matrix increases the concentration of
I–, its counterion, and I3– leading to an increase in the
short-circuit current.23
A typical liquid-electrolyte-based DSSC show an EIS
spectrum consisting of three semi-circle arcs corres-
ponding to the interfacial resistances occurring during
the electrochemical reaction at the platinum/electrolyte
interface, during the charge-transfer process at the
TiO2/dye/electrolyte interface and Warburg diffusion of
I–/I3– in the electrolyte. A frequency-response-analyzer
(FRA) impedance analysis of the -carrageenan/DMSO
DSSC (Figure 6) showed a unique single semi-circle
arc, associated with the charge-recombination process
occurring at the TiO2/dye/electrolyte interface (R2). The
other parameters such as R1 and CPE corresponded to
the series resistance and constant phase element, respec-
tively. A decreasing trend in the semi-circle arc radius
was observed when the -carrageenan concentration was
increased. A quantitative approach of the recombination
resistance was done by fitting the Nyquist plot with its
corresponding equivalent circuit model. Chi-square
values for the fitted equivalent circuit model range from
0.02–0.07, indicating good value fitting. The charge-
recombination resistances (R2) of the DSSCs without
-carrageenan (control), with 1 % w/v -carrageenan
(liquid state) and 2 % w/v -carrageenan (gel state),
obtained from the plot are 515.19 , 317.88  and
292.86 , respectively. A higher R2 value signifies lower
charge recombination between the photoanode conduc-
tion-band electrons and I3– ions at the TiO2/dye/elec-
trolyte interface resulting in a higher Voc.24 The higher
charge recombination of the polymer electrolytes with
-carrageenan, compared to the electrolyte without the
polymer, is generally attributed to the slow ionic con-
ductivity of the ions in the gel matrix of carrageenan
(Table 3). Moreover, the -carrageenan electrolyte pene-
trates poorly into the TiO2 layer due to its polymeric
nature contributing to the less efficient charge-transfer
process.21 The results were consistent with the photo-
voltaic characteristics of the fabricated dye-sensitized
solar cells.
4 CONCLUSIONS
A hydrophilic polysaccharide, -carrageenan, was
successfully used as the polymer matrix for a gel-elec-
trolyte system in a mixed solvent of water and aceto-
nitrile for dye-sensitized solar cells. The solvent system
improved the dissolution of the components and the
electrolyte properties. Efficiencies of less than 0.1 % of
the fabricated DSSCs are attributed to the significant
amount of water. A novel gel-polymer electrolyte con-
sisting of -carrageenan electrolyte gel in DMSO was
developed for DSSCs. Removal of the water from the
polymer-electrolyte system improved the DSSC per-
formance from 0.049 to 0.36 %, a seven-fold increase in
the magnitude. The presence of -carrageenan in the
DMSO-based gel-electrolyte system helps improve the
solar-cell efficiency from 0.221 to 0.360 %.
Acknowledgments
The authors are grateful to Dr. Erwin Enriquez, Dr.
Arnel Salvador and Dr. Armando Somintac for the use of
the solar simulator and Ms. Anna San Esteban for
guidance in solar measurements. This work was
supported by research grants from DOST-PCIEERD of
the Republic of the Philippines; DOST-SEI-ASTHRDP
and DLSU-PhD Fellowship Grants to J. P. O Bantang;
DLSU Research Faculty Grant, DLSU Science
Foundation and DLSU-URCO to D. H. Camacho.
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
828 Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
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J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829 829
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
to the interaction of the cation with the sulfate functional
groups.5
The ionic conductivity and diffusivity of the
tri-iodide ion (Table 3) in DMSO showed that the
addition of -carrageenan decreases the ionic conduc-
tivity of the electrolyte system in comparison to the
DMSO-based liquid-electrolyte system. EIS spectra
(Figure 5a) of the electrolytes in DMSO show a broader
semi-circle arc upon the addition of -carrageenan,
which is associated with an increase in the
charge-transfer resistance at the platinum/electrolyte
interface (RPt). A higher RPt indicates a less efficient
electro-activity at the Pt/electrolyte interface attributed to
a slower ionic conductivity.21 The concentration of iodide
ions is in excess in comparison with the amount of
tri-iodide ions. Hence, the limiting current density of the
gel electrolytes is associated with the diffusion coeffi-
cient of the tri-iodide ions. The limiting current densities
of the electrolytes (Figure 5b) show negative and
positive values due to the transport of charge carriers
from one electrode to the other through the whole elec-
trolyte system.8 A slow diffusion of the tri-iodide ions is
consistent with the slow ionic conductivity of the
carrageenan-electrolyte system. The slow tri-iodide ion
diffusion in the -carrageenan electrolytes resulted in a
higher recombination.22 These results were consistent
with the decreasing trend in the open-circuit voltage of
the electrolytes with -carrageenan. The photovoltaic
performance of the -carrageenan-based dye-sensitized
solar cells (Figure 4b and Table 3) showed that the
addition of -carrageenan to the I–/I3– redox electrolyte
system in dimethylsulfoxide decreases the open-circuit
voltage and increases the short-circuit current of a solar
cell. The improvement in the short-circuit current of the
DSSCs with a -carrageenan/DMSO electrolyte system
is attributed to the ability of the hydroxyl functional
group and negatively charged sulfate ions of -carra-
geenan to allow dissociation of the ions in the polymer
matrix through the electrostatic interaction with the
cations. The enhancement in the short-circuit current of
the -carrageenan/DMSO electrolytes is the major
contribution to the improvement of the overall efficiency
of the dye-sensitized solar cells. Good dissociation of
ions in the polymer matrix increases the concentration of
I–, its counterion, and I3– leading to an increase in the
short-circuit current.23
A typical liquid-electrolyte-based DSSC show an EIS
spectrum consisting of three semi-circle arcs corres-
ponding to the interfacial resistances occurring during
the electrochemical reaction at the platinum/electrolyte
interface, during the charge-transfer process at the
TiO2/dye/electrolyte interface and Warburg diffusion of
I–/I3– in the electrolyte. A frequency-response-analyzer
(FRA) impedance analysis of the -carrageenan/DMSO
DSSC (Figure 6) showed a unique single semi-circle
arc, associated with the charge-recombination process
occurring at the TiO2/dye/electrolyte interface (R2). The
other parameters such as R1 and CPE corresponded to
the series resistance and constant phase element, respec-
tively. A decreasing trend in the semi-circle arc radius
was observed when the -carrageenan concentration was
increased. A quantitative approach of the recombination
resistance was done by fitting the Nyquist plot with its
corresponding equivalent circuit model. Chi-square
values for the fitted equivalent circuit model range from
0.02–0.07, indicating good value fitting. The charge-
recombination resistances (R2) of the DSSCs without
-carrageenan (control), with 1 % w/v -carrageenan
(liquid state) and 2 % w/v -carrageenan (gel state),
obtained from the plot are 515.19 , 317.88  and
292.86 , respectively. A higher R2 value signifies lower
charge recombination between the photoanode conduc-
tion-band electrons and I3– ions at the TiO2/dye/elec-
trolyte interface resulting in a higher Voc.24 The higher
charge recombination of the polymer electrolytes with
-carrageenan, compared to the electrolyte without the
polymer, is generally attributed to the slow ionic con-
ductivity of the ions in the gel matrix of carrageenan
(Table 3). Moreover, the -carrageenan electrolyte pene-
trates poorly into the TiO2 layer due to its polymeric
nature contributing to the less efficient charge-transfer
process.21 The results were consistent with the photo-
voltaic characteristics of the fabricated dye-sensitized
solar cells.
4 CONCLUSIONS
A hydrophilic polysaccharide, -carrageenan, was
successfully used as the polymer matrix for a gel-elec-
trolyte system in a mixed solvent of water and aceto-
nitrile for dye-sensitized solar cells. The solvent system
improved the dissolution of the components and the
electrolyte properties. Efficiencies of less than 0.1 % of
the fabricated DSSCs are attributed to the significant
amount of water. A novel gel-polymer electrolyte con-
sisting of -carrageenan electrolyte gel in DMSO was
developed for DSSCs. Removal of the water from the
polymer-electrolyte system improved the DSSC per-
formance from 0.049 to 0.36 %, a seven-fold increase in
the magnitude. The presence of -carrageenan in the
DMSO-based gel-electrolyte system helps improve the
solar-cell efficiency from 0.221 to 0.360 %.
Acknowledgments
The authors are grateful to Dr. Erwin Enriquez, Dr.
Arnel Salvador and Dr. Armando Somintac for the use of
the solar simulator and Ms. Anna San Esteban for
guidance in solar measurements. This work was
supported by research grants from DOST-PCIEERD of
the Republic of the Philippines; DOST-SEI-ASTHRDP
and DLSU-PhD Fellowship Grants to J. P. O Bantang;
DLSU Research Faculty Grant, DLSU Science
Foundation and DLSU-URCO to D. H. Camacho.
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
828 Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
5 REFERENCES
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Ionic Liquid Composite: Development of Polysaccharide-Based
Solid Electrolyte System, The Manila J. of Science, 6 (2011), 8–15
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water in bis-benzimidazole-derivative electrolyte additives to the
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18 S. Yuan, Q. Tang, B. He, P. Yang, Efficient quasi-solid-state dye-sen-
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19 S. Chan, J. P. Bantang, D. Camacho, Influence of Nanomaterial
Fillers in Biopolymer Electrolyte System for Squaraine-Based
Dye-Sensitized Solar Cells, Int. J. Electrochem. Sci., 10 (2015),
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20 J. P. Bantang, D. Camacho, A Novel Biopolymer Gel Electrolyte
System for DSSC Applications, Proceedings of the DLSU Research
Congress, 3, 2015
21 S. Venkatesan, N. Obadja, T.-W. Chang, L.-T. Chen, Y.-L. Lee,
Performance improvement of gel- and solid-state dye-sensitized solar
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doi:10.1016/j.jpowsour.2014.06.016
22 H.-L. Hsu, C.-F. Tien, Y.-T. Yang, J. Leu, Dye-sensitized solar cells
based on agarose gel electrolytes using allylimidazolium iodides and
environmentally benign solvents, Electrochim. Acta, 91 (2013),
208–213, doi:10.1016/j.electacta.2012.12.133
23 K. Hara, T. Horiguchi, T. Kinoshita, K. Sayama, H. Arakawa,
Influence of electrolytes on the photovoltaic performance of organic
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Sol. Cells, 70 (2001), 151–161, doi:10.1016/S0927-0248(01)00019-8
24 Y. Yang, J. Cui, P. Yi, X. Zheng, X. Guo, W. Wang, Effects of
nanoparticle additives on the properties of agarose polymer
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j.jpowsour.2013.10.016
J. P. BANTANG, D. CAMACHO: GELLING POLYSACCHARIDE AS THE ELECTROLYTE MATRIX ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 823–829 829
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
831–836
DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING
THE MECHANICAL PROPERTIES AND STRESS CORROSION
RESISTANCE OF 7000 Al ALLOYS
RAZVOJ TOPLOTNE OBDELAVE ZA IZBOLJ[ANJE MEHANSKIH
LASTNOSTI IN NAPETOSTNO KOROZIJSKO ODPORNOST 7000
Al ZLITIN
Mehdi Shakouri1, Mohammad Esmailian1, Saeed Shabestari2
1Advanced Material and Renewable Energy Department, Iranian Research Organization for Science and Technology (IROST),
P. O. Box 3353-5111, Tehran, Iran
2Center of Excellence for High Strength Alloys Technology (CEHSAT), School of Metallurgy and Materials Engineering,
Iran University of Science and Technology (IUST), Narmak, 16846, Tehran, Iran
mshakoory@yahoo.com
Prejem rokopisa – received: 2016-10-10; sprejem za objavo – accepted for publication: 2017-04-19
doi:10.17222/mit.2016.297
A retrogression and re-ageing (RRA) treatment is a three-step heat treatment that can improve both the mechanical strength and
the corrosion resistance in aluminum alloys. In this work, the mechanical and stress corrosion properties under various ageing
treatment conditions were investigated in an Al–8.5Zn–2.1Mg–2Cu–0.2Ag (w/%) alloy. The treatments were the T6
conventional method followed by a retrogression and re-ageing (RRA) treatment. Tensile test, scanning electron microscopy
(SEM), energy-dispersive X-ray spectroscopy (EDS) and differential scanning calorimetry (DSC) were used to investigate the
mechanical and stress corrosion cracking (SCC) properties. The results showed that both, the strength and the corrosion
resistance criteria (SCR), which is defined as the ratio between the remaining strength percent in stressed and un-stressed
conditions, improve after the RRA treatment. The tensile strength and SCR criteria for the T6 heat treatment were 547 MPa and
71 % initially and then increased to 612 MPa and 95 % for the RRA treatment, respectively. Moreover, the EDS results showed
that in grain-boundary precipitates the Cu concentration is much higher for the RRA in comparison with the T6 treatment, but it
is lower for the matrix precipitates in the RRA treatment.
Keywords: aluminium alloys, stress corrosion resistance, precipitation, retrogression and re-ageing
Obdelava z retrogresijo in ponovno o`ivitvijo (angl. RRA) je tristopenjska toplotna obdelava, s katero lahko izbolj{amo tako
mehansko trdnost kot odpornosti proti koroziji pri aluminijevih zlitinah. V pri~ujo~em delu so bile raziskovane mehanske
lastnosti in lastnosti odpornosti proti koroziji pri zlitini Al–8.5Zn–2.1Mg–2Cu–0.2Ag (w/%) pod razli~nimi pogoji obdelave
staranja. Izvedeni so bili testi po konvencionalni T6 metodi, z upo{tevanjem RRA obdelave. Natezni preiskus, vrsti~na
elektronska mikroskopija (SEM), energijsko disperzijska rentgenska spektroskopija (EDS) in DSC-testiranja so bili izvedeni, da
bi raziskali mehanske in napetostno-korozijske lomne lastnosti (angl. SCC). Rezultati so pokazali, da se oba kriterija tako
trdnost kot odpornost proti koroziji, ki definirata razmerje med procentom zaostale trdnosti pri napetostnih in nenapetostnih
pogojih, po RRA-obdelavi izbolj{ata.
Klju~ne besede: aluminijeve zlitine, odpornost proti koroziji, oborine, retrogresija in ponovno staranje
1 INTRODUCTION
The Al–Zn–Mg–Cu series aluminum alloys are preci-
pitation-hardening alloys that are used extensively for
light-weight structural applications, in particular in the
aircraft industry. They have a combination of good
strength and good stress corrosion resistance.1
The most popular heat treatment to gain the peak-
aged strength (T6X temper) is a solution treatment,
quenching, stretching (for stress relief purposes)
followed by artificial aging at 120 °C for 24 h, but this
temper is highly susceptible to SCC. In order to reduce
this susceptibility, an over-aging treatment (T7X temper)
is needed.2 This requirement becomes increasingly de-
manding as the solute contents are increased in commer-
cial alloys in order to improve the mechanical properties
even further.
To improve both the mechanical strength and corro-
sion resistance, it has been proposed to utilize an RRA
treatment. This three-step heat treatment has been shown
to offer a stress corrosion resistance as good as that of a
T7X heat treatment, while keeping a strength compar-
able to that of a T6X temper.3,4 This type of heat treat-
ment comprises three steps. First, an ageing step that
leads to a T6 state. A second step (called retrogression or
reversion) of short duration at high temperature dissolves
part of the initially formed precipitates. Third heat
treatment step at lower temperature leads to the desired
microstructure.5
Stress corrosion cracking occurs under loading in a
corrosive environment. Several investigations have re-
ported that the SCC mechanism involves anodic dissolu-
tion, hydrogen-induced cracking, passive film rupture,
hydrogen embrittlement, magnesium segregation to the
Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836 831
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 621.78:67.017:620.193:669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)831(2017)
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
831–836
DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING
THE MECHANICAL PROPERTIES AND STRESS CORROSION
RESISTANCE OF 7000 Al ALLOYS
RAZVOJ TOPLOTNE OBDELAVE ZA IZBOLJ[ANJE MEHANSKIH
LASTNOSTI IN NAPETOSTNO KOROZIJSKO ODPORNOST 7000
Al ZLITIN
Mehdi Shakouri1, Mohammad Esmailian1, Saeed Shabestari2
1Advanced Material and Renewable Energy Department, Iranian Research Organization for Science and Technology (IROST),
P. O. Box 3353-5111, Tehran, Iran
2Center of Excellence for High Strength Alloys Technology (CEHSAT), School of Metallurgy and Materials Engineering,
Iran University of Science and Technology (IUST), Narmak, 16846, Tehran, Iran
mshakoory@yahoo.com
Prejem rokopisa – received: 2016-10-10; sprejem za objavo – accepted for publication: 2017-04-19
doi:10.17222/mit.2016.297
A retrogression and re-ageing (RRA) treatment is a three-step heat treatment that can improve both the mechanical strength and
the corrosion resistance in aluminum alloys. In this work, the mechanical and stress corrosion properties under various ageing
treatment conditions were investigated in an Al–8.5Zn–2.1Mg–2Cu–0.2Ag (w/%) alloy. The treatments were the T6
conventional method followed by a retrogression and re-ageing (RRA) treatment. Tensile test, scanning electron microscopy
(SEM), energy-dispersive X-ray spectroscopy (EDS) and differential scanning calorimetry (DSC) were used to investigate the
mechanical and stress corrosion cracking (SCC) properties. The results showed that both, the strength and the corrosion
resistance criteria (SCR), which is defined as the ratio between the remaining strength percent in stressed and un-stressed
conditions, improve after the RRA treatment. The tensile strength and SCR criteria for the T6 heat treatment were 547 MPa and
71 % initially and then increased to 612 MPa and 95 % for the RRA treatment, respectively. Moreover, the EDS results showed
that in grain-boundary precipitates the Cu concentration is much higher for the RRA in comparison with the T6 treatment, but it
is lower for the matrix precipitates in the RRA treatment.
Keywords: aluminium alloys, stress corrosion resistance, precipitation, retrogression and re-ageing
Obdelava z retrogresijo in ponovno o`ivitvijo (angl. RRA) je tristopenjska toplotna obdelava, s katero lahko izbolj{amo tako
mehansko trdnost kot odpornosti proti koroziji pri aluminijevih zlitinah. V pri~ujo~em delu so bile raziskovane mehanske
lastnosti in lastnosti odpornosti proti koroziji pri zlitini Al–8.5Zn–2.1Mg–2Cu–0.2Ag (w/%) pod razli~nimi pogoji obdelave
staranja. Izvedeni so bili testi po konvencionalni T6 metodi, z upo{tevanjem RRA obdelave. Natezni preiskus, vrsti~na
elektronska mikroskopija (SEM), energijsko disperzijska rentgenska spektroskopija (EDS) in DSC-testiranja so bili izvedeni, da
bi raziskali mehanske in napetostno-korozijske lomne lastnosti (angl. SCC). Rezultati so pokazali, da se oba kriterija tako
trdnost kot odpornost proti koroziji, ki definirata razmerje med procentom zaostale trdnosti pri napetostnih in nenapetostnih
pogojih, po RRA-obdelavi izbolj{ata.
Klju~ne besede: aluminijeve zlitine, odpornost proti koroziji, oborine, retrogresija in ponovno staranje
1 INTRODUCTION
The Al–Zn–Mg–Cu series aluminum alloys are preci-
pitation-hardening alloys that are used extensively for
light-weight structural applications, in particular in the
aircraft industry. They have a combination of good
strength and good stress corrosion resistance.1
The most popular heat treatment to gain the peak-
aged strength (T6X temper) is a solution treatment,
quenching, stretching (for stress relief purposes)
followed by artificial aging at 120 °C for 24 h, but this
temper is highly susceptible to SCC. In order to reduce
this susceptibility, an over-aging treatment (T7X temper)
is needed.2 This requirement becomes increasingly de-
manding as the solute contents are increased in commer-
cial alloys in order to improve the mechanical properties
even further.
To improve both the mechanical strength and corro-
sion resistance, it has been proposed to utilize an RRA
treatment. This three-step heat treatment has been shown
to offer a stress corrosion resistance as good as that of a
T7X heat treatment, while keeping a strength compar-
able to that of a T6X temper.3,4 This type of heat treat-
ment comprises three steps. First, an ageing step that
leads to a T6 state. A second step (called retrogression or
reversion) of short duration at high temperature dissolves
part of the initially formed precipitates. Third heat
treatment step at lower temperature leads to the desired
microstructure.5
Stress corrosion cracking occurs under loading in a
corrosive environment. Several investigations have re-
ported that the SCC mechanism involves anodic dissolu-
tion, hydrogen-induced cracking, passive film rupture,
hydrogen embrittlement, magnesium segregation to the
Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836 831
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 621.78:67.017:620.193:669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)831(2017)
grain boundaries and a precipitate-free zone (PFZ) along
the grain boundary.6–10 However, the microstructural
characteristics of Al-Zn-Mg-Cu high strength aluminum
alloys are well known to have a strong influence not only
on the mechanical properties but also on the SCC sus-
ceptibility. Larger grain-boundary precipitates can trap
more atomic hydrogen to nucleate hydrogen bubbles,
thereby decreasing the hydrogen concentration at grain
boundaries below a critical value is considered to prevent
intergranular SCC fracture.11 Furthermore, the cathodic
grain-boundary precipitates grow by depleting solute
atoms. Studies showed that this leads to the broadening
of the anodic PFZ, which contained no strengthen pre-
cipitation phase and as a result was soft and weak.12 The
combination of tensile stress and anodic dissolution
caused SCC.
SCC resistance is of practical importance for the
industrial applications of the Al–Zn–Mg–Cu series
aluminum alloys. Various heat treatments in these alloys
offer very different SCC properties. The SCR criteria are
proposed to compare the SCC resistance.
The aim of this paper is to investigate the mechanical
strength and SCC of a 7000 series aluminum alloy after
conventional T6 and RRA treatments. The SCR criteria
have been defined to easily compare the SCC resistance
of this alloy after various treatments. SEM investigations
were used to determine the effect of the heat treatments
on the alloy’s microstructure with the aim of studying its
effect on SCC susceptibility. For comparison, the stress
corrosion resistance of the alloy with different heat
treatments was studied by the breaking load method
according to ASTM G139 standard.13
2 EXPERIMENTAL PART
The samples used in this study were received as
8-mm-thick sheets that were homogenized, hot rolled
and heat treated after being alloyed and casted. The
chemical composition of the alloy is shown in Table 1.
Table 1: Chemical composition (in mass fractions, w/%) of fabricated
alloy
Alloy
No. Zn Cu Mg Fe Si Zr Ag
1 8.5 2 2.1 0.18 0.16 0.20 0.19
An induction melting furnace used for melting and
the melt was poured in a water-cooled copper mold. The
as-cast specimens were homogenized at 460 °C for 24 h
and hot rolled to about 33 % reduction. Hot rolling was
performed with 40 min–1 rolling speed at 430 °C. The
specimens were milled to tensile samples according to
the ASTM E8M standard. A schematic showing the pre-
paration of tensile samples from the specimens is given
in Figure 1.
A solution heat treatment was done at 471 °C for 6 h
followed by a water quenching. T6 ageing was per-
formed for 24 h at 120 °C (T6 temper). The scheme of
the retrogression and re-ageing treatment is shown in
Figure 2.
Mechanical properties measurements were made at
ambient temperature on the specimens machined
according to the ASTM E8M-04 small size standard.
The average of three tests was used for each result. The
test’s strain rate was 10–3 /s.
The SEM analysis was performed on a TESCAN
scanning electron microscope. Thermal analysis was per-
formed in a DSC 1 Mettler Toledo differential scanning
calorimeter. Polished alloy disks with a diameter of
5 mm and 0.6 mm thick were sealed in aluminum pans
and heated in a flowing argon atmosphere at a constant
heating rate of 10 °C/min. The flowing rate of the argon
was 100 mL/min.
The stress corrosion tests were performed according
to the ASTM G139 standard. The Neutral 3.5 % Sodium
Chloride Solution was prepared in accordance with the
requirements of the ASTM G44 standard. The ASTM
G49 standard was used for preparation of direct tension
stress corrosion test specimens. A 207 MPa stress for the
cycle of 4 d was applied as per section 8-2 of the ASTM
G139 standard. In order to eliminate the corrosion other
than SCC (such as pitting), some specimens were tested
unstressed. The ratio of remaining strength of stressed
and unstressed samples represented in percent, were
calculated as SCR criteria for different treatment.
3 RESULTS AND DISCUSSION
3.1 Mechanical strength
The mechanical strength of the alloy was 547 MPa
after the T6 treatment and was 612 MPa after the RRA
treatment. The RRA treatment leads to an about 12 %
increase in mechanical strength. While T6 treatment is
carried out to ensure maximum strength; the RRA
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
832 Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Scheme of the RRA treatment
Figure 1: Schematic showing the preparation of tensile samples from
the specimens
treated samples show higher strength. It seems that the
performed T6 treatment does not lead to precipitation of
all the alloying elements and yet there are elements in
the solution. To investigate this further, DSC analysis
was conducted on the samples.
3.2 DSC results
Figure 3 shows the DSC analysis for the T6 and
RRA treatment. Unlike the RRA treatment, the T6 curve
shows two exothermic peaks at temperatures between
200 °C and 245 °C. The RRA curve in this region has
only one small peak at about 240 °C. These peaks are
results of  and  precipitation from solid solution.14–16
In the RRA curve, there exists only one peak related to 
precipitation. This means that there are still a lot of ele-
ments in solid solution after T6 treatment suggesting that
either time or temperature or both are not enough for
diffusion and precipitation of the elements. As a result,
the maximum strength has not achieved.
3.3 Stress corrosion cracking
Stress corrosion tests were performed according to
the ASTM G139 standard for direct tension stress corro-
sion test method. Figure 4 shows the strength of T6 and
RRA samples after exposure to corrosive 3.5 % sodium
chloride solution without stressing and with a 207 MPa
stress after 4 d. As expected, strength is dropped in both
the T6 and RRA treatment after the SCC test. After the
SCC test with 207 MPa constant stresses for 4 d in corro-
sive solution, the strength of RRA treated samples
decreased from 612 MPa to 545 MPa. Accordingly, the
strength of the T6 treated samples dropped from 547
MPa to 367 MPa. Calculating the results in percentages,
gives us a better understanding of the strength reduction.
The test was conducted in corrosive solution in two
stress conditions: un-stressed and with constant 207 MPa
tension stress. The mechanical strength decreased for
both conditions. In order to better demonstrate the SCC
results, it is most useful to calculate the remaining
strength after SCC test and subtract the effect of corro-
sion methods other than SCC (such as pitting). To
achieve this, the percentage of SCR ratio of remaining
strength of stressed and unstressed samples is defined as
stress corrosion resistance SCR. The SCR results of the
alloy after the T6 and RRA heat treatments are shown in
Figure 5. The remaining strengths of the unstressed
samples after RRA and T6 treatments were 94.7 % and
93.3 %, respectively, which are not much different. It
means that unstressed corrosion results for both treat-
ments were almost identical. However, the remaining
strength percentages in the constant stress condition
were very different. The remaining strength percentage
was 89 % for the RRA treatment and 68 % for the T6
treatment. The SCR for the RRA treatment was 95 %
and for the T6 treatment it was 71 %. A significant im-
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836 833
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: The SCR results of the alloy after T6 and RRA heat
treatments
Figure 3: DSC curves for T6 and RRA treatment Figure 4: Direct tension stress corrosion cracking test after T6 and
RRA treatment
grain boundaries and a precipitate-free zone (PFZ) along
the grain boundary.6–10 However, the microstructural
characteristics of Al-Zn-Mg-Cu high strength aluminum
alloys are well known to have a strong influence not only
on the mechanical properties but also on the SCC sus-
ceptibility. Larger grain-boundary precipitates can trap
more atomic hydrogen to nucleate hydrogen bubbles,
thereby decreasing the hydrogen concentration at grain
boundaries below a critical value is considered to prevent
intergranular SCC fracture.11 Furthermore, the cathodic
grain-boundary precipitates grow by depleting solute
atoms. Studies showed that this leads to the broadening
of the anodic PFZ, which contained no strengthen pre-
cipitation phase and as a result was soft and weak.12 The
combination of tensile stress and anodic dissolution
caused SCC.
SCC resistance is of practical importance for the
industrial applications of the Al–Zn–Mg–Cu series
aluminum alloys. Various heat treatments in these alloys
offer very different SCC properties. The SCR criteria are
proposed to compare the SCC resistance.
The aim of this paper is to investigate the mechanical
strength and SCC of a 7000 series aluminum alloy after
conventional T6 and RRA treatments. The SCR criteria
have been defined to easily compare the SCC resistance
of this alloy after various treatments. SEM investigations
were used to determine the effect of the heat treatments
on the alloy’s microstructure with the aim of studying its
effect on SCC susceptibility. For comparison, the stress
corrosion resistance of the alloy with different heat
treatments was studied by the breaking load method
according to ASTM G139 standard.13
2 EXPERIMENTAL PART
The samples used in this study were received as
8-mm-thick sheets that were homogenized, hot rolled
and heat treated after being alloyed and casted. The
chemical composition of the alloy is shown in Table 1.
Table 1: Chemical composition (in mass fractions, w/%) of fabricated
alloy
Alloy
No. Zn Cu Mg Fe Si Zr Ag
1 8.5 2 2.1 0.18 0.16 0.20 0.19
An induction melting furnace used for melting and
the melt was poured in a water-cooled copper mold. The
as-cast specimens were homogenized at 460 °C for 24 h
and hot rolled to about 33 % reduction. Hot rolling was
performed with 40 min–1 rolling speed at 430 °C. The
specimens were milled to tensile samples according to
the ASTM E8M standard. A schematic showing the pre-
paration of tensile samples from the specimens is given
in Figure 1.
A solution heat treatment was done at 471 °C for 6 h
followed by a water quenching. T6 ageing was per-
formed for 24 h at 120 °C (T6 temper). The scheme of
the retrogression and re-ageing treatment is shown in
Figure 2.
Mechanical properties measurements were made at
ambient temperature on the specimens machined
according to the ASTM E8M-04 small size standard.
The average of three tests was used for each result. The
test’s strain rate was 10–3 /s.
The SEM analysis was performed on a TESCAN
scanning electron microscope. Thermal analysis was per-
formed in a DSC 1 Mettler Toledo differential scanning
calorimeter. Polished alloy disks with a diameter of
5 mm and 0.6 mm thick were sealed in aluminum pans
and heated in a flowing argon atmosphere at a constant
heating rate of 10 °C/min. The flowing rate of the argon
was 100 mL/min.
The stress corrosion tests were performed according
to the ASTM G139 standard. The Neutral 3.5 % Sodium
Chloride Solution was prepared in accordance with the
requirements of the ASTM G44 standard. The ASTM
G49 standard was used for preparation of direct tension
stress corrosion test specimens. A 207 MPa stress for the
cycle of 4 d was applied as per section 8-2 of the ASTM
G139 standard. In order to eliminate the corrosion other
than SCC (such as pitting), some specimens were tested
unstressed. The ratio of remaining strength of stressed
and unstressed samples represented in percent, were
calculated as SCR criteria for different treatment.
3 RESULTS AND DISCUSSION
3.1 Mechanical strength
The mechanical strength of the alloy was 547 MPa
after the T6 treatment and was 612 MPa after the RRA
treatment. The RRA treatment leads to an about 12 %
increase in mechanical strength. While T6 treatment is
carried out to ensure maximum strength; the RRA
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
832 Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Scheme of the RRA treatment
Figure 1: Schematic showing the preparation of tensile samples from
the specimens
treated samples show higher strength. It seems that the
performed T6 treatment does not lead to precipitation of
all the alloying elements and yet there are elements in
the solution. To investigate this further, DSC analysis
was conducted on the samples.
3.2 DSC results
Figure 3 shows the DSC analysis for the T6 and
RRA treatment. Unlike the RRA treatment, the T6 curve
shows two exothermic peaks at temperatures between
200 °C and 245 °C. The RRA curve in this region has
only one small peak at about 240 °C. These peaks are
results of  and  precipitation from solid solution.14–16
In the RRA curve, there exists only one peak related to 
precipitation. This means that there are still a lot of ele-
ments in solid solution after T6 treatment suggesting that
either time or temperature or both are not enough for
diffusion and precipitation of the elements. As a result,
the maximum strength has not achieved.
3.3 Stress corrosion cracking
Stress corrosion tests were performed according to
the ASTM G139 standard for direct tension stress corro-
sion test method. Figure 4 shows the strength of T6 and
RRA samples after exposure to corrosive 3.5 % sodium
chloride solution without stressing and with a 207 MPa
stress after 4 d. As expected, strength is dropped in both
the T6 and RRA treatment after the SCC test. After the
SCC test with 207 MPa constant stresses for 4 d in corro-
sive solution, the strength of RRA treated samples
decreased from 612 MPa to 545 MPa. Accordingly, the
strength of the T6 treated samples dropped from 547
MPa to 367 MPa. Calculating the results in percentages,
gives us a better understanding of the strength reduction.
The test was conducted in corrosive solution in two
stress conditions: un-stressed and with constant 207 MPa
tension stress. The mechanical strength decreased for
both conditions. In order to better demonstrate the SCC
results, it is most useful to calculate the remaining
strength after SCC test and subtract the effect of corro-
sion methods other than SCC (such as pitting). To
achieve this, the percentage of SCR ratio of remaining
strength of stressed and unstressed samples is defined as
stress corrosion resistance SCR. The SCR results of the
alloy after the T6 and RRA heat treatments are shown in
Figure 5. The remaining strengths of the unstressed
samples after RRA and T6 treatments were 94.7 % and
93.3 %, respectively, which are not much different. It
means that unstressed corrosion results for both treat-
ments were almost identical. However, the remaining
strength percentages in the constant stress condition
were very different. The remaining strength percentage
was 89 % for the RRA treatment and 68 % for the T6
treatment. The SCR for the RRA treatment was 95 %
and for the T6 treatment it was 71 %. A significant im-
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836 833
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: The SCR results of the alloy after T6 and RRA heat
treatments
Figure 3: DSC curves for T6 and RRA treatment Figure 4: Direct tension stress corrosion cracking test after T6 and
RRA treatment
provement in the stress corrosion cracking has been
achieved after the RRA treatment.
3.4 SEM and EDS analysis
The SEM micrographs of the T6 and RRA samples
are shown in Figures 6 and 7.
Comparison of these two micrographs shows that,
after the RRA treatment the grain-boundary precipitates
have lost their continuity and divided to relatively larger
particles. This is an essential factor for reducing the
intergranular corrosion and consequently reducing the
stress corrosion cracking. Two basic mechanisms of the
SCC in 7000 series aluminum alloys have been pro-
posed: anodic dissolution and cathodic dissolution
(hydrogen embrittlement).6 The precipitates in the grain
boundaries have the electrode potential different from
the Al matrix. This would result in the anodic dissolution
and form defects in a chloride solution. The hydrogen
atoms produced in the crack tip also lead to the hydrogen
embrittlement in the grain boundaries. The large size and
distance of the grain-boundary particles could decrease
the anodic dissolution speed. These particles can also act
as the trapping sites for atomic hydrogen and transform
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
834 Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 9: EDS analysis of RRA sample: a) matrix; b) grain boundary
precipitates
Figure 7: SEM micrograph of RRA sample showing less continuous
precipitates along the grain boundary in comparison to the T6 sample
Figure 6: SEM micrograph of T6 sample showing continuous
precipitates along the grain boundary
Figure 8: EDS analysis of T6 sample: a) matrix, b) grain boundary
precipitates
them to molecular hydrogen bubbles to reduce the con-
centration of the atomic hydrogen at the grain boun-
daries. Both the size and distance of the grain-boundary
precipitates in RRA treated alloy were larger than that in
T6 treated alloy, therefore, the SCR criteria after RRA
treatment is higher than the T6 treatment.
Another factor showing SCC resistance improvement
after RRA treatment is the concentration changes of pre-
cipitates and matrix. The Zn and Cu atoms have opposite
effects on the electrochemical properties of alumi-
num.17,18 In a corrosive medium, the  phase in grain
boundary is very active and anodic with respect to the
matrix.19 Cu enrichment of precipitates after RRA
increases the electrode potential of the grain boundaries
and decreases the activity of grain-boundary phases.
Figures 9 and 10 show the EDS analysis of the grain-
boundary precipitates and matrix after RRA and T6
treatment. Figure 8a shows the Cu-rich matrix after the
T6 treatment while according to Figure 8b, the preci-
pitates are Zn-reach. After RRA treatment, the matrix
depleted from Cu atoms and the grain boundary preci-
pitates were Cu-enrichment (Figures 9a to 9b). Large
variations in concentration of solute atoms happen
during the RRA treatment. Zn atoms have a higher diffu-
sivity than Cu atoms and leave the precipitates first after
the reversion stage.20 As a result, the Zn content of the
matrix increases and a corresponding increase in the Cu
precipitate content happens. This happens reversely in
the T6 treatment, where Zn incorporates first the preci-
pitates, resulting in a high Zn content of the precipitates
and high Cu matrix content. The Zn and Cu atoms have
opposite effects on the electrochemical properties of
aluminum.21 A change in the corrosion behavior of the
7000 series aluminum alloy is related to changes in pre-
cipitate composition and particularly to a Cu enrichment
of the precipitates.22 During the RRA treatment, the
precipitates enrich by Cu atoms (Figure 9b) and this
factor as well as other factors improves the SCC
resistance.
4 CONCLUSIONS
The mechanical strength of an Al–8.5Zn–2.1Mg–
2Cu–0.2Ag (w/%) alloy after conventional T6 treatment
was lower than the mechanical strength after RRA
treatment.
SCR is a useful criteria to compare the SCC in
Al-Zn-Mg-Cu alloys after various heat treatments. The
SCR of the Al–8.5Zn–2.1Mg–2Cu–0.2Ag (w/%) alloy
after conventional T6 treatment was 71 % and was 95 %
after RRA treatment. The SCC improved significantly
after the RRA treatment.
The SCC improvement after the RRA treatment
related to Cu enrichment of the precipitates along with
the discontinuities of the grain-boundary precipitates.
5 REFERENCES
1 T. Dursun, C. Soutis, Recent developments in advanced aircraft
aluminium alloys, Materials & Design, 56 (2014), 862–871,
doi:10.1016/j.matdes.2013.12.002
2 D. Liu, B. Xiong, F. Bian, Z. Li, X. Li, Y. Zhang, F. Wang, H. Liu, In
situ studies of microstructure evolution and properties of an
Al–7.5Zn–1.7Mg–1.4Cu–0.12Zr alloy during retrogression and
reaging, Materials & Design, 56 (2014), 1020–1024, doi:10.1016/
j.matdes.2013.12.006
3 J. Li, N. Birbilis, C. Li, Z. Jia, B. Cai, Z. Zheng, Influence of
retrogression temperature and time on the mechanical properties and
exfoliation corrosion behavior of aluminium alloy AA7150,
Materials Characterization, 60 (2009) 11, 1334–1341, doi:10.1016/
j.matchar.2009.06.007
4 R. Bucci, C. Warren, E. Starke, The Need for New Materials in
Aging Aircraft Structures, Journal of Aircraft, 37 (2000) 1,
122–129, doi:10.2514/2.2571
5 D. Feng, X. M. Zhang, S. D. Liu, and Y. L. Deng, Non-isothermal
retrogression and re-ageing treatment schedule for AA7055 thick
plate, Materials & Design, 60 (2014), 208–217, doi:10.1016/
j.matdes.2014.03.064
6 D. Najjar, T. Magnin, T. J. Warner, Influence of critical surface
defects and localized competition between anodic dissolution and
hydrogen effects during stress corrosion cracking of a 7050 alumi-
nium alloy, Materials Science and Engineering: A, 238 (1997) 2,
293–302, doi:10.1016/S0921-5093(97)00369-9
7 T. Burleigh, The postulated mechanisms for stress corrosion cracking
of aluminum alloys: a review of the literature 1980-1989, Corrosion,
47 (1991) 2, 89–98, doi:10.5006/1.3585235
8 R. Jones, D. Baer, M. Danielson, J. Vetrano, Role of Mg in the stress
corrosion cracking of an Al-Mg alloy, Metallurgical and Materials
Transactions A, 32 (2001) 7, 1699–1711, doi:10.1007/s11661-
001-0148-0
9 E. Pugh, Progress toward understanding the stress corrosion
problem, Corrosion, 41 (1985) 9, 517–526, doi:10.5006/1.3583022
10 A. Sedriks, J. Green, D. Novak, Corrosion processes and solution
chemistry within stress corrosion cracks in aluminum alloys,
Proceedings of NACE, 1971, 569–575
11 T. Tsai, T. Chuang, Role of grain size on the stress corrosion crack-
ing of 7475 aluminum alloys, Materials Science and Engineering: A,
225 (1997) 1–2, 135–144, doi:10.1016/S0921-5093(96)10840-6
12 T. Pardoen, D. Dumont, A. Deschamps, Y. Brechet, Grain boundary
versus transgranular ductile failure, Journal of the Mechanics and
Physics of Solids, 51 (2003) 4, 637–665, doi:10.1016/S0022-
5096(02)00102-3
13 ASTM 139:2005(G)- Standard test method for determining stress-
corrosion cracking resistance of heat-treatable aluminum alloy
products using breaking load method, ASTM International, West
Conshohocken
14 K. Ghosh, N. Gao, Determination of kinetic parameters from
calorimetric study of solid state reactions in 7150 Al-Zn-Mg alloy,
Transactions of Nonferrous Metals Society of China, 21 (2011) 6,
1199–1209, doi:10.1016/S1003-6326(11)60843-1
15 T. Marlaud, A. Deschamps, F. Bley, W. Lefebvre, B. Baroux,
Influence of alloy composition and heat treatment on precipitate
composition in Al–Zn–Mg–Cu alloys, Acta Materialia, 58 (2010) 1,
248–260, doi:10.1016/j.actamat.2009.09.003
16 J. G. Tang, H. Chen, X. M. Zhang, S. D. Liu, W. J. Liu, H. Ouyang,
H. P. Li, Influence of quench-induced precipitation on aging behavior
of Al-Zn-Mg-Cu alloy, Transactions of Nonferrous Metals Society of
China, 22 (2012) 6, 1255–1263, doi:10.1016/S1003-6326(11)
61313-7
17 R. W. Revie, H. H. Uhlig, Uhlig’s corrosion handbook: vol. 51, John
Wiley & Sons, Hoboken 2011
18 G. Silva, B. Rivolta, R. Gerosa, U. Derudi, Study of the SCC
Behavior of 7075 Aluminum Alloy After One-Step Aging at 163° C,
Journal of Materials Engineering and Performance, 22 (2013) 1,
210–214, doi:10.1007/s11665-012-0221-4
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836 835
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
provement in the stress corrosion cracking has been
achieved after the RRA treatment.
3.4 SEM and EDS analysis
The SEM micrographs of the T6 and RRA samples
are shown in Figures 6 and 7.
Comparison of these two micrographs shows that,
after the RRA treatment the grain-boundary precipitates
have lost their continuity and divided to relatively larger
particles. This is an essential factor for reducing the
intergranular corrosion and consequently reducing the
stress corrosion cracking. Two basic mechanisms of the
SCC in 7000 series aluminum alloys have been pro-
posed: anodic dissolution and cathodic dissolution
(hydrogen embrittlement).6 The precipitates in the grain
boundaries have the electrode potential different from
the Al matrix. This would result in the anodic dissolution
and form defects in a chloride solution. The hydrogen
atoms produced in the crack tip also lead to the hydrogen
embrittlement in the grain boundaries. The large size and
distance of the grain-boundary particles could decrease
the anodic dissolution speed. These particles can also act
as the trapping sites for atomic hydrogen and transform
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
834 Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 9: EDS analysis of RRA sample: a) matrix; b) grain boundary
precipitates
Figure 7: SEM micrograph of RRA sample showing less continuous
precipitates along the grain boundary in comparison to the T6 sample
Figure 6: SEM micrograph of T6 sample showing continuous
precipitates along the grain boundary
Figure 8: EDS analysis of T6 sample: a) matrix, b) grain boundary
precipitates
them to molecular hydrogen bubbles to reduce the con-
centration of the atomic hydrogen at the grain boun-
daries. Both the size and distance of the grain-boundary
precipitates in RRA treated alloy were larger than that in
T6 treated alloy, therefore, the SCR criteria after RRA
treatment is higher than the T6 treatment.
Another factor showing SCC resistance improvement
after RRA treatment is the concentration changes of pre-
cipitates and matrix. The Zn and Cu atoms have opposite
effects on the electrochemical properties of alumi-
num.17,18 In a corrosive medium, the  phase in grain
boundary is very active and anodic with respect to the
matrix.19 Cu enrichment of precipitates after RRA
increases the electrode potential of the grain boundaries
and decreases the activity of grain-boundary phases.
Figures 9 and 10 show the EDS analysis of the grain-
boundary precipitates and matrix after RRA and T6
treatment. Figure 8a shows the Cu-rich matrix after the
T6 treatment while according to Figure 8b, the preci-
pitates are Zn-reach. After RRA treatment, the matrix
depleted from Cu atoms and the grain boundary preci-
pitates were Cu-enrichment (Figures 9a to 9b). Large
variations in concentration of solute atoms happen
during the RRA treatment. Zn atoms have a higher diffu-
sivity than Cu atoms and leave the precipitates first after
the reversion stage.20 As a result, the Zn content of the
matrix increases and a corresponding increase in the Cu
precipitate content happens. This happens reversely in
the T6 treatment, where Zn incorporates first the preci-
pitates, resulting in a high Zn content of the precipitates
and high Cu matrix content. The Zn and Cu atoms have
opposite effects on the electrochemical properties of
aluminum.21 A change in the corrosion behavior of the
7000 series aluminum alloy is related to changes in pre-
cipitate composition and particularly to a Cu enrichment
of the precipitates.22 During the RRA treatment, the
precipitates enrich by Cu atoms (Figure 9b) and this
factor as well as other factors improves the SCC
resistance.
4 CONCLUSIONS
The mechanical strength of an Al–8.5Zn–2.1Mg–
2Cu–0.2Ag (w/%) alloy after conventional T6 treatment
was lower than the mechanical strength after RRA
treatment.
SCR is a useful criteria to compare the SCC in
Al-Zn-Mg-Cu alloys after various heat treatments. The
SCR of the Al–8.5Zn–2.1Mg–2Cu–0.2Ag (w/%) alloy
after conventional T6 treatment was 71 % and was 95 %
after RRA treatment. The SCC improved significantly
after the RRA treatment.
The SCC improvement after the RRA treatment
related to Cu enrichment of the precipitates along with
the discontinuities of the grain-boundary precipitates.
5 REFERENCES
1 T. Dursun, C. Soutis, Recent developments in advanced aircraft
aluminium alloys, Materials & Design, 56 (2014), 862–871,
doi:10.1016/j.matdes.2013.12.002
2 D. Liu, B. Xiong, F. Bian, Z. Li, X. Li, Y. Zhang, F. Wang, H. Liu, In
situ studies of microstructure evolution and properties of an
Al–7.5Zn–1.7Mg–1.4Cu–0.12Zr alloy during retrogression and
reaging, Materials & Design, 56 (2014), 1020–1024, doi:10.1016/
j.matdes.2013.12.006
3 J. Li, N. Birbilis, C. Li, Z. Jia, B. Cai, Z. Zheng, Influence of
retrogression temperature and time on the mechanical properties and
exfoliation corrosion behavior of aluminium alloy AA7150,
Materials Characterization, 60 (2009) 11, 1334–1341, doi:10.1016/
j.matchar.2009.06.007
4 R. Bucci, C. Warren, E. Starke, The Need for New Materials in
Aging Aircraft Structures, Journal of Aircraft, 37 (2000) 1,
122–129, doi:10.2514/2.2571
5 D. Feng, X. M. Zhang, S. D. Liu, and Y. L. Deng, Non-isothermal
retrogression and re-ageing treatment schedule for AA7055 thick
plate, Materials & Design, 60 (2014), 208–217, doi:10.1016/
j.matdes.2014.03.064
6 D. Najjar, T. Magnin, T. J. Warner, Influence of critical surface
defects and localized competition between anodic dissolution and
hydrogen effects during stress corrosion cracking of a 7050 alumi-
nium alloy, Materials Science and Engineering: A, 238 (1997) 2,
293–302, doi:10.1016/S0921-5093(97)00369-9
7 T. Burleigh, The postulated mechanisms for stress corrosion cracking
of aluminum alloys: a review of the literature 1980-1989, Corrosion,
47 (1991) 2, 89–98, doi:10.5006/1.3585235
8 R. Jones, D. Baer, M. Danielson, J. Vetrano, Role of Mg in the stress
corrosion cracking of an Al-Mg alloy, Metallurgical and Materials
Transactions A, 32 (2001) 7, 1699–1711, doi:10.1007/s11661-
001-0148-0
9 E. Pugh, Progress toward understanding the stress corrosion
problem, Corrosion, 41 (1985) 9, 517–526, doi:10.5006/1.3583022
10 A. Sedriks, J. Green, D. Novak, Corrosion processes and solution
chemistry within stress corrosion cracks in aluminum alloys,
Proceedings of NACE, 1971, 569–575
11 T. Tsai, T. Chuang, Role of grain size on the stress corrosion crack-
ing of 7475 aluminum alloys, Materials Science and Engineering: A,
225 (1997) 1–2, 135–144, doi:10.1016/S0921-5093(96)10840-6
12 T. Pardoen, D. Dumont, A. Deschamps, Y. Brechet, Grain boundary
versus transgranular ductile failure, Journal of the Mechanics and
Physics of Solids, 51 (2003) 4, 637–665, doi:10.1016/S0022-
5096(02)00102-3
13 ASTM 139:2005(G)- Standard test method for determining stress-
corrosion cracking resistance of heat-treatable aluminum alloy
products using breaking load method, ASTM International, West
Conshohocken
14 K. Ghosh, N. Gao, Determination of kinetic parameters from
calorimetric study of solid state reactions in 7150 Al-Zn-Mg alloy,
Transactions of Nonferrous Metals Society of China, 21 (2011) 6,
1199–1209, doi:10.1016/S1003-6326(11)60843-1
15 T. Marlaud, A. Deschamps, F. Bley, W. Lefebvre, B. Baroux,
Influence of alloy composition and heat treatment on precipitate
composition in Al–Zn–Mg–Cu alloys, Acta Materialia, 58 (2010) 1,
248–260, doi:10.1016/j.actamat.2009.09.003
16 J. G. Tang, H. Chen, X. M. Zhang, S. D. Liu, W. J. Liu, H. Ouyang,
H. P. Li, Influence of quench-induced precipitation on aging behavior
of Al-Zn-Mg-Cu alloy, Transactions of Nonferrous Metals Society of
China, 22 (2012) 6, 1255–1263, doi:10.1016/S1003-6326(11)
61313-7
17 R. W. Revie, H. H. Uhlig, Uhlig’s corrosion handbook: vol. 51, John
Wiley & Sons, Hoboken 2011
18 G. Silva, B. Rivolta, R. Gerosa, U. Derudi, Study of the SCC
Behavior of 7075 Aluminum Alloy After One-Step Aging at 163° C,
Journal of Materials Engineering and Performance, 22 (2013) 1,
210–214, doi:10.1007/s11665-012-0221-4
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836 835
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19 S. Maitra, G. English, Mechanism of localized corrosion of 7075
alloy plate, Metallurgical and Materials Transactions A, 12 (1981) 3,
535–541, doi:10.1007/BF02648553
20 H. Bakker, H. Bonzel, C. Bruff, M. Dayananda, W. Gust, J. Horvath,
I. Kaur, G. Kidson, A. LeClaire, H. Mehrer, Diffusion in solid metals
and alloys/Diffusion in festen metallen und legierungen: vol. 26,
Springer, Berlin 1990
21 E. Hollingsworth, H. Hunsicker, Corrosion of aluminum and alumi-
num alloys, ASM Handbook, 13 (1987), 583–609
22 J. T. Staley, S. Byrne, E. Colvin, K. Kinnear, Corrosion and stress-
corrosion of 7xxx-w products, Materials Science Forum, 217 (1996),
1587–1592, doi:10.4028/MSF.217-222.1587
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836 Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
837–842
CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED
ELECTROLESS DUBLEX Ni-P/Ni-B-W COATINGS
KOROZIJSKA ODPORNOST PLATIRANIH IN NEELEKTRI^NO
TOPOLOTNO OBDELANIH DUPLEKS Ni-P/Ni-B-W PREVLEK
Behiye Yüksel1, Garip Erdogan2, Fatih Erdem Bastan2, Rasid Ahmed Yýldýz3
1Istanbul Aydin University, Mechanical Engineering Department, 34668 Istanbul, Turkey
2Sakarya University, Engineering Faculty, Department of Metallurgy and Materials Engineering, 54187 Sakarya, Turkey
3TU Istanbul, Mechanical Engineering Faculty, 34437 Istanbul, Turkey
behiyeyuksel@aydin.edu.tr
Prejem rokopisa – received: 2016-10-20; sprejem za objavo – accepted for publication: 2017-03-10
doi:10.17222/mit.2016.304
In this study, Ni-P/Ni-B-W dublex coatings were deposited on carbon steel substrates (AISI 1020) using the electroless plating
process and their microstructure and corrosion properties were systematically evaluated based on different heat-treatment
temperatures. Both, the surface morphology and cross-sectional morphology of the Ni-P/Ni-B-W coatings were studied using a
scanning electron microscope (SEM), while X-ray diffraction (XRD) was applied for examining the structural modifications.
The amorphous coating began to crystallize at a heat-treatment temperature of 350 °C. Potentiodynamic polarization
measurements were carried out in an aqueous medium containing 3.5 % NaCl for evaluating the corrosion resistance of
as-plated and heat-treated dublex coatings. The corrosion potentials of dublex coatings were observed to shift toward more
positive values with increased heat-treatment temperatures. Depending on the heat-treatment temperature, it was identified that
the crystallized dublex coatings generally had better corrosion resistance than the amorphous coating.
Keywords: Ni-P/Ni-B-W coating, corrosion, heat treatment
V {tudiji so bile dvojne Ni-P/Ni-B-W prevleke nane{ene na substrate iz ogljikovega jekla (AISI 1020) s postopkom neelek-
tri~nega platiranja. Sistemati~no so bile ocenjene mikrostruktura ter korozijske lastnosti izdelanih prevlek, glede na razli~ne
temperature toplotne obdelave. Morfologije povr{in in profili izdelanih Ni-P/Ni-B-W prevlek so bili pregledani z vrsti~nim
elektronskim mikroskopom (SEM), medtem ko so bile z rentgensko difrakcijsko analizo (XRD) ugotovljene strukturne
spremembe. Amorfna prevleka je za~ela kristalizirati pri temperaturi toplotne obdelave 350 °C. Korozijska odpornost platiranih
in toplotno obdelanih dupleks prevlek je bila ocenjena s pomo~jo meritev potenciodinami~ne polarizacije v 3,5 % vodni razto-
pini NaCl. Opazovali so korozijske potenciale dupleks prevlek, da bi dobili bolj pozitivne vrednosti z zvi{animi temperaturami
topolotne obdelave. Ugotovljeno je bilo, da imajo kristalizirane dupleks prevleke na splo{no bolj{o korozijsko odpornost kot
amorfne prevleke.
Klju~ne besede: Ni-P/Ni-B-W dvojna prevleka, korozija, toplotna obdelava
1 INTRODUCTION
Studied for the first time in 1946 by Brenner and
Riddell, the electroless plating process has been used in
many industrial applications because of its superior
characteristics over electroplating, such as the ability to
plate insulation materials and having a homogeneous
coating-thickness distribution.1 Among electroless coat-
ing types, electroless nickel coating is the most popular
for having good hardness, wear and corrosion resistance
properties.2 If the electroless Ni coating family is re-
viewed, besides pure Ni coatings, Ni-P and Ni-B coat-
ings deposited using reducing agents such as hypo-
phosphite, borohydride or dimethylamine borane stand
out. The major advantages of Ni-P coatings are their low
cost, high corrosion resistance and easy process control.3
Nonetheless, a borohydride reduced nickel coating is
harder and has higher wear resistance than tool steel and
hard chrome coatings.4,5 On the other hand, some studies
– although they are limited in number – have highlighted
the addition of tungsten to this coating system for in-
creasing its corrosion resistance, which is lower than that
of a Ni-P coating.6–9 As a result of the co-deposition of
this refractory material (which cannot be reduced in an
aqueous solution as metallic tungsten) together with iron
group metals, developing coatings with attractive corro-
sion and tribological properties is possible.10–12 Recently,
research for obtaining multilayer coatings to increase
existing coating corrosion resistance to higher levels has
also been conducted.13,14
The aim of the present study was to evaluate the
effect of different heat-treatment temperatures on the
microstructure and corrosion properties of electroless
Ni-P/Ni-B-W coatings. To the best of our knowledge, no
others studies are available on this topic.
2 EXPERIMENTAL PART
For this experiment, 10 mm × 10 mm × 60 mm plain
carbon steel (AISI 1020) plates were used as a substrate
material for developing Ni-P/Ni-B-W dublex coatings.
Prior to the plating process, the surfaces of all the
Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842 837
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.193:621.78:669.058 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)837(2017)
19 S. Maitra, G. English, Mechanism of localized corrosion of 7075
alloy plate, Metallurgical and Materials Transactions A, 12 (1981) 3,
535–541, doi:10.1007/BF02648553
20 H. Bakker, H. Bonzel, C. Bruff, M. Dayananda, W. Gust, J. Horvath,
I. Kaur, G. Kidson, A. LeClaire, H. Mehrer, Diffusion in solid metals
and alloys/Diffusion in festen metallen und legierungen: vol. 26,
Springer, Berlin 1990
21 E. Hollingsworth, H. Hunsicker, Corrosion of aluminum and alumi-
num alloys, ASM Handbook, 13 (1987), 583–609
22 J. T. Staley, S. Byrne, E. Colvin, K. Kinnear, Corrosion and stress-
corrosion of 7xxx-w products, Materials Science Forum, 217 (1996),
1587–1592, doi:10.4028/MSF.217-222.1587
M. SHAKOURI et al.: DEVELOPMENT OF A HEAT TREATMENT FOR INCREASING THE MECHANICAL PROPERTIES ...
836 Materiali in tehnologije / Materials and technology 51 (2017) 5, 831–836
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
837–842
CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED
ELECTROLESS DUBLEX Ni-P/Ni-B-W COATINGS
KOROZIJSKA ODPORNOST PLATIRANIH IN NEELEKTRI^NO
TOPOLOTNO OBDELANIH DUPLEKS Ni-P/Ni-B-W PREVLEK
Behiye Yüksel1, Garip Erdogan2, Fatih Erdem Bastan2, Rasid Ahmed Yýldýz3
1Istanbul Aydin University, Mechanical Engineering Department, 34668 Istanbul, Turkey
2Sakarya University, Engineering Faculty, Department of Metallurgy and Materials Engineering, 54187 Sakarya, Turkey
3TU Istanbul, Mechanical Engineering Faculty, 34437 Istanbul, Turkey
behiyeyuksel@aydin.edu.tr
Prejem rokopisa – received: 2016-10-20; sprejem za objavo – accepted for publication: 2017-03-10
doi:10.17222/mit.2016.304
In this study, Ni-P/Ni-B-W dublex coatings were deposited on carbon steel substrates (AISI 1020) using the electroless plating
process and their microstructure and corrosion properties were systematically evaluated based on different heat-treatment
temperatures. Both, the surface morphology and cross-sectional morphology of the Ni-P/Ni-B-W coatings were studied using a
scanning electron microscope (SEM), while X-ray diffraction (XRD) was applied for examining the structural modifications.
The amorphous coating began to crystallize at a heat-treatment temperature of 350 °C. Potentiodynamic polarization
measurements were carried out in an aqueous medium containing 3.5 % NaCl for evaluating the corrosion resistance of
as-plated and heat-treated dublex coatings. The corrosion potentials of dublex coatings were observed to shift toward more
positive values with increased heat-treatment temperatures. Depending on the heat-treatment temperature, it was identified that
the crystallized dublex coatings generally had better corrosion resistance than the amorphous coating.
Keywords: Ni-P/Ni-B-W coating, corrosion, heat treatment
V {tudiji so bile dvojne Ni-P/Ni-B-W prevleke nane{ene na substrate iz ogljikovega jekla (AISI 1020) s postopkom neelek-
tri~nega platiranja. Sistemati~no so bile ocenjene mikrostruktura ter korozijske lastnosti izdelanih prevlek, glede na razli~ne
temperature toplotne obdelave. Morfologije povr{in in profili izdelanih Ni-P/Ni-B-W prevlek so bili pregledani z vrsti~nim
elektronskim mikroskopom (SEM), medtem ko so bile z rentgensko difrakcijsko analizo (XRD) ugotovljene strukturne
spremembe. Amorfna prevleka je za~ela kristalizirati pri temperaturi toplotne obdelave 350 °C. Korozijska odpornost platiranih
in toplotno obdelanih dupleks prevlek je bila ocenjena s pomo~jo meritev potenciodinami~ne polarizacije v 3,5 % vodni razto-
pini NaCl. Opazovali so korozijske potenciale dupleks prevlek, da bi dobili bolj pozitivne vrednosti z zvi{animi temperaturami
topolotne obdelave. Ugotovljeno je bilo, da imajo kristalizirane dupleks prevleke na splo{no bolj{o korozijsko odpornost kot
amorfne prevleke.
Klju~ne besede: Ni-P/Ni-B-W dvojna prevleka, korozija, toplotna obdelava
1 INTRODUCTION
Studied for the first time in 1946 by Brenner and
Riddell, the electroless plating process has been used in
many industrial applications because of its superior
characteristics over electroplating, such as the ability to
plate insulation materials and having a homogeneous
coating-thickness distribution.1 Among electroless coat-
ing types, electroless nickel coating is the most popular
for having good hardness, wear and corrosion resistance
properties.2 If the electroless Ni coating family is re-
viewed, besides pure Ni coatings, Ni-P and Ni-B coat-
ings deposited using reducing agents such as hypo-
phosphite, borohydride or dimethylamine borane stand
out. The major advantages of Ni-P coatings are their low
cost, high corrosion resistance and easy process control.3
Nonetheless, a borohydride reduced nickel coating is
harder and has higher wear resistance than tool steel and
hard chrome coatings.4,5 On the other hand, some studies
– although they are limited in number – have highlighted
the addition of tungsten to this coating system for in-
creasing its corrosion resistance, which is lower than that
of a Ni-P coating.6–9 As a result of the co-deposition of
this refractory material (which cannot be reduced in an
aqueous solution as metallic tungsten) together with iron
group metals, developing coatings with attractive corro-
sion and tribological properties is possible.10–12 Recently,
research for obtaining multilayer coatings to increase
existing coating corrosion resistance to higher levels has
also been conducted.13,14
The aim of the present study was to evaluate the
effect of different heat-treatment temperatures on the
microstructure and corrosion properties of electroless
Ni-P/Ni-B-W coatings. To the best of our knowledge, no
others studies are available on this topic.
2 EXPERIMENTAL PART
For this experiment, 10 mm × 10 mm × 60 mm plain
carbon steel (AISI 1020) plates were used as a substrate
material for developing Ni-P/Ni-B-W dublex coatings.
Prior to the plating process, the surfaces of all the
Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842 837
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.193:621.78:669.058 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)837(2017)
samples were mechanically cleaned (up to 1200 grade)
and then soaked in trichloroethylene and cleaned with
detergent in an ultrasonic bath at 70 °C. Finally, sample
surfaces were activated in 30 % of volume fractions. HCl
for two minutes and then rinsed in distilled water. The
cleaned substrates were then soaked in Ni-P and Ni-B-W
electroless plating baths for two hours, respectively.
Reagent-grade chemicals and distilled water were
used for the preparation of all the electrolytes and the pH
was adjusted to 4.8±0.2 and 13.5±0.2 with H2SO4 or
NaOH in Ni-P and Ni-B-W plating baths, respectively.
The temperature of the baths was maintained at 90±2 °C
for Ni-P and 88±2 °C for Ni-B-W.
Deposition of Ni-P and Ni-B-W was carried out in an
aqueous bath containing 15 g/L NiSO4.6H2O, 26 g/L
Na2H2PO2·H2O, 13 g/L NaC2H3O2, 12 mL/L HF and
8 g/L NH4HF2 for the Ni-P coating, and 24 g/L NiCl2,
60 mL/L EDTA, 26.5 g/L KOH, 120 g/L NaBH4,
263 g/L NaOH, 2.6g/L PbWO4, 13 g/L EDTA and 40 g/L
Na2WO4·2H2O for Ni-B-W coatings, respectively.
The coated specimens were heat treated at 250 °C,
300 °C, 350 °C and 650 °C for 1 h each at a heating rate
of 5 °C/min in an air-circulated furnace, followed by
slow furnace cooling to room temperature.
The composition of the dublex coatings was deter-
mined by energy-dispersive X-ray analysis (EDX) using
a Link Analytical QX-2000 attached to the SEM appa-
ratus. The crystal structures of the dublex coatings were
examined using a Philips PW 3710 grazing incidence
x-ray diffractometer (Cu-K radiation). A Jeol
JSM-7000F FE-SEM was used to characterize the
microstructures and morphology of the coatings. The
corrosion behavior was investigated using a PGZ 301
Dynamic Voltammetry and VoltaMaster4 software.
Electrochemical experiments were carried out in a
3.5 % NaCl aqueous solution in a three-electrode cell at
room temperature. In this cell, a platinum electrode was
used as a counter electrode and a saturated calomel
electrode was used as a reference electrode; the dublex
coating was used as a working electrode and by masking
with silicon, only a 1 cm2 surface area was left on the
samples for exposure to the electrolyte. The dynamic
potential scanning technique was used for obtaining the
polarization curves of dublex coating samples. Prior to
the potentiodynamic measurement, samples were held
for approximately 15 min in the electrolyte and then
electrode potential was raised from -600mV to +200 mV
at a rate of 10 mV/min.
3 RESULTS
The XRD patterns of the as-plated and heat-treated
Ni-P/Ni-B-W coatings, and their chemical compositions
are in Figure 1 and in Table 1. SEM micrographs of
Ni-P/Ni-B-W coatings before and after the heat treat-
ments are in Figure 2.
Table 1: Chemical compositions of the Ni-P/Ni-B-W dublex coatings
Type of
coating Ni (w/%) P (w/%) B (w/%) W (w/%)
Ni-P 87.27 12.73 - -
Ni-B-W 87.45 - 7.86 4.69
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
838 Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Surface morphology of dublex coating a), b) as-plated
Ni-P/Ni-B-W, c) heat treated at 300 °C, d) heat treated at 650 °C
Figure 1: XRD spectra for a): as-plated Ni-P/Ni-B-W, b) 250 °C, c)
300 °C, d) 350 °C and e) 650 °C
The cross-sections of the Ni-P/Ni-B-W dublex coat-
ings are shown in Figure 3.
Potentiodynamic curves of as-plated and heat-treated
Ni-P/Ni-B-W coatings at different temperatures obtained
in a 3.5 % NaCl aqueous medium are shown in Figure 4.
The changes in the corrosion current density (Icorr)
values with heat treatment are given in Table 2.
Table 2: Corrosion resistance of as-plated and heat-treated
Ni-P/Ni-B-W dublex coatings at different temperatures in 3.5 % NaCl
solution
NiP/NiBW Icorr (μA/cm2) Ecorr (mV)
As-plated 1.55±0.59 -506±32
300 °C 1.70±0.41 -470±23
350 °C 0.71±0.20 -382 ±18
650 °C 0.57±0.09 -280±11
The surface morphologies of the as-plated sample
and heat-treated sample at 650 °C following potentio-
dynamic polarization measurements are shown in Fig-
ure 5.
The corrosion current density and corrosion potential
values obtained by Tafel interpolation are shown in
Table 3.
Table 3: Corrosion resistance of substrate and Ni-B-W and
Ni-P/Ni-B-W dublex coatings heat treated at 650 °C in 3.5 % NaCl
solution
Sample Icorr (μA/cm2) Ecorr (mV)
Carbon steel
substrate 4.59±1.01 -598±46
NiBW 3.40±0.35 -298±17
NiP/NiBW 0.57±0.09 -280±11
Figure 6 shows the polarization curves of the
Ni-B-W and dublex Ni-P/Ni-B-W coatings heat treated
at 650 °C and the steel substrate obtained in a 3.5 %
NaCl aqueous medium.
4 DISCUSSION
The Ni-P/Ni-B-W coating diffraction patterns up to
300 °C reveal a single broad peak indicating an amor-
phous coating structure under these conditions (Fig-
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842 839
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Cross-section of Ni-P/Ni-B-W dublex coating: a) as-plated,
b) heat treated at 250 °C c) heat treated at 350 °C, d) heat treated at
Figure 4: Polarization curves in a 3.5 % NaCl aqueous medium of
as-plated and heat-treated Ni-P/Ni-B-W coatings: a) as-plated
NiP/NiBW, b) heat treated at 300 °C, c) heat treated at 350 °C d) heat
treated at 650 °C
Figure 5: Surface morphology of NiP/NiBW dublex coatings detected
by SEM after potentiodynamic measurement: a) as-plated, b) heat
treated at 650 °C
samples were mechanically cleaned (up to 1200 grade)
and then soaked in trichloroethylene and cleaned with
detergent in an ultrasonic bath at 70 °C. Finally, sample
surfaces were activated in 30 % of volume fractions. HCl
for two minutes and then rinsed in distilled water. The
cleaned substrates were then soaked in Ni-P and Ni-B-W
electroless plating baths for two hours, respectively.
Reagent-grade chemicals and distilled water were
used for the preparation of all the electrolytes and the pH
was adjusted to 4.8±0.2 and 13.5±0.2 with H2SO4 or
NaOH in Ni-P and Ni-B-W plating baths, respectively.
The temperature of the baths was maintained at 90±2 °C
for Ni-P and 88±2 °C for Ni-B-W.
Deposition of Ni-P and Ni-B-W was carried out in an
aqueous bath containing 15 g/L NiSO4.6H2O, 26 g/L
Na2H2PO2·H2O, 13 g/L NaC2H3O2, 12 mL/L HF and
8 g/L NH4HF2 for the Ni-P coating, and 24 g/L NiCl2,
60 mL/L EDTA, 26.5 g/L KOH, 120 g/L NaBH4,
263 g/L NaOH, 2.6g/L PbWO4, 13 g/L EDTA and 40 g/L
Na2WO4·2H2O for Ni-B-W coatings, respectively.
The coated specimens were heat treated at 250 °C,
300 °C, 350 °C and 650 °C for 1 h each at a heating rate
of 5 °C/min in an air-circulated furnace, followed by
slow furnace cooling to room temperature.
The composition of the dublex coatings was deter-
mined by energy-dispersive X-ray analysis (EDX) using
a Link Analytical QX-2000 attached to the SEM appa-
ratus. The crystal structures of the dublex coatings were
examined using a Philips PW 3710 grazing incidence
x-ray diffractometer (Cu-K radiation). A Jeol
JSM-7000F FE-SEM was used to characterize the
microstructures and morphology of the coatings. The
corrosion behavior was investigated using a PGZ 301
Dynamic Voltammetry and VoltaMaster4 software.
Electrochemical experiments were carried out in a
3.5 % NaCl aqueous solution in a three-electrode cell at
room temperature. In this cell, a platinum electrode was
used as a counter electrode and a saturated calomel
electrode was used as a reference electrode; the dublex
coating was used as a working electrode and by masking
with silicon, only a 1 cm2 surface area was left on the
samples for exposure to the electrolyte. The dynamic
potential scanning technique was used for obtaining the
polarization curves of dublex coating samples. Prior to
the potentiodynamic measurement, samples were held
for approximately 15 min in the electrolyte and then
electrode potential was raised from -600mV to +200 mV
at a rate of 10 mV/min.
3 RESULTS
The XRD patterns of the as-plated and heat-treated
Ni-P/Ni-B-W coatings, and their chemical compositions
are in Figure 1 and in Table 1. SEM micrographs of
Ni-P/Ni-B-W coatings before and after the heat treat-
ments are in Figure 2.
Table 1: Chemical compositions of the Ni-P/Ni-B-W dublex coatings
Type of
coating Ni (w/%) P (w/%) B (w/%) W (w/%)
Ni-P 87.27 12.73 - -
Ni-B-W 87.45 - 7.86 4.69
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
838 Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Surface morphology of dublex coating a), b) as-plated
Ni-P/Ni-B-W, c) heat treated at 300 °C, d) heat treated at 650 °C
Figure 1: XRD spectra for a): as-plated Ni-P/Ni-B-W, b) 250 °C, c)
300 °C, d) 350 °C and e) 650 °C
The cross-sections of the Ni-P/Ni-B-W dublex coat-
ings are shown in Figure 3.
Potentiodynamic curves of as-plated and heat-treated
Ni-P/Ni-B-W coatings at different temperatures obtained
in a 3.5 % NaCl aqueous medium are shown in Figure 4.
The changes in the corrosion current density (Icorr)
values with heat treatment are given in Table 2.
Table 2: Corrosion resistance of as-plated and heat-treated
Ni-P/Ni-B-W dublex coatings at different temperatures in 3.5 % NaCl
solution
NiP/NiBW Icorr (μA/cm2) Ecorr (mV)
As-plated 1.55±0.59 -506±32
300 °C 1.70±0.41 -470±23
350 °C 0.71±0.20 -382 ±18
650 °C 0.57±0.09 -280±11
The surface morphologies of the as-plated sample
and heat-treated sample at 650 °C following potentio-
dynamic polarization measurements are shown in Fig-
ure 5.
The corrosion current density and corrosion potential
values obtained by Tafel interpolation are shown in
Table 3.
Table 3: Corrosion resistance of substrate and Ni-B-W and
Ni-P/Ni-B-W dublex coatings heat treated at 650 °C in 3.5 % NaCl
solution
Sample Icorr (μA/cm2) Ecorr (mV)
Carbon steel
substrate 4.59±1.01 -598±46
NiBW 3.40±0.35 -298±17
NiP/NiBW 0.57±0.09 -280±11
Figure 6 shows the polarization curves of the
Ni-B-W and dublex Ni-P/Ni-B-W coatings heat treated
at 650 °C and the steel substrate obtained in a 3.5 %
NaCl aqueous medium.
4 DISCUSSION
The Ni-P/Ni-B-W coating diffraction patterns up to
300 °C reveal a single broad peak indicating an amor-
phous coating structure under these conditions (Fig-
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842 839
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Cross-section of Ni-P/Ni-B-W dublex coating: a) as-plated,
b) heat treated at 250 °C c) heat treated at 350 °C, d) heat treated at
Figure 4: Polarization curves in a 3.5 % NaCl aqueous medium of
as-plated and heat-treated Ni-P/Ni-B-W coatings: a) as-plated
NiP/NiBW, b) heat treated at 300 °C, c) heat treated at 350 °C d) heat
treated at 650 °C
Figure 5: Surface morphology of NiP/NiBW dublex coatings detected
by SEM after potentiodynamic measurement: a) as-plated, b) heat
treated at 650 °C
ure 1). For heat-treatment temperatures up to 300 °C, the
dublex coating preserved its amorphous structure.
However, after heat treatment at and above 300 °C, the
diffraction patterns showed the formation of Ni3B inter-
metallic crystals. For the higher heat-treatment tempe-
ratures up to 650 °C, formation of the nickel borides
progressively increased. Drovosekov et al., reported
chemical bond formations between boron and tungsten
in the Ni-B-W coating system. However, the major
nickel borides formed during the heat treatments might
have prevented the formation of the B-W bond.9,15 In
other words, the quantity of the boron in the coating was
high enough to prevent the formation of Ni-W com-
pounds in the structure.16
SEM micrographs of the Ni-P/Ni-B-W coatings be-
fore and after the heat treatments showed the formation
of typical spherical nodular structures (Figure 2). Each
nodule appears to be formed by the unification of many
grains existing as colonies (Figure 2b), similar to the
case suggested by S. Ruan et al.17 Some microcracks
were noticeable on the surfaces of the as-plated
Ni-P/Ni-B-W samples. The microcracks may be due to
the internal stress caused by the relatively larger tungsten
atoms in the coating structure.18,19
With the heat treatments, the structure progressively
crystallized and, as consequence, a substantial reduction
of the microcracks was noticed (Figure 2c and 2d). In a
Ni-B-W coating system, W atoms may selectively exist
at the boundaries of the nodules, from there, they may
segregates to the coating surface, during the heat treat-
ments.16,17 This process may decrease the internal stress
inside the coating, resulting in a substantial decrease of
the microcracks on the coating surface.
The thickness of the Ni-P coating in the dublex
coating system was measured as 2 μm, while the average
thickness of the Ni-B-W (the other coating in the
system) was measured as 11 μm (Figure 3). On the other
hand, it was obvious that the nodular structure seen in
the surface morphology of the Ni-B-W coating showed
columnar growth starting from the Ni-P layer to the
topmost surface of the coating (Figure 3a). Furthermore,
it was observed that adherence of the Ni-P coating to the
substrate was good and there was no pore formation
between the substrate and Ni-P coating. Pores seen in the
interface of Ni-P/Ni-B-W, in addition to pores partially
observed inside the Ni-B-W coating had a columnar
structure and they were considered to have been formed
due to the capturing of released hydrogen – even if in
small amounts – during the oxidation reaction of sodium
borohydride used as a the boron source.5 The pores
disappear with increasing heat-treatment temperatures.
This may be a result of the diffusion of hydrogen atoms
to the surface. They may promote better crystallization
of the coating and ease the formation of the Ni3B and
Ni2B phases (Figure 3b and 3c). After 650 °C heat
treatment the major phase of the coating was noticed to
be Ni3B. At that stage adherence between the coatings
reached its best level, while the Ni-B-W coating showed
columnar grains with a denser structure (Figure 3d).
As Figure 4, a difference of approximately 220 mV
was detected between corrosion potentials (Ecorr) of
as-plated and 650 °C heat-treated dublex coatings. This
is the shifting of Ecorr values of the coatings toward
nobler values with increased heat treatments applied to
the coatings. The anodic branch of the polarization curve
of the coating before heat treatment showed the break-
down potential at values around -310 mV (Figure 4a).
But the current density increased at a constant rate,
indicating that the unstable passive film formation on the
coating has disappeared rapidly. For heat treatment at
300 °C, the breakdown potential was observed to shift
toward -220 mV. At 300 °C, which is the onset tem-
perature of crystallization, corrosion potential shifted
towards a positive end approximately 40 mV, compared
to the as-plated sample. At 350 °C, crystallization was
nearly completed and the corrosion potential had shifted
towards more positive values. However, the breakdown
potential was observed (as with other samples) at around
-180 mV. Nonetheless, it was detected that the current
density of this sample became higher as the potential
value increased compared with that for the to as-plated
and heat-treated samples at 300 °C. For the heat treat-
ment at 650 °C, a passive film formation on the coating
surface was seen at around -190 mV, suggesting that
corrosion rate was restrained at such potential range.
However, as this protective film began to dissolve at
around -90 mV, a steep increase in the current density
was observed.
The major reason for the difference between the
corrosion resistance of as-plated and heat-treated
Ni-P/Ni-B-W coatings exposed to chloride ions was
undoubtedly associated with the transformation of the
microstructure from an amorphous to a crystalline mor-
phology. The main reason for having a higher Ecorr value
the heat-treated coating at 300 °C (where an amorphous
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
840 Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Polarization curves: a) steel substrate, b) Ni-B-W coating
heat treated at 650 °C, c) dublex Ni-P/Ni-B-W coating heat treated at
650 °C
structure still prevailed) compared to the as-plated
sample was the result of the borides formed during the
onset of crystallization. The heat-treated sample at
350 °C (where the microstructure had completely been
transformed to a crystalline structure) had a higher
corrosion resistance compared to the sample heat-treated
at 300 °C.20 On the other hand, it is a known fact that
corrosion resistance generally increases with addition of
tungsten to a electroless nickel-boron coating system,
because of the tendency to form an oxide film on the
surfaces.7,18 However, as seen in Figure 4, the corrosion
resistance of the as-plated Ni-P/Ni-B-W sample having a
nodular structure did not increase, with the addition of
tungsten. This is because the tungsten atoms were stuck
in the colony boundaries, not to be segregated to the
surfaces during the corrosion process and therefore,
could not be oxidized to promote the formation of a
mixed oxide film. However, depending on the increased
heat treatment temperatures segregation of W atoms
toward surfaces at 650 °C was more likely to occur. The
mixed oxide film a passive zone formed this way may be
the main reason for the substantial decrease of the
current density as observed here.
It can be seen that uniform corrosion occurred on the
surface of the as-plated sample in the amorphous
structure that has the lowest corrosion resistance (Fig-
ure 5a). Conversely, on the surface of the heat-treated
sample at 650 °C complete crystallization were noted
and the pits (indicated by arrows) in the grains within the
colonies were observed (Figure 5b).
To interpret the effect of the dublex coating on
corrosion resistance, the potentiodynamic polarization
measurements of the Ni-P/Ni-B-W and Ni-B-W
coatings, as well as which were deposited under the
same process conditions and heat-treated at 650 °C, and
carbon steel substrate were studied (Table 3).
As can be observed from Figure 6, while the Ni-B-W
coating and the dublex coating almost had the same
corrosion potential, more positive values at about 300
mV, compared with that of the substrate. Among the
coating studied the dublex coating has the lowest corro-
sion current density. The anodic branches of polarization
curves of both the dublex and the Ni-B-W coatings have
a passive zone, and both coatings had an Epit value of
approximately -90 mV. Additionally, the current density
of the Ni-B-W coating at a constant rate increases,
depending on the increased potential value. On the other
hand, the anodic branch of the dublex coating indicated
that an entrance to a second passive zone at approxi-
mately 20 mV. This passive zone may be related to the
existence of Ni-P in the dublex coating. A similar inter-
pretation was reported by Zhang et al. for the corrosion
of Ni-B-W surfaces.14 In general, corrosion starts at the
outermost surface of the substrate and proceed inward.
However, the stable passive film formed on the Ni-P
layer on the substrate may act as a barrier for corrosion
propagation, causing the dublex coating to have better
corrosion resistance than the single-layer Ni-B-W
coating.
5 CONCLUSIONS
The effect of heat treatments on the corrosion resis-
tance of a Ni-P/Ni-B-W dublex coating deposited on
carbon steel was studied. The Ni-P/Ni-B-W coatings
have an amorphous structure for a heat treatments up to
300 °C. The coatings start to crystallize at about 350 °C
and completed at about 650 °C with the major phases
being the nickel borides. The heat treatment, besides
causing nickel boride formation within Ni-B-W, also
caused obtaining a denser Ni-P/Ni-B-W coating when
increasing heat treatment temperature.
The corrosion resistance of the dublex coating in-
creases substantially with increasing the heat treatment
temperature. The dublex coating and the single-layer
Ni-B-W coating heat-treated at 650 °C indicated higher
corrosion resistance as compared to that of steel sub-
strate. When these two coatings were compared in terms
of corrosion resistance, the dublex coating showed better
performance.
6 REFERENCES
1 A. Brenner, G. Riddell, Nickel Plating on Steel by Chemical
Reduction, J. Res. Nat. Bur. Std., 37 (1946), 31–35
2 J. Sudagar, J. Lian, W. Sha, Electroless nickel, alloy, composite and
nano coatings-A critical review, J. Alloys Compd., 571(2013),
183–204, doi:10.1016/j.jallcom.2013.03.107
3 Y. F. Shen, W.Y. Xue, Z.Y. Liu, L. Zuo, Nanoscratching Deformation
and Fracture Toughness of Electroless Ni-P Coatings, Surf. Coat.
Technol., 205 (2010), 632–640, doi:10.1016/j.surfcoat.2010.07.066
4 K. M. Gorbunova, M.V. Ivanov, V. P. Moissev, Electroless Depo-
sition of Nickel-Boron Alloys – Mechanýsm of Process, Structure,
and Some Propertýes of Deposits, J. Electrochem. Soc., 120 (1973),
613–618, doi:10.1149/1.2403514
5 R. N. Duncan, T. L. Arney, Operatýon and Use of Sodium-Boro-
hydride-Reduced Electroless Nickel, Plat. Surf. Finish., 71 (1984)
12, 49–54
6 R. A. C. Santana, S. Prasad, A. R. N. Campos, F. O. Arau´ jo, G. P.
DA Silva, P. Lýma-Neto, Electrodeposition and Mechanical Pro-
perties of Ni–W–B Composites from Tartrate Bath, Prot Met Phys
Chem., 46 (2010) 1, 117–122, doi:10.1134/S207020511001017X
7 M. G. Hosseini, M. Abdolmaleki , H. Ebrahimzadeh, S. A. Seyed
Sadjadi, Effect of 2-Butyne-1, 4-Diol on the Nanostructure and
Corrosion Resistance Properties of Electrodeposited Ni-W-B
Coatings, Int. J. Electrochem. Sci., 6 (2011) 4, 1189–1205
8 T. Osaka, N. Takano, T. Kurokawa, T. Kaneko, K. Ueno, Characteri-
zation of chemically-deposited NiB and NiWB thin films as a
capping layer for ULSI application, Surf. Coat. Technol., 169–170
(2003), 124–127, doi:10.1016/S0257-8972(03)00186-5
9 G. Graef, K. Anderson, J. Groza, A. Palazoglu, Phase evolution in
electrodeposited Ni-W-B alloy, Mater. Sci. Eng., B41 (1996),
253–257, doi:10.1016/S0921-5107(96)01656-X
10 E. Lassner, W. D. Schubert, Tungsten: Properties, Chemistry, Tech-
nology of the Element, Alloys, and Chemical Compounds, Kluwer
Academic, New York, 1999
11 E. Beltowska-Lehman, Kinetics of Induced Electrodeposition of
Alloys Containing Mo from Citrate Solutions, Phys. Stat. Sol., 5
(2008) 11, 3514–3520, doi:10.1002/pssc.200779404
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842 841
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
ure 1). For heat-treatment temperatures up to 300 °C, the
dublex coating preserved its amorphous structure.
However, after heat treatment at and above 300 °C, the
diffraction patterns showed the formation of Ni3B inter-
metallic crystals. For the higher heat-treatment tempe-
ratures up to 650 °C, formation of the nickel borides
progressively increased. Drovosekov et al., reported
chemical bond formations between boron and tungsten
in the Ni-B-W coating system. However, the major
nickel borides formed during the heat treatments might
have prevented the formation of the B-W bond.9,15 In
other words, the quantity of the boron in the coating was
high enough to prevent the formation of Ni-W com-
pounds in the structure.16
SEM micrographs of the Ni-P/Ni-B-W coatings be-
fore and after the heat treatments showed the formation
of typical spherical nodular structures (Figure 2). Each
nodule appears to be formed by the unification of many
grains existing as colonies (Figure 2b), similar to the
case suggested by S. Ruan et al.17 Some microcracks
were noticeable on the surfaces of the as-plated
Ni-P/Ni-B-W samples. The microcracks may be due to
the internal stress caused by the relatively larger tungsten
atoms in the coating structure.18,19
With the heat treatments, the structure progressively
crystallized and, as consequence, a substantial reduction
of the microcracks was noticed (Figure 2c and 2d). In a
Ni-B-W coating system, W atoms may selectively exist
at the boundaries of the nodules, from there, they may
segregates to the coating surface, during the heat treat-
ments.16,17 This process may decrease the internal stress
inside the coating, resulting in a substantial decrease of
the microcracks on the coating surface.
The thickness of the Ni-P coating in the dublex
coating system was measured as 2 μm, while the average
thickness of the Ni-B-W (the other coating in the
system) was measured as 11 μm (Figure 3). On the other
hand, it was obvious that the nodular structure seen in
the surface morphology of the Ni-B-W coating showed
columnar growth starting from the Ni-P layer to the
topmost surface of the coating (Figure 3a). Furthermore,
it was observed that adherence of the Ni-P coating to the
substrate was good and there was no pore formation
between the substrate and Ni-P coating. Pores seen in the
interface of Ni-P/Ni-B-W, in addition to pores partially
observed inside the Ni-B-W coating had a columnar
structure and they were considered to have been formed
due to the capturing of released hydrogen – even if in
small amounts – during the oxidation reaction of sodium
borohydride used as a the boron source.5 The pores
disappear with increasing heat-treatment temperatures.
This may be a result of the diffusion of hydrogen atoms
to the surface. They may promote better crystallization
of the coating and ease the formation of the Ni3B and
Ni2B phases (Figure 3b and 3c). After 650 °C heat
treatment the major phase of the coating was noticed to
be Ni3B. At that stage adherence between the coatings
reached its best level, while the Ni-B-W coating showed
columnar grains with a denser structure (Figure 3d).
As Figure 4, a difference of approximately 220 mV
was detected between corrosion potentials (Ecorr) of
as-plated and 650 °C heat-treated dublex coatings. This
is the shifting of Ecorr values of the coatings toward
nobler values with increased heat treatments applied to
the coatings. The anodic branch of the polarization curve
of the coating before heat treatment showed the break-
down potential at values around -310 mV (Figure 4a).
But the current density increased at a constant rate,
indicating that the unstable passive film formation on the
coating has disappeared rapidly. For heat treatment at
300 °C, the breakdown potential was observed to shift
toward -220 mV. At 300 °C, which is the onset tem-
perature of crystallization, corrosion potential shifted
towards a positive end approximately 40 mV, compared
to the as-plated sample. At 350 °C, crystallization was
nearly completed and the corrosion potential had shifted
towards more positive values. However, the breakdown
potential was observed (as with other samples) at around
-180 mV. Nonetheless, it was detected that the current
density of this sample became higher as the potential
value increased compared with that for the to as-plated
and heat-treated samples at 300 °C. For the heat treat-
ment at 650 °C, a passive film formation on the coating
surface was seen at around -190 mV, suggesting that
corrosion rate was restrained at such potential range.
However, as this protective film began to dissolve at
around -90 mV, a steep increase in the current density
was observed.
The major reason for the difference between the
corrosion resistance of as-plated and heat-treated
Ni-P/Ni-B-W coatings exposed to chloride ions was
undoubtedly associated with the transformation of the
microstructure from an amorphous to a crystalline mor-
phology. The main reason for having a higher Ecorr value
the heat-treated coating at 300 °C (where an amorphous
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
840 Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Polarization curves: a) steel substrate, b) Ni-B-W coating
heat treated at 650 °C, c) dublex Ni-P/Ni-B-W coating heat treated at
650 °C
structure still prevailed) compared to the as-plated
sample was the result of the borides formed during the
onset of crystallization. The heat-treated sample at
350 °C (where the microstructure had completely been
transformed to a crystalline structure) had a higher
corrosion resistance compared to the sample heat-treated
at 300 °C.20 On the other hand, it is a known fact that
corrosion resistance generally increases with addition of
tungsten to a electroless nickel-boron coating system,
because of the tendency to form an oxide film on the
surfaces.7,18 However, as seen in Figure 4, the corrosion
resistance of the as-plated Ni-P/Ni-B-W sample having a
nodular structure did not increase, with the addition of
tungsten. This is because the tungsten atoms were stuck
in the colony boundaries, not to be segregated to the
surfaces during the corrosion process and therefore,
could not be oxidized to promote the formation of a
mixed oxide film. However, depending on the increased
heat treatment temperatures segregation of W atoms
toward surfaces at 650 °C was more likely to occur. The
mixed oxide film a passive zone formed this way may be
the main reason for the substantial decrease of the
current density as observed here.
It can be seen that uniform corrosion occurred on the
surface of the as-plated sample in the amorphous
structure that has the lowest corrosion resistance (Fig-
ure 5a). Conversely, on the surface of the heat-treated
sample at 650 °C complete crystallization were noted
and the pits (indicated by arrows) in the grains within the
colonies were observed (Figure 5b).
To interpret the effect of the dublex coating on
corrosion resistance, the potentiodynamic polarization
measurements of the Ni-P/Ni-B-W and Ni-B-W
coatings, as well as which were deposited under the
same process conditions and heat-treated at 650 °C, and
carbon steel substrate were studied (Table 3).
As can be observed from Figure 6, while the Ni-B-W
coating and the dublex coating almost had the same
corrosion potential, more positive values at about 300
mV, compared with that of the substrate. Among the
coating studied the dublex coating has the lowest corro-
sion current density. The anodic branches of polarization
curves of both the dublex and the Ni-B-W coatings have
a passive zone, and both coatings had an Epit value of
approximately -90 mV. Additionally, the current density
of the Ni-B-W coating at a constant rate increases,
depending on the increased potential value. On the other
hand, the anodic branch of the dublex coating indicated
that an entrance to a second passive zone at approxi-
mately 20 mV. This passive zone may be related to the
existence of Ni-P in the dublex coating. A similar inter-
pretation was reported by Zhang et al. for the corrosion
of Ni-B-W surfaces.14 In general, corrosion starts at the
outermost surface of the substrate and proceed inward.
However, the stable passive film formed on the Ni-P
layer on the substrate may act as a barrier for corrosion
propagation, causing the dublex coating to have better
corrosion resistance than the single-layer Ni-B-W
coating.
5 CONCLUSIONS
The effect of heat treatments on the corrosion resis-
tance of a Ni-P/Ni-B-W dublex coating deposited on
carbon steel was studied. The Ni-P/Ni-B-W coatings
have an amorphous structure for a heat treatments up to
300 °C. The coatings start to crystallize at about 350 °C
and completed at about 650 °C with the major phases
being the nickel borides. The heat treatment, besides
causing nickel boride formation within Ni-B-W, also
caused obtaining a denser Ni-P/Ni-B-W coating when
increasing heat treatment temperature.
The corrosion resistance of the dublex coating in-
creases substantially with increasing the heat treatment
temperature. The dublex coating and the single-layer
Ni-B-W coating heat-treated at 650 °C indicated higher
corrosion resistance as compared to that of steel sub-
strate. When these two coatings were compared in terms
of corrosion resistance, the dublex coating showed better
performance.
6 REFERENCES
1 A. Brenner, G. Riddell, Nickel Plating on Steel by Chemical
Reduction, J. Res. Nat. Bur. Std., 37 (1946), 31–35
2 J. Sudagar, J. Lian, W. Sha, Electroless nickel, alloy, composite and
nano coatings-A critical review, J. Alloys Compd., 571(2013),
183–204, doi:10.1016/j.jallcom.2013.03.107
3 Y. F. Shen, W.Y. Xue, Z.Y. Liu, L. Zuo, Nanoscratching Deformation
and Fracture Toughness of Electroless Ni-P Coatings, Surf. Coat.
Technol., 205 (2010), 632–640, doi:10.1016/j.surfcoat.2010.07.066
4 K. M. Gorbunova, M.V. Ivanov, V. P. Moissev, Electroless Depo-
sition of Nickel-Boron Alloys – Mechanýsm of Process, Structure,
and Some Propertýes of Deposits, J. Electrochem. Soc., 120 (1973),
613–618, doi:10.1149/1.2403514
5 R. N. Duncan, T. L. Arney, Operatýon and Use of Sodium-Boro-
hydride-Reduced Electroless Nickel, Plat. Surf. Finish., 71 (1984)
12, 49–54
6 R. A. C. Santana, S. Prasad, A. R. N. Campos, F. O. Arau´ jo, G. P.
DA Silva, P. Lýma-Neto, Electrodeposition and Mechanical Pro-
perties of Ni–W–B Composites from Tartrate Bath, Prot Met Phys
Chem., 46 (2010) 1, 117–122, doi:10.1134/S207020511001017X
7 M. G. Hosseini, M. Abdolmaleki , H. Ebrahimzadeh, S. A. Seyed
Sadjadi, Effect of 2-Butyne-1, 4-Diol on the Nanostructure and
Corrosion Resistance Properties of Electrodeposited Ni-W-B
Coatings, Int. J. Electrochem. Sci., 6 (2011) 4, 1189–1205
8 T. Osaka, N. Takano, T. Kurokawa, T. Kaneko, K. Ueno, Characteri-
zation of chemically-deposited NiB and NiWB thin films as a
capping layer for ULSI application, Surf. Coat. Technol., 169–170
(2003), 124–127, doi:10.1016/S0257-8972(03)00186-5
9 G. Graef, K. Anderson, J. Groza, A. Palazoglu, Phase evolution in
electrodeposited Ni-W-B alloy, Mater. Sci. Eng., B41 (1996),
253–257, doi:10.1016/S0921-5107(96)01656-X
10 E. Lassner, W. D. Schubert, Tungsten: Properties, Chemistry, Tech-
nology of the Element, Alloys, and Chemical Compounds, Kluwer
Academic, New York, 1999
11 E. Beltowska-Lehman, Kinetics of Induced Electrodeposition of
Alloys Containing Mo from Citrate Solutions, Phys. Stat. Sol., 5
(2008) 11, 3514–3520, doi:10.1002/pssc.200779404
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842 841
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
12 A. I. Aydeniz, A. Göksenli, G. Dil, F. Muhaffel, C. Calli, B. Yüksel,
Electroless Ni-B-W Coatings for Improving Hardness, Wear And
Corrosion Resistance, MTAEC9, 47 (2013) 6, 803–806
13 T. S. N. Sankara Narayanan, K. Krishnaveni, S. K. Seshadri, Elec-
troless Ni–P/Ni–B Dublex Coatings: Preparation and Evaluation of
Microhardness, Wear and Corrosion Resistance, Mater. Chem. Phys.,
82 (2003), 771–779, doi:10.1016/S0254-0584(03)00390-0
14 W. X. Zhang, Z. H. Jiang, G. Y. Li, Q. Jiang, J. S. Lian, Electroless
NiP/NiB Dublex Coatings For Improving The Hardness and The
Corrosion Resistance of AZ91 D Magnesium Alloy, Appl. Surf. Sci.,
254 (2008), 4949–4955, doi:10.1016/j.apsusc.2008.01.144
15 A. B. Drovosekov, M.V. Ivanov, V. M. Krutskikh, E. N. Lubnin, Y.
M. Polukarov, Chemically Deposited Ni-W-B Coatings: Composi-
tion, Structure and Properties, Prot. Met., 41(2005) 1, 55–62,
doi:10.1007/s11124-005-0008-1
16 F. Z. Yang, Z. H. Ma, L. Huang, S. K. Xu, S. M. Zhou, Electro-
deposition and Properties of an Amorphous Ni-W-B Alloy before
and after Heat Treatment, Chin. J. Chem., 24 (2006), 114–118,
doi:10.1002/cjoc.200690004
17 S. Ruan, C. A. Schu, Mesoscale Structure and Segregation in
Electrodeposited Nanocrystalline Alloys, Scripta Mater., 59 (2008)
11, 1218–1221, doi:10.1016/j.scriptamat.2008.08.010
18 R. A.C. Santana, S. Prasad, A. R. N. Campos, F. O. Arau´ jo, G.P.
Sýlva, P. Lima-Neto, Electrodeposition and Corrosion Behaviour of a
Ni–W–B Amorphous Alloy, J. Appl. Electrochem., 36 (2006),
105–113, doi:10.1007/s10800-005-9046-2
19 R. A. Yildiz, A. Göksenli, B. Yüksel, F. Muhaffel, A. Aydeniz, Effect
of Annealing Temperature on The Corrosion Resistance of Elec-
troless Produced Ni-B-W Coatings, Adv Mat Res., 652 (2013),
263–268, doi:10.4028/www.scientific.net/AMR.651.263
20 Z. Z. Hamid, H. B. Hassan, A. M. Attyia, Influence of Deposition
Temperature and Heat Treatment on The Performance of Electroless
Ni–B Films, Surf. Coat. Technol., 205 (2010) 7, 2348–2354,
doi:10.1016/j.surfcoat.2010.09.025
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
842 Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
J. ZHE et al.: SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW STRESSES ...
843–847
SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW
STRESSES AND AN INSTANTANEOUS-STRESS-CHANGE
TESTING
KRATKOTRAJNO LEZENJE TOPLOTNO ODPORNEGA JEKLA P91
PRI NIZKIH NAPETOSTIH IN NENADNI MENJAVI NAPETOSTI
OBREMENJEVANJA
Jiang Zhe, Shen Junjie, Zhao Pengshuo
Tianjin University of Technology, Tianjin Key Laboratory for Advanced Mechatronic System Design and Intelligent Control, National
Demonstration Center for Experimental Mechanical and Electrical Engineering Education, 391 Binshui Road, Xiqing District, Tianjin, China
sjj1982428@sina.com, j211209977@live.com
Prejem rokopisa – received: 2016-1'-20; sprejem za objavo – accepted for publication: 2017-03-21
doi:10.17222/mit.2016.305
For the short-term creep behavior to be evaluated and the creep mechanism of P91 heat-resistant steels at low stresses and high
temperatures to be clarified, stress-change testing was conducted with a “helicoidal-spring creep test” demonstrating a high
strain resolution. The creep deformation consists of the primary creep stage, whereas no secondary creep stage was observed.
Blackburn’s law was suggested to be the best choice for a short-term-creep-behavior description because it provides a good
representation of an experimental creep curve. An anelastic backflow at a low stress was confirmed, following a high reduction
in the stress. The absolute value of the instantaneous strain for a load increase was equal to the value for a load decrease and the
creep of the P91 steels at low stresses might have been controlled by the viscous glide of dislocations.
Keywords: P91 heat-resistant steel, creep, anelastic, stress change
Da bi bilo mo~ oceniti kratkotrajno lezenje in njegov mehanizem toplotno obstojnega jekla P91 pri nizkih napetostih in visokih
temperaturah, so bili izvedeni preizkusi spremembe napetosti z uporabo t.i. testa spiralne vzmeti, ki omogo~a veliko lo~ljivost
deformacije. Deformacija lezenja sestoji iz primarne in sekundarne faze lezenja, vendar sekundarne faze niso analizirali oz.
opazovali. Blackburnov zakon je najbolj uporaben zakon za opis obna{anja jekla med kratkotrajnim lezenjem, ker se z njim
najbolj pribli`amo rezultatom, dobljenim z eksperimenti. Neelasti~ni povratni tok pri nizki napetosti je bil potrjen z ve~jim
zmanj{anjem napetosti. Zaradi absolutne vrednosti trenutne deformacije za dano obremenitev je bil narastek deformacije enak
zmanj{anju obremenitve, medtem ko je lezenje pri nizkih napetostih jekla P91 kontrolirano z viskoznim drsenjem dislokacij.
Klju~ne besede: P91 jeklo za delo v vro~em, lezenje, neelasti~nost, sprememba napetosti
1 INTRODUCTION
The P91 ferritic heat-resistant steel is utilized for the
material production of thermal-power-plant components,
which are exposed to elevated temperatures for extensive
periods. This requires that the P91 structural materials
resist creep deformation at both high temperatures and
low stresses. The creep strength obtained with the extra-
polation method based on the data of the creep deforma-
tion at a high stress is higher than the true creep strength
under these conditions (at low stresses).1,2 The creep-
strength degradation is not only based on the microstruc-
tural degradation3–6, but it is also related to the change in
the creep mechanism.7–9 The creep-mechanism clarifica-
tion has to depend on the instantaneous creep behavior
because the long-term creep is associated with the micro-
structural degradation. The stress-change test constitutes
an effective method for a creep-deformation-mechanism
clarification based on the instantaneous-strain and
creep-rate-variation evaluation at the stress changes
during the creep testing.10–12
At a high strain rate (a high stress) and a high tem-
perature, a sudden-stress-change experiment has been
widely utilized for pure metals and solution-strengthened
alloys.10–12 In contrast, at a low strain rate 0, especially
the ultra-low 0, lower than 10–10 s–1, it is usually impossi-
ble for a conventional tension-creep technique to dist-
inguish the instantaneous plastic strain from the total
strain under a sudden stress change, due to the unsatis-
factory strain resolution, existing between 10–6 and 10–5.
As an additional creep technique, the helicoidal-spring
creep test based on torsion deformation provides a
significantly high strain resolution, even up to 10–9.13
Due to this high strain resolution, it is possible for very
low instantaneous strains to be measured during a
sudden stress change.14
In the present study, the instantaneous creep of the
P91 ferritic heat-resistant steels was studied during
sudden-stress-change experiments using the helicoidal-
spring-specimen technique for the creep mechanism
under both a low stress and a high temperature.
Materiali in tehnologije / Materials and technology 51 (2017) 5, 843–847 843
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 621.9.011:620.172.24:691.714 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)843(2017)
12 A. I. Aydeniz, A. Göksenli, G. Dil, F. Muhaffel, C. Calli, B. Yüksel,
Electroless Ni-B-W Coatings for Improving Hardness, Wear And
Corrosion Resistance, MTAEC9, 47 (2013) 6, 803–806
13 T. S. N. Sankara Narayanan, K. Krishnaveni, S. K. Seshadri, Elec-
troless Ni–P/Ni–B Dublex Coatings: Preparation and Evaluation of
Microhardness, Wear and Corrosion Resistance, Mater. Chem. Phys.,
82 (2003), 771–779, doi:10.1016/S0254-0584(03)00390-0
14 W. X. Zhang, Z. H. Jiang, G. Y. Li, Q. Jiang, J. S. Lian, Electroless
NiP/NiB Dublex Coatings For Improving The Hardness and The
Corrosion Resistance of AZ91 D Magnesium Alloy, Appl. Surf. Sci.,
254 (2008), 4949–4955, doi:10.1016/j.apsusc.2008.01.144
15 A. B. Drovosekov, M.V. Ivanov, V. M. Krutskikh, E. N. Lubnin, Y.
M. Polukarov, Chemically Deposited Ni-W-B Coatings: Composi-
tion, Structure and Properties, Prot. Met., 41(2005) 1, 55–62,
doi:10.1007/s11124-005-0008-1
16 F. Z. Yang, Z. H. Ma, L. Huang, S. K. Xu, S. M. Zhou, Electro-
deposition and Properties of an Amorphous Ni-W-B Alloy before
and after Heat Treatment, Chin. J. Chem., 24 (2006), 114–118,
doi:10.1002/cjoc.200690004
17 S. Ruan, C. A. Schu, Mesoscale Structure and Segregation in
Electrodeposited Nanocrystalline Alloys, Scripta Mater., 59 (2008)
11, 1218–1221, doi:10.1016/j.scriptamat.2008.08.010
18 R. A.C. Santana, S. Prasad, A. R. N. Campos, F. O. Arau´ jo, G.P.
Sýlva, P. Lima-Neto, Electrodeposition and Corrosion Behaviour of a
Ni–W–B Amorphous Alloy, J. Appl. Electrochem., 36 (2006),
105–113, doi:10.1007/s10800-005-9046-2
19 R. A. Yildiz, A. Göksenli, B. Yüksel, F. Muhaffel, A. Aydeniz, Effect
of Annealing Temperature on The Corrosion Resistance of Elec-
troless Produced Ni-B-W Coatings, Adv Mat Res., 652 (2013),
263–268, doi:10.4028/www.scientific.net/AMR.651.263
20 Z. Z. Hamid, H. B. Hassan, A. M. Attyia, Influence of Deposition
Temperature and Heat Treatment on The Performance of Electroless
Ni–B Films, Surf. Coat. Technol., 205 (2010) 7, 2348–2354,
doi:10.1016/j.surfcoat.2010.09.025
B. YÜKSEL et al.: CORROSION RESISTANCE OF AS-PLATED AND HEAT-TREATED ELECTROLESS ...
842 Materiali in tehnologije / Materials and technology 51 (2017) 5, 837–842
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
J. ZHE et al.: SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW STRESSES ...
843–847
SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW
STRESSES AND AN INSTANTANEOUS-STRESS-CHANGE
TESTING
KRATKOTRAJNO LEZENJE TOPLOTNO ODPORNEGA JEKLA P91
PRI NIZKIH NAPETOSTIH IN NENADNI MENJAVI NAPETOSTI
OBREMENJEVANJA
Jiang Zhe, Shen Junjie, Zhao Pengshuo
Tianjin University of Technology, Tianjin Key Laboratory for Advanced Mechatronic System Design and Intelligent Control, National
Demonstration Center for Experimental Mechanical and Electrical Engineering Education, 391 Binshui Road, Xiqing District, Tianjin, China
sjj1982428@sina.com, j211209977@live.com
Prejem rokopisa – received: 2016-1'-20; sprejem za objavo – accepted for publication: 2017-03-21
doi:10.17222/mit.2016.305
For the short-term creep behavior to be evaluated and the creep mechanism of P91 heat-resistant steels at low stresses and high
temperatures to be clarified, stress-change testing was conducted with a “helicoidal-spring creep test” demonstrating a high
strain resolution. The creep deformation consists of the primary creep stage, whereas no secondary creep stage was observed.
Blackburn’s law was suggested to be the best choice for a short-term-creep-behavior description because it provides a good
representation of an experimental creep curve. An anelastic backflow at a low stress was confirmed, following a high reduction
in the stress. The absolute value of the instantaneous strain for a load increase was equal to the value for a load decrease and the
creep of the P91 steels at low stresses might have been controlled by the viscous glide of dislocations.
Keywords: P91 heat-resistant steel, creep, anelastic, stress change
Da bi bilo mo~ oceniti kratkotrajno lezenje in njegov mehanizem toplotno obstojnega jekla P91 pri nizkih napetostih in visokih
temperaturah, so bili izvedeni preizkusi spremembe napetosti z uporabo t.i. testa spiralne vzmeti, ki omogo~a veliko lo~ljivost
deformacije. Deformacija lezenja sestoji iz primarne in sekundarne faze lezenja, vendar sekundarne faze niso analizirali oz.
opazovali. Blackburnov zakon je najbolj uporaben zakon za opis obna{anja jekla med kratkotrajnim lezenjem, ker se z njim
najbolj pribli`amo rezultatom, dobljenim z eksperimenti. Neelasti~ni povratni tok pri nizki napetosti je bil potrjen z ve~jim
zmanj{anjem napetosti. Zaradi absolutne vrednosti trenutne deformacije za dano obremenitev je bil narastek deformacije enak
zmanj{anju obremenitve, medtem ko je lezenje pri nizkih napetostih jekla P91 kontrolirano z viskoznim drsenjem dislokacij.
Klju~ne besede: P91 jeklo za delo v vro~em, lezenje, neelasti~nost, sprememba napetosti
1 INTRODUCTION
The P91 ferritic heat-resistant steel is utilized for the
material production of thermal-power-plant components,
which are exposed to elevated temperatures for extensive
periods. This requires that the P91 structural materials
resist creep deformation at both high temperatures and
low stresses. The creep strength obtained with the extra-
polation method based on the data of the creep deforma-
tion at a high stress is higher than the true creep strength
under these conditions (at low stresses).1,2 The creep-
strength degradation is not only based on the microstruc-
tural degradation3–6, but it is also related to the change in
the creep mechanism.7–9 The creep-mechanism clarifica-
tion has to depend on the instantaneous creep behavior
because the long-term creep is associated with the micro-
structural degradation. The stress-change test constitutes
an effective method for a creep-deformation-mechanism
clarification based on the instantaneous-strain and
creep-rate-variation evaluation at the stress changes
during the creep testing.10–12
At a high strain rate (a high stress) and a high tem-
perature, a sudden-stress-change experiment has been
widely utilized for pure metals and solution-strengthened
alloys.10–12 In contrast, at a low strain rate 0, especially
the ultra-low 0, lower than 10–10 s–1, it is usually impossi-
ble for a conventional tension-creep technique to dist-
inguish the instantaneous plastic strain from the total
strain under a sudden stress change, due to the unsatis-
factory strain resolution, existing between 10–6 and 10–5.
As an additional creep technique, the helicoidal-spring
creep test based on torsion deformation provides a
significantly high strain resolution, even up to 10–9.13
Due to this high strain resolution, it is possible for very
low instantaneous strains to be measured during a
sudden stress change.14
In the present study, the instantaneous creep of the
P91 ferritic heat-resistant steels was studied during
sudden-stress-change experiments using the helicoidal-
spring-specimen technique for the creep mechanism
under both a low stress and a high temperature.
Materiali in tehnologije / Materials and technology 51 (2017) 5, 843–847 843
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 621.9.011:620.172.24:691.714 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)843(2017)
2 EXPERIMENTAL PART
2.1 Materials
The P91 heat-resistant steel was utilized in this expe-
riment. The chemical components are presented in Table 1.
2.2 Specimen processing
The steel was processed into helicoidal-spring speci-
mens using the method of high-speed wire-electrode
cutting machining; the outer and inner diameters of the
specimens were 12 mm and 8 mm, respectively. The
cross-section of the helicoidal-spring specimens was a
rectangle. The side length of the rectangle was 2 mm, as
presented in Figure 1.
The helicoidal-spring specimens were annealed at
1313 K for 60 min and consequently tempered at 1033 K
for 60 min.
2.3 Test methods
Figure 2 presents the experimental apparatus, which
includes an electric furnace with three adjacent heaters, a
high-precision optical micrometer, a load, a weight and a
system for receiving and processing information. By
measuring the helicoidal-spring-specimen pitch change,
the creep curves could be obtained. The test-apparatus
details were previously reported.14
During the testing, the number of weights was in-
creased or decreased for the instantaneous stress change
to be observed. The weight was suspended under a load
with a flammable thread. In order for a load to be
increased, the load was hooked quickly. For a load to be
decreased, the flammable thread was burned.
The test was performed at a temperature of 923 K
and a slight oscillation. The specimen was loaded sub-
sequently for five days when the strain rate was down to
10–9–10–8 s–1. Also, the stress consequently increased or
decreased at various applied levels and the instantaneous
strain was measured. The following equations15,16 were
utilized for calculating the mean surface shear stress, ,
and the surface shear strain, , with the assumption of the
pure torsion of the helicoidal-spring specimen:
 =
PD
a b2 2
(1)
 =
2
2
a
Dπ
Δ (2)
where P is the average load, D is the coil diameter
(12 mm) and  is the displacement of the mean
coil-pitch spacing. In this study, the torsion was the
dominant component of deformation, because D was
quite higher than d (D/d > 12)17 and the value of  was
between 2 mm and 4 mm.18 Since the stress and strain in
the helicoidal spring had essentially shear components,
the former could be transformed into the equivalent
tensile quantities with the von Mises equations for the
tensile stress  = 3 and the tensile strain  = / 3.
3 RESULTS AND DISCUSSION
Figure 3 demonstrates the creep curves obtained at
923 K and at various stresses: (34.85, 27.75 and 19.35) MPa.
J. ZHE et al.: SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW STRESSES ...
844 Materiali in tehnologije / Materials and technology 51 (2017) 5, 843–847
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Table 1: Chemical components of P91
Element Cr Mo Si V C Nb Cu N P Ni Ti Al S
Content (%) 9.50 0.91 0.60 0.20 0.10 0.10 0.07 0.04 0.02 0.02 0.01 0.01 0.01
Figure 2: Experimental apparatus including electric furnace, high-
precision optical micrometer, load, weight and information-receiving
and processing systemFigure 1: Helicoidal-spring specimen for the stress-change testing
The creep rate which corresponded to the slope of the
curves decreased as the duration increased.
Figure 4 presents the strain rate on a logarithmic plot
versus the strain at 923 K and various stresses. The creep
deformation of the P91 steels consisted of the primary
creep stages where the creep rate decreased along with
the strain and no apparent secondary (steady-state) creep
stage was observed where the creep rate would not
change along with the strain.
The constitutive creep equations expressing the pri-
mary stages should be utilized in the experimental
creep-curve analysis. The following constitutive creep
equations were utilized, being widely accepted as the
basic creep equations.19,20
Power law:
 = 0 + at
b (3)
Exponential law:
 = 0 + a[1-exp(-bt)] (4)
Logarithmic law:
 = 0 + a ln(1+bt) (5)
Blackburn’s law:
 = 0 + ac[1 - exp(-bt)]+c[1 - exp(-dt)] (6)
where  is the strain, 0 is the instantaneous strain on
dead weight, t is the elapsed time and a, b, c and d are
the parameters characterizing the primary creep region.
The a, b, c and d values were called the "scaling fac-
tors".21
Figure 5 presents representative results of a re-
gression analysis, for the creep at 923 K and 27.75 MPa.
The exponential-law and power-law equations did not
reproduce the experimental data. Compared to the loga-
rithmic-law equation, the Blackburn equation provided a
better representation of the experimental creep curve.
Therefore, the short-term creep for the P91 heat-resistant
steels at low stresses could be described with the
Blackburn equation.
In order for the creep mechanism of the P91 steels at
low stresses to be studied, the instantaneous creep
behavior was investigated. The specimens were loaded
for 4.3 × 105 s until the strain reached 6 × 10–3. In this
case, the dislocations could be expected to move. The
stress increased or decreased at various levels that were
consequently applied and the instantaneous strain was
measured.
J. ZHE et al.: SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW STRESSES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 843–847 845
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Strain/time creep curves for P91 at 923 K and various
stresses
Figure 4: Strain rate/strain creep curves at 923 K and various stresses
Figure 6: Example of instantaneous elongation and contraction upon
low stress changes at a high temperature
Figure 5: Comparison of experimental data and predicted curves of
strain versus time at 923 K and 27.75 MPa
2 EXPERIMENTAL PART
2.1 Materials
The P91 heat-resistant steel was utilized in this expe-
riment. The chemical components are presented in Table 1.
2.2 Specimen processing
The steel was processed into helicoidal-spring speci-
mens using the method of high-speed wire-electrode
cutting machining; the outer and inner diameters of the
specimens were 12 mm and 8 mm, respectively. The
cross-section of the helicoidal-spring specimens was a
rectangle. The side length of the rectangle was 2 mm, as
presented in Figure 1.
The helicoidal-spring specimens were annealed at
1313 K for 60 min and consequently tempered at 1033 K
for 60 min.
2.3 Test methods
Figure 2 presents the experimental apparatus, which
includes an electric furnace with three adjacent heaters, a
high-precision optical micrometer, a load, a weight and a
system for receiving and processing information. By
measuring the helicoidal-spring-specimen pitch change,
the creep curves could be obtained. The test-apparatus
details were previously reported.14
During the testing, the number of weights was in-
creased or decreased for the instantaneous stress change
to be observed. The weight was suspended under a load
with a flammable thread. In order for a load to be
increased, the load was hooked quickly. For a load to be
decreased, the flammable thread was burned.
The test was performed at a temperature of 923 K
and a slight oscillation. The specimen was loaded sub-
sequently for five days when the strain rate was down to
10–9–10–8 s–1. Also, the stress consequently increased or
decreased at various applied levels and the instantaneous
strain was measured. The following equations15,16 were
utilized for calculating the mean surface shear stress, ,
and the surface shear strain, , with the assumption of the
pure torsion of the helicoidal-spring specimen:
 =
PD
a b2 2
(1)
 =
2
2
a
Dπ
Δ (2)
where P is the average load, D is the coil diameter
(12 mm) and  is the displacement of the mean
coil-pitch spacing. In this study, the torsion was the
dominant component of deformation, because D was
quite higher than d (D/d > 12)17 and the value of  was
between 2 mm and 4 mm.18 Since the stress and strain in
the helicoidal spring had essentially shear components,
the former could be transformed into the equivalent
tensile quantities with the von Mises equations for the
tensile stress  = 3 and the tensile strain  = / 3.
3 RESULTS AND DISCUSSION
Figure 3 demonstrates the creep curves obtained at
923 K and at various stresses: (34.85, 27.75 and 19.35) MPa.
J. ZHE et al.: SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW STRESSES ...
844 Materiali in tehnologije / Materials and technology 51 (2017) 5, 843–847
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Table 1: Chemical components of P91
Element Cr Mo Si V C Nb Cu N P Ni Ti Al S
Content (%) 9.50 0.91 0.60 0.20 0.10 0.10 0.07 0.04 0.02 0.02 0.01 0.01 0.01
Figure 2: Experimental apparatus including electric furnace, high-
precision optical micrometer, load, weight and information-receiving
and processing systemFigure 1: Helicoidal-spring specimen for the stress-change testing
The creep rate which corresponded to the slope of the
curves decreased as the duration increased.
Figure 4 presents the strain rate on a logarithmic plot
versus the strain at 923 K and various stresses. The creep
deformation of the P91 steels consisted of the primary
creep stages where the creep rate decreased along with
the strain and no apparent secondary (steady-state) creep
stage was observed where the creep rate would not
change along with the strain.
The constitutive creep equations expressing the pri-
mary stages should be utilized in the experimental
creep-curve analysis. The following constitutive creep
equations were utilized, being widely accepted as the
basic creep equations.19,20
Power law:
 = 0 + at
b (3)
Exponential law:
 = 0 + a[1-exp(-bt)] (4)
Logarithmic law:
 = 0 + a ln(1+bt) (5)
Blackburn’s law:
 = 0 + ac[1 - exp(-bt)]+c[1 - exp(-dt)] (6)
where  is the strain, 0 is the instantaneous strain on
dead weight, t is the elapsed time and a, b, c and d are
the parameters characterizing the primary creep region.
The a, b, c and d values were called the "scaling fac-
tors".21
Figure 5 presents representative results of a re-
gression analysis, for the creep at 923 K and 27.75 MPa.
The exponential-law and power-law equations did not
reproduce the experimental data. Compared to the loga-
rithmic-law equation, the Blackburn equation provided a
better representation of the experimental creep curve.
Therefore, the short-term creep for the P91 heat-resistant
steels at low stresses could be described with the
Blackburn equation.
In order for the creep mechanism of the P91 steels at
low stresses to be studied, the instantaneous creep
behavior was investigated. The specimens were loaded
for 4.3 × 105 s until the strain reached 6 × 10–3. In this
case, the dislocations could be expected to move. The
stress increased or decreased at various levels that were
consequently applied and the instantaneous strain was
measured.
J. ZHE et al.: SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW STRESSES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 843–847 845
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Strain/time creep curves for P91 at 923 K and various
stresses
Figure 4: Strain rate/strain creep curves at 923 K and various stresses
Figure 6: Example of instantaneous elongation and contraction upon
low stress changes at a high temperature
Figure 5: Comparison of experimental data and predicted curves of
strain versus time at 923 K and 27.75 MPa
Figure 7 presents instantaneous-elongation and con-
traction examples upon low changes in the stress of the
P91 steels. In the figure, the – was the instantaneous
strain under a stress decrease and the + was the in-
stantaneous strain under a stress increase.
The relationship between the stress increment ||
and the instantaneous strain |E| is presented in Fig-
ure 7. The creep demonstrates a viscous behavior be-
cause the absolute values of the instantaneous strain for
the load increase were equal to the values for the load
decrease.
Two types of creep at low stress exist. One is creep
controlled by the diffusion including the lattice diffusion
creep (Nabarro-Herring type)22 at a high-T and grain
boundary diffusion creep (Coble type)22 at an inter-
mediate-T. The other is creep associated with dislocation
movement including free flight motion (climb con-
trolled) and viscous motion (glide controlled)23. When
creep is controlled by the diffusion, the creep should be a
non-viscous behavior. In this case, the absolute values of
the instantaneous strain for a small-load increase should
be apparently larger than the values for a small-load
decrease. If creep is controlled by free flight motion of
dislocation, plastic strain can occur instantaneously
when the stress is increased by a small amount. There-
fore, creep shows non-viscous behavior. When creep is
controlled by the viscous glide of dislocations, instant-
aneous plastic strain does not occur even if the applied
stress is increase suddenly. Thus, creep shows viscous
behavior. In this study, the creep demonstrates a viscous
behavior, because non instantaneous strain is observed
during the stress increase. It means creep may be con-
trolled by the viscous glide of dislocations.
Figure 8 gives an example of the anelastic backflow,
observed for a high reduction in the stress during the
transient-creep stage. Specifically, the helicoidal spring
specimen was loaded at 27.75 MPa and 923 K for 120 h
and consequently unloaded all the stress for 120 h. Two
different parts of the creep strain existed: the anelastic
strain was reversible, which is presented with solid
double arrows, whereas the plastic strain was irreversible
and it is presented with dashed double arrows. Therefore,
the instantaneous strain that occurred under a sudden
stress change included both elastic and anelastic compo-
nents.
4 CONCLUSIONS
The short-term creep behavior in the P91 heat-resis-
tant steels was investigated using instantaneous-stress-
change tests. The results were as follows:
1) The creep deformation consisted of the primary creep
stages, but no secondary creep stage was observed.
2) The Blackburn equation could be utilized for the
creep-curve description because it provided an im-
proved representation of the experimental creep
curve.
3) The viscous glide of dislocations might have been the
dominant creep mechanism because the creep dis-
played a viscous behavior.
4) The short-term creep at a low stress displayed an
anelastic behavior.
Acknowledgment
This project was supported by the National Natural
Science Foundation of China (Grant No. 51605330).
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MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
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doi:10.2320/matertrans1960.16.671
12 K. Abe, H. Yoshinaga, S. Morozumi, A Method of Discerning Fric-
tional Stress and Internal Stress by the Stress Relaxation Test,
Journal of the Japan Institute of Metals and Materials, 18 (1997) 6,
479–487, doi:10.2320/matertrans1960.18.479
13 H. Oikawa, H. Sugawara, Instantaneous plastic strain associated with
stress increments during the steady state creep of Al and Al - 5.5 At.
Pct. Mg alloy, Scripta Metall., 12 (1978) 1, 85–89, doi:10.1016/
0036-9748(78)90234-x
14 A. Magnin, Measurement of Very Low Creep Strains, Journal of
Testing and Evaluation, 37 (2009) 1, 53–58, doi:10.1122/1.5550242
15 J. J. Shen, K. Ikeda, S. Hata , H. Nakashima, Instantaneous creep in
face-centered cubic metals at ultra-low strain rates by a high-
resolution strain measurement, Journal of Wuhan University of
Technology (Materials Science Edition), 28 (2013) 6, 1096–1100,
doi:10.1007/s11595-013-0826-y
16 S. Timoshenko, D. H. Yong, Elements of Strength of Materials, 5th
ed., New York, Van Nostrand, 1968, 77–80
17 A. M. Wahi, Mechanical Springs, 2nd ed., New York, McGraw Hill,
1963, 229–230
18 M. R. Mitchel, R. E. Link, P. Marecek, Measurement of Very Low
Creep Strain, A Review, J. Test. Eval., 37 (2009) 1, 53–58, doi:10. 15
20/jte101475
19 P. S. Zhang, J. J. Shen, H. Zhang, Short-Term Creep Behavior in P91
Heat-Resistant Steel at Low Stress, J. Material Science Forum, 850
(2016) 1, 922–926, doi:10.4028/www. scientific. net/ MSF.850.922
20 S. Holdsworth, Developments in the assessment of creep strain and
ductility data, Materials at High Temperatures, 21 (2004) 1, 25–32,
doi:10.3184/096034004782750041
21 R. P. Reed, N. J. Simon , R. P. Walsh, Creep of copper: 4–300 K,
Materials Science and Engineering A, 147 (1991) 1, 23–32,
doi:10.1016/0921-5093(91)90801-s
22 W. D. Nix, Effects of Grain Shape on Nabarro-Herring and Coble
Creep Processes, Metals Forum, 4 (1981) 1, 38–43
23 H. Hiroyuki, N. Satoshi, K. Junichi, K. Akihiro, Creep deformation
characterization of heat resistant steel by stress change test,
International Journal of Pressure Vessels and Piping, 24 (2009) 9,
556–562, doi:10.1016/j.ijpvp.2008.06.003
24 K. Maruyama, S. Karashima, Theorethical Consideration of
Measurement of Wor-Hardening and Recovery Rates during high
temperature Creep, Trans. Japan Inst. Metals, 16 (1975) 11,
671–678, doi:10.2320/matertrans1960.16.671
J. ZHE et al.: SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW STRESSES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 843–847 847
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7 presents instantaneous-elongation and con-
traction examples upon low changes in the stress of the
P91 steels. In the figure, the – was the instantaneous
strain under a stress decrease and the + was the in-
stantaneous strain under a stress increase.
The relationship between the stress increment ||
and the instantaneous strain |E| is presented in Fig-
ure 7. The creep demonstrates a viscous behavior be-
cause the absolute values of the instantaneous strain for
the load increase were equal to the values for the load
decrease.
Two types of creep at low stress exist. One is creep
controlled by the diffusion including the lattice diffusion
creep (Nabarro-Herring type)22 at a high-T and grain
boundary diffusion creep (Coble type)22 at an inter-
mediate-T. The other is creep associated with dislocation
movement including free flight motion (climb con-
trolled) and viscous motion (glide controlled)23. When
creep is controlled by the diffusion, the creep should be a
non-viscous behavior. In this case, the absolute values of
the instantaneous strain for a small-load increase should
be apparently larger than the values for a small-load
decrease. If creep is controlled by free flight motion of
dislocation, plastic strain can occur instantaneously
when the stress is increased by a small amount. There-
fore, creep shows non-viscous behavior. When creep is
controlled by the viscous glide of dislocations, instant-
aneous plastic strain does not occur even if the applied
stress is increase suddenly. Thus, creep shows viscous
behavior. In this study, the creep demonstrates a viscous
behavior, because non instantaneous strain is observed
during the stress increase. It means creep may be con-
trolled by the viscous glide of dislocations.
Figure 8 gives an example of the anelastic backflow,
observed for a high reduction in the stress during the
transient-creep stage. Specifically, the helicoidal spring
specimen was loaded at 27.75 MPa and 923 K for 120 h
and consequently unloaded all the stress for 120 h. Two
different parts of the creep strain existed: the anelastic
strain was reversible, which is presented with solid
double arrows, whereas the plastic strain was irreversible
and it is presented with dashed double arrows. Therefore,
the instantaneous strain that occurred under a sudden
stress change included both elastic and anelastic compo-
nents.
4 CONCLUSIONS
The short-term creep behavior in the P91 heat-resis-
tant steels was investigated using instantaneous-stress-
change tests. The results were as follows:
1) The creep deformation consisted of the primary creep
stages, but no secondary creep stage was observed.
2) The Blackburn equation could be utilized for the
creep-curve description because it provided an im-
proved representation of the experimental creep
curve.
3) The viscous glide of dislocations might have been the
dominant creep mechanism because the creep dis-
played a viscous behavior.
4) The short-term creep at a low stress displayed an
anelastic behavior.
Acknowledgment
This project was supported by the National Natural
Science Foundation of China (Grant No. 51605330).
5 REFERENCES
1 K. Kimura, Y. Takahashi, ASME 2012 Pressure Vessels & Piping
Division Conference, American Society of Mechanical Engineers, 18
(2012) 2, 177, doi:10.1016/0308-0161(93)90110-f
2 K. Maruyama, J. S. Lee, Creep & Fracture in High Temperature
Components: Design & Life Assessment Issues, DEStech
3 H. G. Armaki, R. Chen, K. Maruyama, M. Igaras, Premature creep
failure in strength enhanced high Cr ferritic steels caused by static
recovery of tempered martensite lath structures, Materials Science
and Engineering A, 527 (2010) 24–25, 6581–6588, doi:10.1016/ j.
Msea. 2010. 07. 037
J. ZHE et al.: SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW STRESSES ...
846 Materiali in tehnologije / Materials and technology 51 (2017) 5, 843–847
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7: Instantaneous strain at a high temperature (T = 923 K) ver-
sus the stress change during P91 steady-state creep
Figure 8: Elongation versus time during a high stress reduction with-
out an elastic strain
4 O. Prat, J. Garcia, D. Rojas, G. Sauthoff, G. Inden, The role of Laves
phase on microstructure evolution and creep strength of novel 9%Cr
heat resistant steels, Intermetallics, 32 (2013) 32, 362–372,
doi:10.1016/j.internet.2012.08.016
5 L. Kyu-Ho, H. Sung-Min, S. Jae-Hyeok, S. Jin-Yoo, H. Joo-Youl, J.
Woo-Sang, Effect of Nb addition on Z-phase formation and creep
strength in high-Cr martensitic heat-resistant steels, Materials
Characterization, 102 (2015) 2, 79–84, doi:10.1016/j.Matchar. 2014.
12. 028
6 X. Dong, Q. Z. Gao, C. Li, 9Cr-1.7W-0.4Mo-Co Ferritic Heat-Resis-
tant Steel Isothermal Aging Microstructural Evolution, Publications,
Lancaster, 85 (2008) 1–2, 372–379, doi:10.1016/j.ijpvp. 2007. 06.
006
7 Transactions of Materials and Heat Treatment, 36 (2015) 3, 81–86,
doi:10. 1016/j.Msea.2013.09.033
8 K. Kimura, H. Kushima, K. Sawada, Long-term creep deformation
property of modified 9Cr–1Mo steel, Materials Science and Engi-
neering A, 511 (2009) 18, 58–63, doi:10.1016/j.msea.2008.04.095
9 K. Kimura, Y. Toda, H. Kushima, K. Sawada, Creep strength of high
chromium steel with ferrite matrix, International Journal of Pressure
Vessels and Piping, 87 (2010) 6, 282–288, doi:10.1016/j.ij2010.03.
016
10 E. M. Haney, F. Dalle, M. Sauzay, L. Vincent, I. Tournié, L. Allais,
B. Fournier, Macroscopic results of long-term creep on a modified
9Cr–1Mo steel (T91), Materials Science and Engineering A, 510
(2009) 18, 99–103, doi: 10.1016/ j.msea. 2008. 04.099
11 K. Maruyama, S. Karashima, Theoretical Consideration of Measure-
ment of Work-Hardening and Recovery Rates during High Tem-
perature Creep, Trans. Japan Inst. Metals, 16 (1975) 11, 671–678,
doi:10.2320/matertrans1960.16.671
12 K. Abe, H. Yoshinaga, S. Morozumi, A Method of Discerning Fric-
tional Stress and Internal Stress by the Stress Relaxation Test,
Journal of the Japan Institute of Metals and Materials, 18 (1997) 6,
479–487, doi:10.2320/matertrans1960.18.479
13 H. Oikawa, H. Sugawara, Instantaneous plastic strain associated with
stress increments during the steady state creep of Al and Al - 5.5 At.
Pct. Mg alloy, Scripta Metall., 12 (1978) 1, 85–89, doi:10.1016/
0036-9748(78)90234-x
14 A. Magnin, Measurement of Very Low Creep Strains, Journal of
Testing and Evaluation, 37 (2009) 1, 53–58, doi:10.1122/1.5550242
15 J. J. Shen, K. Ikeda, S. Hata , H. Nakashima, Instantaneous creep in
face-centered cubic metals at ultra-low strain rates by a high-
resolution strain measurement, Journal of Wuhan University of
Technology (Materials Science Edition), 28 (2013) 6, 1096–1100,
doi:10.1007/s11595-013-0826-y
16 S. Timoshenko, D. H. Yong, Elements of Strength of Materials, 5th
ed., New York, Van Nostrand, 1968, 77–80
17 A. M. Wahi, Mechanical Springs, 2nd ed., New York, McGraw Hill,
1963, 229–230
18 M. R. Mitchel, R. E. Link, P. Marecek, Measurement of Very Low
Creep Strain, A Review, J. Test. Eval., 37 (2009) 1, 53–58, doi:10. 15
20/jte101475
19 P. S. Zhang, J. J. Shen, H. Zhang, Short-Term Creep Behavior in P91
Heat-Resistant Steel at Low Stress, J. Material Science Forum, 850
(2016) 1, 922–926, doi:10.4028/www. scientific. net/ MSF.850.922
20 S. Holdsworth, Developments in the assessment of creep strain and
ductility data, Materials at High Temperatures, 21 (2004) 1, 25–32,
doi:10.3184/096034004782750041
21 R. P. Reed, N. J. Simon , R. P. Walsh, Creep of copper: 4–300 K,
Materials Science and Engineering A, 147 (1991) 1, 23–32,
doi:10.1016/0921-5093(91)90801-s
22 W. D. Nix, Effects of Grain Shape on Nabarro-Herring and Coble
Creep Processes, Metals Forum, 4 (1981) 1, 38–43
23 H. Hiroyuki, N. Satoshi, K. Junichi, K. Akihiro, Creep deformation
characterization of heat resistant steel by stress change test,
International Journal of Pressure Vessels and Piping, 24 (2009) 9,
556–562, doi:10.1016/j.ijpvp.2008.06.003
24 K. Maruyama, S. Karashima, Theorethical Consideration of
Measurement of Wor-Hardening and Recovery Rates during high
temperature Creep, Trans. Japan Inst. Metals, 16 (1975) 11,
671–678, doi:10.2320/matertrans1960.16.671
J. ZHE et al.: SHORT-TERM CREEP OF P91 HEAT-RESISTANT STEELS AT LOW STRESSES ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 843–847 847
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
J. PALÁN et al.: EFFECT OF SEVERE PLASTIC AND HEAVY COLD DEFORMATION ...
849–853
EFFECT OF SEVERE PLASTIC AND HEAVY COLD
DEFORMATION ON THE STRUCTURAL AND MECHANICAL
PROPERTIES OF COMMERCIALLY PURE TITANIUM
U^INEK PLASTI^NOSTI IN DEFORMACIJE PRI
PODHLAJEVANJU NA STRUKTURNE IN MEHANSKE LASTNOSTI
^ISTEGA KOMERCIALNEGA TITANA
Jan Palán1, Pavol [utta2, Tomá{ Kubina1, Mária Dománková3
1COMTES FHT, Prumyslova 995, 334 41 Dobrany, Czech Republic
2New Technologies – Research Centre Westbohemian Region, University of West Bohemia, 306 14 Plzen, Czech Republic
3Institute of Materials Science, Faculty of Materials Science and Technology Trnava, STU Bratislava, Bottova 25, 917 24 Trnava, Slovakia
jpalan@comtesfht.cz
Prejem rokopisa – received: 2016-10-25; sprejem za objavo – accepted for publication: 2017-01-23
doi:10.17222/mit.2016.307
SPD (severe plastic deformation) processing of materials provides a great potential associated with the enhancement of their
properties by refining the initial grain structure. The present experiments involved mechanical working of commercial-purity
titanium (Ti Grade 2) with the CONFORM SPD technique, which is one of the SPD methods, and with rotary swaging. The
objective was to process the material at as low temperatures as possible in order to avoid softening processes and, therefore, to
achieve the maximum strengthening through a microstructure refinement. Three passes through a CONFORM SPD machine
were completed and the resulting ultimate strength was 673 MPa. The average grain size was 330 nm. The greatest improvement
of the mechanical properties was achieved in the first pass. In the subsequent passes, the contributions were minor. The pro-
cessing in the CONFORM SPD machine did not impair the ductility of the material. Subsequently, the wires were rotary
swaged. The ultimate strength achieved was 1070 MPa. The response of the properties to this forming method was markedly
different. The reason is that rotary swaging does not belong to SPD techniques. It causes rapid work hardening and reduces the
ductility of the material. The workpiece was subsequently investigated with the aid of several techniques. Light and trans-
mission electron microscopy and X-ray diffraction were employed for evaluating the grain size, distribution and orientation.
Keywords: equal-channel angular pressing, CONFORM SPD technique, rotary swaging, titanium, extrusion
Velika plasti~na deformacija (angl. SPD) pri obdelavi materialov omogo~a veliko mo`nosti, povezane z izbolj{anjem njihovih
lastnosti z izbolj{avo izvorne strukture zrn. Pri~ujo~i preizkus je vklju~eval mehansko obdelavo komercialno dostopnega ~istega
titana (Ti Grade 2) s CONFORM SPD-tehniko, ki je ena od SPD-metod, in s kovanjem. Namen je bil, da bi se material obdelalo
pri ~im ni`jih temperaturah, kot je mogo~e, da bi se izognili procesom meh~anja in, da bi dosegli maksimalno krepitev s pre-
~i{~enjem strukture. Narejeni so bili trije nizi v CONFORM SPD-stroju in kon~na mo~ je dosegla 673 MPa. Povpre~na velikost
zrn je bila 330 nm. Najve~je izbolj{ave mehanskih lastnosti so bile dose`ene v prvem nizu. V slede~ih nizih so bile le-te manj{e.
Obdelava v CONFORM SPD-stroju ni {kodovala duktilnosti materiala. Posledi~no so bile `ice zvite. Kon~na mo~ je bila 1070
MPa. Odziv lastnosti na to metodo oblikovanja je bil izrazito druga~en. Razlog je v tem, da rotacijsko zvijanje ne sodi v tehnike
SPD. Rotacijsko zvijanje povzro~i hitro utrjevanje in zmanj{uje duktilnost materiala. Vzorec je bil nato raziskan s pomo~jo ve~
tehnik. Za ovrednotenje velikosti, razporeditve in orientacije zrn, sta bili uporabljeni transmisijska elektronska mikroskopija in
rentgenska difrakcija.
Klju~ne besede: enakomerno kotno stiskanje, CONFORM SPD-tehnika, rotacijsko zvijanje, titan, iztiskanje
1 INTRODUCTION
Severe plastic deformation (SPD) is a generic term
for a group of methods which cause grain refinement in a
material and produce equiaxed grains.1 The occurrence
of an ultrafine structure or nanostructure is conditional
on high hydrostatic pressure (P >1 GPa) and large shear
deformation applied at relatively low temperatures. The
temperatures of SPD processes should meet the
condition of T(SPD) < 0.4 T(melting).1 Another aspect is
the strain magnitude, defined as e(true strain) > 6–8.
When the above conditions are met, the forming process
leads to a high density of lattice defects, predominantly
dislocations, and to the formation of subgrains, which
reduces the stored energy.2 The relationship between the
mechanical properties and the grain size is described
with the Hall-Petch equation3,4. In the present study,
commercially pure (CP) titanium Grade 2 was worked by
means of the CONFORM SPD (CONSPD) method. This
method uses continuous angular pressing through a die
chamber of a modified design to produce ultrafine
structures and nanostructures of the materials.5–7 The
design modification reflects the principles of the ECAP
method (equal-channel angular pressing), as described in
studies.5–7 Continuous operation is the key advantage of
this forming method. The processing of CP titanium
Grade 2 with the CONFORM SDP machine was des-
cribed in multiple papers.5,7 The main outcomes were the
improved mechanical properties (ultimate tensile
strength (UTS) = 698 MPa, 0.2 % offset yield stress
(0.2 OYS) = 637 MPa), no significant decrease in the
Materiali in tehnologije / Materials and technology 51 (2017) 5, 849–853 849
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:620.1:536.421.48 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)849(2017)
J. PALÁN et al.: EFFECT OF SEVERE PLASTIC AND HEAVY COLD DEFORMATION ...
849–853
EFFECT OF SEVERE PLASTIC AND HEAVY COLD
DEFORMATION ON THE STRUCTURAL AND MECHANICAL
PROPERTIES OF COMMERCIALLY PURE TITANIUM
U^INEK PLASTI^NOSTI IN DEFORMACIJE PRI
PODHLAJEVANJU NA STRUKTURNE IN MEHANSKE LASTNOSTI
^ISTEGA KOMERCIALNEGA TITANA
Jan Palán1, Pavol [utta2, Tomá{ Kubina1, Mária Dománková3
1COMTES FHT, Prumyslova 995, 334 41 Dobrany, Czech Republic
2New Technologies – Research Centre Westbohemian Region, University of West Bohemia, 306 14 Plzen, Czech Republic
3Institute of Materials Science, Faculty of Materials Science and Technology Trnava, STU Bratislava, Bottova 25, 917 24 Trnava, Slovakia
jpalan@comtesfht.cz
Prejem rokopisa – received: 2016-10-25; sprejem za objavo – accepted for publication: 2017-01-23
doi:10.17222/mit.2016.307
SPD (severe plastic deformation) processing of materials provides a great potential associated with the enhancement of their
properties by refining the initial grain structure. The present experiments involved mechanical working of commercial-purity
titanium (Ti Grade 2) with the CONFORM SPD technique, which is one of the SPD methods, and with rotary swaging. The
objective was to process the material at as low temperatures as possible in order to avoid softening processes and, therefore, to
achieve the maximum strengthening through a microstructure refinement. Three passes through a CONFORM SPD machine
were completed and the resulting ultimate strength was 673 MPa. The average grain size was 330 nm. The greatest improvement
of the mechanical properties was achieved in the first pass. In the subsequent passes, the contributions were minor. The pro-
cessing in the CONFORM SPD machine did not impair the ductility of the material. Subsequently, the wires were rotary
swaged. The ultimate strength achieved was 1070 MPa. The response of the properties to this forming method was markedly
different. The reason is that rotary swaging does not belong to SPD techniques. It causes rapid work hardening and reduces the
ductility of the material. The workpiece was subsequently investigated with the aid of several techniques. Light and trans-
mission electron microscopy and X-ray diffraction were employed for evaluating the grain size, distribution and orientation.
Keywords: equal-channel angular pressing, CONFORM SPD technique, rotary swaging, titanium, extrusion
Velika plasti~na deformacija (angl. SPD) pri obdelavi materialov omogo~a veliko mo`nosti, povezane z izbolj{anjem njihovih
lastnosti z izbolj{avo izvorne strukture zrn. Pri~ujo~i preizkus je vklju~eval mehansko obdelavo komercialno dostopnega ~istega
titana (Ti Grade 2) s CONFORM SPD-tehniko, ki je ena od SPD-metod, in s kovanjem. Namen je bil, da bi se material obdelalo
pri ~im ni`jih temperaturah, kot je mogo~e, da bi se izognili procesom meh~anja in, da bi dosegli maksimalno krepitev s pre-
~i{~enjem strukture. Narejeni so bili trije nizi v CONFORM SPD-stroju in kon~na mo~ je dosegla 673 MPa. Povpre~na velikost
zrn je bila 330 nm. Najve~je izbolj{ave mehanskih lastnosti so bile dose`ene v prvem nizu. V slede~ih nizih so bile le-te manj{e.
Obdelava v CONFORM SPD-stroju ni {kodovala duktilnosti materiala. Posledi~no so bile `ice zvite. Kon~na mo~ je bila 1070
MPa. Odziv lastnosti na to metodo oblikovanja je bil izrazito druga~en. Razlog je v tem, da rotacijsko zvijanje ne sodi v tehnike
SPD. Rotacijsko zvijanje povzro~i hitro utrjevanje in zmanj{uje duktilnost materiala. Vzorec je bil nato raziskan s pomo~jo ve~
tehnik. Za ovrednotenje velikosti, razporeditve in orientacije zrn, sta bili uporabljeni transmisijska elektronska mikroskopija in
rentgenska difrakcija.
Klju~ne besede: enakomerno kotno stiskanje, CONFORM SPD-tehnika, rotacijsko zvijanje, titan, iztiskanje
1 INTRODUCTION
Severe plastic deformation (SPD) is a generic term
for a group of methods which cause grain refinement in a
material and produce equiaxed grains.1 The occurrence
of an ultrafine structure or nanostructure is conditional
on high hydrostatic pressure (P >1 GPa) and large shear
deformation applied at relatively low temperatures. The
temperatures of SPD processes should meet the
condition of T(SPD) < 0.4 T(melting).1 Another aspect is
the strain magnitude, defined as e(true strain) > 6–8.
When the above conditions are met, the forming process
leads to a high density of lattice defects, predominantly
dislocations, and to the formation of subgrains, which
reduces the stored energy.2 The relationship between the
mechanical properties and the grain size is described
with the Hall-Petch equation3,4. In the present study,
commercially pure (CP) titanium Grade 2 was worked by
means of the CONFORM SPD (CONSPD) method. This
method uses continuous angular pressing through a die
chamber of a modified design to produce ultrafine
structures and nanostructures of the materials.5–7 The
design modification reflects the principles of the ECAP
method (equal-channel angular pressing), as described in
studies.5–7 Continuous operation is the key advantage of
this forming method. The processing of CP titanium
Grade 2 with the CONFORM SDP machine was des-
cribed in multiple papers.5,7 The main outcomes were the
improved mechanical properties (ultimate tensile
strength (UTS) = 698 MPa, 0.2 % offset yield stress
(0.2 OYS) = 637 MPa), no significant decrease in the
Materiali in tehnologije / Materials and technology 51 (2017) 5, 849–853 849
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 67.017:620.1:536.421.48 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)849(2017)
elongation, and an equiaxed ultrafine-grained (UFG)
microstructure.7
The present study proposes a process route which
involves CONFORM SPD forming and rotary swaging
(RS) in order to enhance the mechanical properties of a
workpiece.8–10 The effects of severe plastic deformation
and work hardening are thus combined. The proposed
technology allows the mechanical properties of pure
titanium to be improved dramatically. The ultrafine-
grained CP Ti Grade 2 has a great potential because its
mechanical properties are comparable to the Ti6Al4V
alloy, which is mostly used in medicine. Recently, it was
suggested that Al and V might be toxic for the human
body.11,12
2 EXPERIMENTAL PART
The material under investigation was CP titanium
Grade 2 with the chemical composition given in Table 1.
The composition was measured by means of the Bruker
Q4 Tasman optical emission spectrometer and the Bruker
G8 Galileo gas analyser. The diameter of Ti rods was
10 mm.
Table 1: Chemical composition of feedstock in weight percent
Fe O C H N Ti
0.046 0.12 0.023 0.0026 0.0076 99.822
The processing was carried out in a CONFORM SPD
machine, type 315i with a modified die chamber (Figure
1), and in a HMP R4-4 rotary-swaging machine. During
the CONFORM SPD process, the temperature was
220 °C, the wheel speed was 0.5 min–1, and the angle of
the die chamber was 90°. Three passes through the
CONFORM SPD machine were completed. The pro-
duct’s cross-section was identical to that of the feed-
stock. Rotary swaging, the subsequent process, was
carried out at ambient temperature. In this operation, the
cross-section area was reduced by 20 % in each pass.
The total area reduction was 90 %.
For the purpose of observation with a transmission
electron microscope (TEM), thin foils were prepared
with the final electrolytic thinning in a Tenupol 5 device,
using a solution of 300 mL CH3OH + 175 mL 2-butanol
+ 30 mL HClO4 at –10 °C and a voltage of 40 V. The
TEM analysis was performed with a JEOL 200CX
instrument with an acceleration voltage of 200 kV. Selec-
tive electron diffraction was used for the determination
of the phases. The grain size was measured using the
linear intercept method.
The preferred orientation of crystallites (texture) was
analysed with an automatic powder diffractometer
X’Pert-Pro equipped with an ultra-fast semiconductor
position-sensitive detector Pixcel. Cu-K1 radiation
( = 0.154056 nm) was used. The texture was characte-
rized from the radial X-ray diffraction patterns using the
Harris (1952) texture index, Equation (1):
T
n I /R
I /R
i
i i
j j
j
n=
⋅
=
∑
1
(1)
where n is the number of the reflections investigated, Ii
is the observed intensity and Ri is the corresponding
intensity for the sample with randomly oriented
crystallites. The Ri values were obtained from the ICDD
PDF standard diffraction data file (reference code
00-005-0682). The texture was analysed in the direction
of the material flow, both after CONFORM SPD
(longitudinal direction), and after CONFORM SPD +
rotary swaging.
3 MICROSTRUCTURE
The mean grain size in the initial condition of the
feedstock was 5390 nm. After the first pass, the mean
grain size was daverage ~ 320±35 nm in the transverse
direction. After the first pass, the microstructure was
equiaxed with a non-uniform dislocation density (Figure
2a). Some locations exhibited a preferential orientation
(Figure 2b). After the second pass (Figure 2c), the mean
grain size was daverage ~ 250±25 nm in transverse
direction. The dislocation density was still non-uniform
after the second pass. After the third pass, the mean grain
size was daverage ~ 330±30 nm (Figure 3d): the grain was
larger than after the second pass. This increase can be
attributed, in part, to the non-uniform deformation and to
the high surface activity of the UFG microstructure.
High surface activity and dislocation density (strain mag-
nitude) lower the temperatures of softening processes. A
certain grain growth can be expected to occur during the
forming due to deformation heat and the heat retained in
the die chamber (220 °C). The grain growth upon multi-
ple passes through the CONFORM SPD machine was
already reported in 9.
Figure 3a is a micrograph of the structure in the
transverse direction after three passes through the CON-
FORM SPD machine and after a 35 % area reduction
J. PALÁN et al.: EFFECT OF SEVERE PLASTIC AND HEAVY COLD DEFORMATION ...
850 Materiali in tehnologije / Materials and technology 51 (2017) 5, 849–853
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Schematic illustration of the CONFORM SPD process
with rotary swaging (RS). The mean grain size in the
transverse direction is daverage ~ 300±150 nm. The large
standard deviation is due to a large variation in the grain
size. Figure 3b is a micrograph taken in the longitudinal
direction after CONFORM SPD and RS. The grains are
elongated and aligned in the direction of forming due to
the nature of the deformation introduced by rotary
swaging. Unlike in CONFORM SPD, the grains become
elongated instead of being refined or converted into new
polyhedral grains. After the additional reduction due to
rotary swaging (3 passes through CONFORM SPD + the
total area reduction of 50 % by RS), it became im-
possible to determine the grain size due to the high
dislocation density, as seen in Figure 3c. The summary
of the average grain sizes is given in Table 2.
Table 2: Average grain size in the transverse and longitudinal sections
obtained from TEM images
Condition
Grain size (nm)
Transverse Longitudinal
as received 5390±20
CON SPD – 1 pass 320±35 340±30
CONSPD – 2 passes 250±25 310±30
CONSPD – 3 passes 330±30 420±30
CON SPD – 3 passes +
RS (35 %) 300±150 220±50
CONSPD – 3 passes +
RS (50 %)
Not possible to evaluate due to high
dislocation density
J. PALÁN et al.: EFFECT OF SEVERE PLASTIC AND HEAVY COLD DEFORMATION ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 849–853 851
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: TEM images: a) CON SPD – 3 passes + RS (area reduction
by 35 %) in the transverse direction, b) CON SPD – 3 passes + RS
(35 %) in the longitudinal direction, c) CON SPD – 3 passes + RS
(50 %) in the transverse direction
Figure 2: TEM images: a) CON SPD – 1 pass (transverse), b) CON
SPD – 1 pass (longitudinal), c) CON SPD – 2 passes (transverse),
d) CON SPD – 3 passes (transverse)
elongation, and an equiaxed ultrafine-grained (UFG)
microstructure.7
The present study proposes a process route which
involves CONFORM SPD forming and rotary swaging
(RS) in order to enhance the mechanical properties of a
workpiece.8–10 The effects of severe plastic deformation
and work hardening are thus combined. The proposed
technology allows the mechanical properties of pure
titanium to be improved dramatically. The ultrafine-
grained CP Ti Grade 2 has a great potential because its
mechanical properties are comparable to the Ti6Al4V
alloy, which is mostly used in medicine. Recently, it was
suggested that Al and V might be toxic for the human
body.11,12
2 EXPERIMENTAL PART
The material under investigation was CP titanium
Grade 2 with the chemical composition given in Table 1.
The composition was measured by means of the Bruker
Q4 Tasman optical emission spectrometer and the Bruker
G8 Galileo gas analyser. The diameter of Ti rods was
10 mm.
Table 1: Chemical composition of feedstock in weight percent
Fe O C H N Ti
0.046 0.12 0.023 0.0026 0.0076 99.822
The processing was carried out in a CONFORM SPD
machine, type 315i with a modified die chamber (Figure
1), and in a HMP R4-4 rotary-swaging machine. During
the CONFORM SPD process, the temperature was
220 °C, the wheel speed was 0.5 min–1, and the angle of
the die chamber was 90°. Three passes through the
CONFORM SPD machine were completed. The pro-
duct’s cross-section was identical to that of the feed-
stock. Rotary swaging, the subsequent process, was
carried out at ambient temperature. In this operation, the
cross-section area was reduced by 20 % in each pass.
The total area reduction was 90 %.
For the purpose of observation with a transmission
electron microscope (TEM), thin foils were prepared
with the final electrolytic thinning in a Tenupol 5 device,
using a solution of 300 mL CH3OH + 175 mL 2-butanol
+ 30 mL HClO4 at –10 °C and a voltage of 40 V. The
TEM analysis was performed with a JEOL 200CX
instrument with an acceleration voltage of 200 kV. Selec-
tive electron diffraction was used for the determination
of the phases. The grain size was measured using the
linear intercept method.
The preferred orientation of crystallites (texture) was
analysed with an automatic powder diffractometer
X’Pert-Pro equipped with an ultra-fast semiconductor
position-sensitive detector Pixcel. Cu-K1 radiation
( = 0.154056 nm) was used. The texture was characte-
rized from the radial X-ray diffraction patterns using the
Harris (1952) texture index, Equation (1):
T
n I /R
I /R
i
i i
j j
j
n=
⋅
=
∑
1
(1)
where n is the number of the reflections investigated, Ii
is the observed intensity and Ri is the corresponding
intensity for the sample with randomly oriented
crystallites. The Ri values were obtained from the ICDD
PDF standard diffraction data file (reference code
00-005-0682). The texture was analysed in the direction
of the material flow, both after CONFORM SPD
(longitudinal direction), and after CONFORM SPD +
rotary swaging.
3 MICROSTRUCTURE
The mean grain size in the initial condition of the
feedstock was 5390 nm. After the first pass, the mean
grain size was daverage ~ 320±35 nm in the transverse
direction. After the first pass, the microstructure was
equiaxed with a non-uniform dislocation density (Figure
2a). Some locations exhibited a preferential orientation
(Figure 2b). After the second pass (Figure 2c), the mean
grain size was daverage ~ 250±25 nm in transverse
direction. The dislocation density was still non-uniform
after the second pass. After the third pass, the mean grain
size was daverage ~ 330±30 nm (Figure 3d): the grain was
larger than after the second pass. This increase can be
attributed, in part, to the non-uniform deformation and to
the high surface activity of the UFG microstructure.
High surface activity and dislocation density (strain mag-
nitude) lower the temperatures of softening processes. A
certain grain growth can be expected to occur during the
forming due to deformation heat and the heat retained in
the die chamber (220 °C). The grain growth upon multi-
ple passes through the CONFORM SPD machine was
already reported in 9.
Figure 3a is a micrograph of the structure in the
transverse direction after three passes through the CON-
FORM SPD machine and after a 35 % area reduction
J. PALÁN et al.: EFFECT OF SEVERE PLASTIC AND HEAVY COLD DEFORMATION ...
850 Materiali in tehnologije / Materials and technology 51 (2017) 5, 849–853
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Schematic illustration of the CONFORM SPD process
with rotary swaging (RS). The mean grain size in the
transverse direction is daverage ~ 300±150 nm. The large
standard deviation is due to a large variation in the grain
size. Figure 3b is a micrograph taken in the longitudinal
direction after CONFORM SPD and RS. The grains are
elongated and aligned in the direction of forming due to
the nature of the deformation introduced by rotary
swaging. Unlike in CONFORM SPD, the grains become
elongated instead of being refined or converted into new
polyhedral grains. After the additional reduction due to
rotary swaging (3 passes through CONFORM SPD + the
total area reduction of 50 % by RS), it became im-
possible to determine the grain size due to the high
dislocation density, as seen in Figure 3c. The summary
of the average grain sizes is given in Table 2.
Table 2: Average grain size in the transverse and longitudinal sections
obtained from TEM images
Condition
Grain size (nm)
Transverse Longitudinal
as received 5390±20
CON SPD – 1 pass 320±35 340±30
CONSPD – 2 passes 250±25 310±30
CONSPD – 3 passes 330±30 420±30
CON SPD – 3 passes +
RS (35 %) 300±150 220±50
CONSPD – 3 passes +
RS (50 %)
Not possible to evaluate due to high
dislocation density
J. PALÁN et al.: EFFECT OF SEVERE PLASTIC AND HEAVY COLD DEFORMATION ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 849–853 851
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: TEM images: a) CON SPD – 3 passes + RS (area reduction
by 35 %) in the transverse direction, b) CON SPD – 3 passes + RS
(35 %) in the longitudinal direction, c) CON SPD – 3 passes + RS
(50 %) in the transverse direction
Figure 2: TEM images: a) CON SPD – 1 pass (transverse), b) CON
SPD – 1 pass (longitudinal), c) CON SPD – 2 passes (transverse),
d) CON SPD – 3 passes (transverse)
4 TEXTURE EVALUATION
The preferred orientation of crystallites expressed
with the texture index Ti (Equation 1, Table 3) is, in all
the cases, in the [001] direction perpendicular to the
sample surface analysed. The highest value was found
for CONSPD – 3 passes + RS (75 %) titanium sample,
i.e., the sample that was processed using the CONFORM
SPD device and rotary swaging (basal texture, Figure 5).
An increased value of the texture index can be seen in
the initial state, due to the previous wire-drawing process
(Figure 4). The texture index is lower for the samples
that were processed with the CONFORM SPD device.
Texture indices for the [101] direction are listed for the
sake of comparison. Nevertheless, all their values are
lower than 1. It means that no preferred orientation of
crystallites in the [101] direction is present, as opposed
to the sample surface.
Table 3: Texture evaluation using X-ray diffraction
Plane (h k l)
Texture index
as received CONSPD – 1pass
CONSPD – 2
passes
Ti (002) 3.13 2.49 2.67
Ti (101) 0.66 0.69 0.86
CONSPD – 3
passes
CONSPD – 3
passes + RS
(75 %)
Standard
Ti (002) 2.37 3.48 1
Ti (101) 0.97 0.24 1
4 MECHANICAL PROPERTIES
Table 4 indicates that the largest increase in the
mechanical properties was obtained during the first pass
through CONFORM SPD. The subsequent passes led to
smaller increments. It is important to note that the
increase in the ultimate strength and yield stress upon
CONFORM SPD is not offset by a decrease in the
ductility. The improvement in the mechanical properties
is mainly due to the grain refinement and the increased
dislocation density, which impede the dislocation move-
ment. A further increase in the mechanical properties
was obtained with rotary swaging applied after the three
passes through the CONFORM SPD machine. In compa-
rison with the CONFORM SPD process, rotary swaging
leads to a decrease in the ductility as a result of work
hardening.
Table 4: Mechanical properties in the as-received condition, after
CONFORM SPD processing and after RS
Condition
0.2
OYS UTS A5 RA
MPa MPa % %
as received 370 480 25 52
CONSPD – 1 pass 540 580 23 62
CONSPD – 2 passes 560 600 23 62
CONSPD – 3 passes 570 623 20 64
CONSPD – 3 passes + RS
(50 %) 830 885 13 54
CONSPD – 3 passes + RS
(75 %) 930 1000 12 57
CONSPD – 3 passes + RS
(85%) 950 1070 12 58
RS only (85 %) 780 850 12 50
5 DISCUSSION
The CONFORM SPD processing led to higher ulti-
mate strengths and yield stresses without a reduced
ductility. This is characteristic of the process of the for-
mation of a UFG equiaxed structure shown in Figure 2.
The grain size and distortion are non-uniform due to
non-uniform deformation during the forming process.
The preferred orientation of crystallites in the [001]
direction perpendicular to the surfaces of the investigated
samples is not too significant, except for the sample
which was processed with CONFORM SPD and the
rotary swaging machine where the highest value of the
texture index was observed (the basal texture).
Major increases in the ultimate strength and yield
stress were found after the first pass through the CON-
J. PALÁN et al.: EFFECT OF SEVERE PLASTIC AND HEAVY COLD DEFORMATION ...
852 Materiali in tehnologije / Materials and technology 51 (2017) 5, 849–853
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Diffraction radial profile of the specimen with the strongest
texture: CONSPD – 3 passes + rotary swaging (75 %)
Figure 4: Diffraction radial profile of the as-received specimen
FORM SPD machine, whereas the increments from the
subsequent passes were small. This is closely related to
the grain size, which did not decrease during the sub-
sequent passes. It is possible to say that the additional
deformation causes the subgrains to rotate into high-
angle grain boundaries, typically with an equiaxed shape
of the grains.5,7,9 The subsequent rotary swaging led to a
further increase in the mechanical properties, but the
strengthening mechanism was different this time. Instead
of the formation of equiaxed grains, the grains became
elongated in the direction of forming, with a much
higher dislocation density (Figure 3).10,11 When com-
pared to the CONFORM SPD route, rotary swaging led
to a reduced ductility as a consequence of work harden-
ing.
6 CONCLUSION
This study involved the processing of Ti Grade 2
using CONFORM SPD and rotary swaging with the goal
of improving its mechanical properties. The results and
findings are described below:
• The CONFORM SPD processing substantially
refined the grain from its initial size of 5390 nm to
350 nm. The subsequent rotary swaging did not
provide for a further grain refinement but led to a
grain elongation and to a higher dislocation density.
• The processing led to the basal texture. The most
significant form of texture was found in the specimen
processed with CONFORM SPD and rotary swaging.
• Three passes through the CONFORM SPD machine
led to a strength of 623 MPa and a yield stress of
570 MPa without any notable decrease in the elon-
gation.
• After CONFORM SPD, the specimens were rotary
swaged, which brought their strength to 1070 MPa
and the yield stress to 950 MPa. In the initial con-
dition, the ultimate strength was 480 MPa and the
yield stress was 370 MPa.
Acknowledgment
These results were obtained under the project entitled
"Development of West-Bohemian Centre of Materials
and Metallurgy" No. lo1412, financed by the Ministry of
Education of the Czech Republic.
7 REFERENCES
1 R. Z. Valiev, R. K. Islamgaliev, I. V. Alexandrov, Bulk nanostruc-
tured materials from severe plastic deformation, Progress in
Materials Science, 45 (2000) 2, doi:10.1016/S0079-6425(99)00007-9
2 N. J. Petch, The cleavage strength of polycrystals, J. Iron Steel Ins,
174 (1953), 25–28
3 E. O. Hall, The Deformation and Ageing of Mild Steel: III Dis-
cussion of Results, Proc. Phys. Soc. London, 64 (1951), 747–753,
doi:10.1088/0370-1301/64/9/303
4 A. Mishra, B. Kad, F. Gregori, M. Meyers, Microstructural evolution
in copper subjected to severe plastic deformation, Acta Materialia,
55 (2007), 2, doi:10.1016/j.actamat.2006.07.008
5 M. Duchek, T. Kubina, J. Hodek, J. Dlouhy, Development of the pro-
duction of ultrafine-grained titanium with the conform equipment,
Mater. Tehnol., 47 (2013), 4
6 G. J. Raab, R. Z. Valiev, T. C. Lowe, Y. T. Zhu, Continuous pro-
cessing of ultrafine grained Al by ECAP–Conform, Materials
Science and Engineering, 382 (2004) 1/2, doi:10.1016/j.msea.
2004.04.021
7 T. Kubina, J. Dlouhý, M. Kover, Preparation and thermal stability of
ultra-fine and nano-grained commercially pure titanium wires using
CONFORM equipment, Mater. Tehnol., 49 (2015), 2, doi:10.17222/
mit.2013.226
8 L. Ostrovska, L. Vistejnova, J. Dzugan, P. Slama, T. Kubina, E.
Ukraintsev, D. Kubies, M. Kralickova, M. Kalbacova, Biological
evaluation of ultra-fine titanium with improved mechanical strength
for dental implant engineering, J. Mater. Sci., 23 (2015),
doi:10.1007/s10853-015-9619-3
9 M. Zemko, T. Kubina, J. Dlouhý, J. Hodek, Technological aspects of
preparation of nanostructured titanium wire using a CONFORM
machine, IOP Conference Series: Materials Science and Engineering,
63 (2014), 2, doi:10.1088/1757-899X/63/1/012049
10 D. V. Gunderov, A. V. Polyakov, I. P. Semenova, Evolution of
microstructure, macrotexture and mechanical properties of
commercially pure Ti during ECAP-conform processing and
drawing, Materials Science and Engineering, 562 (2013), 128–136,
doi:10.1016/j.msea.2012.11.007
11 I. P. Semenova, R. Z. Valiev, E. B. Yakushima, G. H. Salimgareeva,
T. C. Lowe, Strength and fatigue properties enhancement in ultra-
fine-grained Ti produced by severe plastic deformation, Journal of
Materials Science, 43 (2008), 23–24, doi:10.1007/s10853-008-
2984-4
12 H. Alkhazraji, E. El-Danaf, M. Wollmann, L. Wagner, Enhanced
Fatigue Strength of Commercially Pure Ti Processed by Rotary
Swaging, Advances in Materials Science and Engineering, (2015),
1–12, doi:10.1155/2015/301837
J. PALÁN et al.: EFFECT OF SEVERE PLASTIC AND HEAVY COLD DEFORMATION ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 849–853 853
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
4 TEXTURE EVALUATION
The preferred orientation of crystallites expressed
with the texture index Ti (Equation 1, Table 3) is, in all
the cases, in the [001] direction perpendicular to the
sample surface analysed. The highest value was found
for CONSPD – 3 passes + RS (75 %) titanium sample,
i.e., the sample that was processed using the CONFORM
SPD device and rotary swaging (basal texture, Figure 5).
An increased value of the texture index can be seen in
the initial state, due to the previous wire-drawing process
(Figure 4). The texture index is lower for the samples
that were processed with the CONFORM SPD device.
Texture indices for the [101] direction are listed for the
sake of comparison. Nevertheless, all their values are
lower than 1. It means that no preferred orientation of
crystallites in the [101] direction is present, as opposed
to the sample surface.
Table 3: Texture evaluation using X-ray diffraction
Plane (h k l)
Texture index
as received CONSPD – 1pass
CONSPD – 2
passes
Ti (002) 3.13 2.49 2.67
Ti (101) 0.66 0.69 0.86
CONSPD – 3
passes
CONSPD – 3
passes + RS
(75 %)
Standard
Ti (002) 2.37 3.48 1
Ti (101) 0.97 0.24 1
4 MECHANICAL PROPERTIES
Table 4 indicates that the largest increase in the
mechanical properties was obtained during the first pass
through CONFORM SPD. The subsequent passes led to
smaller increments. It is important to note that the
increase in the ultimate strength and yield stress upon
CONFORM SPD is not offset by a decrease in the
ductility. The improvement in the mechanical properties
is mainly due to the grain refinement and the increased
dislocation density, which impede the dislocation move-
ment. A further increase in the mechanical properties
was obtained with rotary swaging applied after the three
passes through the CONFORM SPD machine. In compa-
rison with the CONFORM SPD process, rotary swaging
leads to a decrease in the ductility as a result of work
hardening.
Table 4: Mechanical properties in the as-received condition, after
CONFORM SPD processing and after RS
Condition
0.2
OYS UTS A5 RA
MPa MPa % %
as received 370 480 25 52
CONSPD – 1 pass 540 580 23 62
CONSPD – 2 passes 560 600 23 62
CONSPD – 3 passes 570 623 20 64
CONSPD – 3 passes + RS
(50 %) 830 885 13 54
CONSPD – 3 passes + RS
(75 %) 930 1000 12 57
CONSPD – 3 passes + RS
(85%) 950 1070 12 58
RS only (85 %) 780 850 12 50
5 DISCUSSION
The CONFORM SPD processing led to higher ulti-
mate strengths and yield stresses without a reduced
ductility. This is characteristic of the process of the for-
mation of a UFG equiaxed structure shown in Figure 2.
The grain size and distortion are non-uniform due to
non-uniform deformation during the forming process.
The preferred orientation of crystallites in the [001]
direction perpendicular to the surfaces of the investigated
samples is not too significant, except for the sample
which was processed with CONFORM SPD and the
rotary swaging machine where the highest value of the
texture index was observed (the basal texture).
Major increases in the ultimate strength and yield
stress were found after the first pass through the CON-
J. PALÁN et al.: EFFECT OF SEVERE PLASTIC AND HEAVY COLD DEFORMATION ...
852 Materiali in tehnologije / Materials and technology 51 (2017) 5, 849–853
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Diffraction radial profile of the specimen with the strongest
texture: CONSPD – 3 passes + rotary swaging (75 %)
Figure 4: Diffraction radial profile of the as-received specimen
FORM SPD machine, whereas the increments from the
subsequent passes were small. This is closely related to
the grain size, which did not decrease during the sub-
sequent passes. It is possible to say that the additional
deformation causes the subgrains to rotate into high-
angle grain boundaries, typically with an equiaxed shape
of the grains.5,7,9 The subsequent rotary swaging led to a
further increase in the mechanical properties, but the
strengthening mechanism was different this time. Instead
of the formation of equiaxed grains, the grains became
elongated in the direction of forming, with a much
higher dislocation density (Figure 3).10,11 When com-
pared to the CONFORM SPD route, rotary swaging led
to a reduced ductility as a consequence of work harden-
ing.
6 CONCLUSION
This study involved the processing of Ti Grade 2
using CONFORM SPD and rotary swaging with the goal
of improving its mechanical properties. The results and
findings are described below:
• The CONFORM SPD processing substantially
refined the grain from its initial size of 5390 nm to
350 nm. The subsequent rotary swaging did not
provide for a further grain refinement but led to a
grain elongation and to a higher dislocation density.
• The processing led to the basal texture. The most
significant form of texture was found in the specimen
processed with CONFORM SPD and rotary swaging.
• Three passes through the CONFORM SPD machine
led to a strength of 623 MPa and a yield stress of
570 MPa without any notable decrease in the elon-
gation.
• After CONFORM SPD, the specimens were rotary
swaged, which brought their strength to 1070 MPa
and the yield stress to 950 MPa. In the initial con-
dition, the ultimate strength was 480 MPa and the
yield stress was 370 MPa.
Acknowledgment
These results were obtained under the project entitled
"Development of West-Bohemian Centre of Materials
and Metallurgy" No. lo1412, financed by the Ministry of
Education of the Czech Republic.
7 REFERENCES
1 R. Z. Valiev, R. K. Islamgaliev, I. V. Alexandrov, Bulk nanostruc-
tured materials from severe plastic deformation, Progress in
Materials Science, 45 (2000) 2, doi:10.1016/S0079-6425(99)00007-9
2 N. J. Petch, The cleavage strength of polycrystals, J. Iron Steel Ins,
174 (1953), 25–28
3 E. O. Hall, The Deformation and Ageing of Mild Steel: III Dis-
cussion of Results, Proc. Phys. Soc. London, 64 (1951), 747–753,
doi:10.1088/0370-1301/64/9/303
4 A. Mishra, B. Kad, F. Gregori, M. Meyers, Microstructural evolution
in copper subjected to severe plastic deformation, Acta Materialia,
55 (2007), 2, doi:10.1016/j.actamat.2006.07.008
5 M. Duchek, T. Kubina, J. Hodek, J. Dlouhy, Development of the pro-
duction of ultrafine-grained titanium with the conform equipment,
Mater. Tehnol., 47 (2013), 4
6 G. J. Raab, R. Z. Valiev, T. C. Lowe, Y. T. Zhu, Continuous pro-
cessing of ultrafine grained Al by ECAP–Conform, Materials
Science and Engineering, 382 (2004) 1/2, doi:10.1016/j.msea.
2004.04.021
7 T. Kubina, J. Dlouhý, M. Kover, Preparation and thermal stability of
ultra-fine and nano-grained commercially pure titanium wires using
CONFORM equipment, Mater. Tehnol., 49 (2015), 2, doi:10.17222/
mit.2013.226
8 L. Ostrovska, L. Vistejnova, J. Dzugan, P. Slama, T. Kubina, E.
Ukraintsev, D. Kubies, M. Kralickova, M. Kalbacova, Biological
evaluation of ultra-fine titanium with improved mechanical strength
for dental implant engineering, J. Mater. Sci., 23 (2015),
doi:10.1007/s10853-015-9619-3
9 M. Zemko, T. Kubina, J. Dlouhý, J. Hodek, Technological aspects of
preparation of nanostructured titanium wire using a CONFORM
machine, IOP Conference Series: Materials Science and Engineering,
63 (2014), 2, doi:10.1088/1757-899X/63/1/012049
10 D. V. Gunderov, A. V. Polyakov, I. P. Semenova, Evolution of
microstructure, macrotexture and mechanical properties of
commercially pure Ti during ECAP-conform processing and
drawing, Materials Science and Engineering, 562 (2013), 128–136,
doi:10.1016/j.msea.2012.11.007
11 I. P. Semenova, R. Z. Valiev, E. B. Yakushima, G. H. Salimgareeva,
T. C. Lowe, Strength and fatigue properties enhancement in ultra-
fine-grained Ti produced by severe plastic deformation, Journal of
Materials Science, 43 (2008), 23–24, doi:10.1007/s10853-008-
2984-4
12 H. Alkhazraji, E. El-Danaf, M. Wollmann, L. Wagner, Enhanced
Fatigue Strength of Commercially Pure Ti Processed by Rotary
Swaging, Advances in Materials Science and Engineering, (2015),
1–12, doi:10.1155/2015/301837
J. PALÁN et al.: EFFECT OF SEVERE PLASTIC AND HEAVY COLD DEFORMATION ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 849–853 853
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
A. G. ILLARIONOV et al.: EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON THE STRUCTURE ...
855–859
EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON
THE STRUCTURE AND PROPERTIES OF WELD JOINTS OF A
TWO-PHASE TITANIUM ALLOY
U^INEK MIKROLEGIRANJA ITRIJA IN CIRKONIJA NA
STRUKTURO IN LASTNOSTI NA SPOJE ZAVROV DVOFAZNE
ZLITINE TITANA
Anatoly Illarionov, Artemy Popov, Svetlana Illarionova, Dmitry Gadeev
Ural Federal University, 620034, 28 Mira str., Yekaterinburg, Russia
d.v.gadeev@urfu.ru
Prejem rokopisa – received: 2016-11-10; sprejem za objavo – accepted for publication: 2017-04-19
doi:10.17222/mit.2016.317
The effect of microalloying of the Ti-4.8Al-1.2Mo-2.6V-0.6Cr–0.25Fe titanium alloy with yttrium and zirconium on the phase
composition, structure and mechanical properties of welded sheets was studied using light and transmission electron
microscopy, X-ray diffraction analysis and microindentation-hardness testing. A differential thermal analysis was employed to
model the welding thermal cycle. It was found that alloying with yttrium led to the precipitation of Y2O3 oxide particles
resulting in refining the microstructure of the alloy. In addition, yttrium additions were shown to stabilise the -phase of the
alloy that decreases the hardness of the alloy.
Keywords: microalloying, rare-earth elements, titanium alloy, welding
U~inek mikrolegiranja zlitine titana Ti-4.8Al-1.2Mo-2.6V-0.6Cr-0.25Fe z itrijem in cirkonijem na sestavo faze, strukturne in
mehanske lastnosti je bil raziskovan s svetlobno elektronsko mikroskopijo, z rentgensko difrakcijsko analizo in s testiranjem
trdote mikroindentacije. Uporabili smo diferencialno termi~no analizo za pridobitev modela za varilni termi~ni cikel.
Ugotovljeno je bilo, da legiranje z itrijem vodi k razpr{enosti Y2O3 oksidnih delcev, kar se ka`e pri rafiniranju mikrostrukture
zlitine. Poleg tega je bilo dokazano, da so dodatki itrija stabilizirali -fazo v zlitini, ki zmanj{uje trdoto zlitine.
Klju~ne besede: mikrolegiranje, elementi redke zemlje, zlitina titana, varjenje
1 INTRODUCTION
Titanium alloys, due to their high specific strength
and good corrosion resistance, are widely used in the
aerospace and marine industries as structural materials
and it is desirable to further enhance their properties. The
latter is possible by means of complex multicomponent
alloying.1,2 In addition to the specific strength, good
weldability is usually necessary to maximise the material
utilisation.
Since -titanium and + alloys are considered to
have a better weldability compared to metastable -tita-
nium alloys,3 a number of +-alloys were developed in
Russia to meet these requirements. One of them is
Ti-4.8Al-1.2Mo-2.6V-0.6Cr–0.25Fe, which is a typical
martensite-type two-phase alloy.4
Despite good weldability, the welding usually
negatively affects the mechanical properties of titanium
alloys.5 This is partly due to the oxygen absorbed from
the shield atmosphere followed by its diffusion to the
weld-fusion zone (FZ), partly due to a high heat input
resulting in significant grain coarsening. One way to
overcome such a problem is to refine the prior grain size
by the grain-boundary pinning effect of second-phase
particles and promote heterogeneous nucleation in the
FZ.6 It was shown that microalloying with transition
elements, especially rare-earth metals, is a promising
way to refine the structure of titanium alloys.7 One of the
most frequently employed alloying elements is
yttrium.8–11
Although it was already shown what microalloying of
the Ti-4.8Al-1.2Mo-2.6V-0.6Cr–0.25Fe alloy with
0.06 % mass fraction of yttrium decreases the oxygen
concentration in - and -solid solutions due to the
formation of Y2O3 particles,4 detailed studies of a weld
structure were not conducted. This is why our study
aimed at investigating the influence of microalloying
titanium alloys with zirconium and yttrium on the
microstructure formation and mechanical properties of
the weld joints of this alloy.
2 EXPERIMENTAL PART
Two-millimeter-thick hot-rolled sheets of an (+)
martensitic titanium alloy alloyed with yttrium (alloy 1)
and zirconium (alloy 2) were used in this study. The
chemical compositions of the alloys are shown in
Table 1. Prior to their use, the sheets were subjected to
vacuum annealing at 750 °C for 1 h followed by arc
welding using non-consumable tungsten electrodes.
Materiali in tehnologije / Materials and technology 51 (2017) 5, 855–859 855
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 669.794:669.296:67.017:66.046.51 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)855(2017)
A. G. ILLARIONOV et al.: EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON THE STRUCTURE ...
855–859
EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON
THE STRUCTURE AND PROPERTIES OF WELD JOINTS OF A
TWO-PHASE TITANIUM ALLOY
U^INEK MIKROLEGIRANJA ITRIJA IN CIRKONIJA NA
STRUKTURO IN LASTNOSTI NA SPOJE ZAVROV DVOFAZNE
ZLITINE TITANA
Anatoly Illarionov, Artemy Popov, Svetlana Illarionova, Dmitry Gadeev
Ural Federal University, 620034, 28 Mira str., Yekaterinburg, Russia
d.v.gadeev@urfu.ru
Prejem rokopisa – received: 2016-11-10; sprejem za objavo – accepted for publication: 2017-04-19
doi:10.17222/mit.2016.317
The effect of microalloying of the Ti-4.8Al-1.2Mo-2.6V-0.6Cr–0.25Fe titanium alloy with yttrium and zirconium on the phase
composition, structure and mechanical properties of welded sheets was studied using light and transmission electron
microscopy, X-ray diffraction analysis and microindentation-hardness testing. A differential thermal analysis was employed to
model the welding thermal cycle. It was found that alloying with yttrium led to the precipitation of Y2O3 oxide particles
resulting in refining the microstructure of the alloy. In addition, yttrium additions were shown to stabilise the -phase of the
alloy that decreases the hardness of the alloy.
Keywords: microalloying, rare-earth elements, titanium alloy, welding
U~inek mikrolegiranja zlitine titana Ti-4.8Al-1.2Mo-2.6V-0.6Cr-0.25Fe z itrijem in cirkonijem na sestavo faze, strukturne in
mehanske lastnosti je bil raziskovan s svetlobno elektronsko mikroskopijo, z rentgensko difrakcijsko analizo in s testiranjem
trdote mikroindentacije. Uporabili smo diferencialno termi~no analizo za pridobitev modela za varilni termi~ni cikel.
Ugotovljeno je bilo, da legiranje z itrijem vodi k razpr{enosti Y2O3 oksidnih delcev, kar se ka`e pri rafiniranju mikrostrukture
zlitine. Poleg tega je bilo dokazano, da so dodatki itrija stabilizirali -fazo v zlitini, ki zmanj{uje trdoto zlitine.
Klju~ne besede: mikrolegiranje, elementi redke zemlje, zlitina titana, varjenje
1 INTRODUCTION
Titanium alloys, due to their high specific strength
and good corrosion resistance, are widely used in the
aerospace and marine industries as structural materials
and it is desirable to further enhance their properties. The
latter is possible by means of complex multicomponent
alloying.1,2 In addition to the specific strength, good
weldability is usually necessary to maximise the material
utilisation.
Since -titanium and + alloys are considered to
have a better weldability compared to metastable -tita-
nium alloys,3 a number of +-alloys were developed in
Russia to meet these requirements. One of them is
Ti-4.8Al-1.2Mo-2.6V-0.6Cr–0.25Fe, which is a typical
martensite-type two-phase alloy.4
Despite good weldability, the welding usually
negatively affects the mechanical properties of titanium
alloys.5 This is partly due to the oxygen absorbed from
the shield atmosphere followed by its diffusion to the
weld-fusion zone (FZ), partly due to a high heat input
resulting in significant grain coarsening. One way to
overcome such a problem is to refine the prior grain size
by the grain-boundary pinning effect of second-phase
particles and promote heterogeneous nucleation in the
FZ.6 It was shown that microalloying with transition
elements, especially rare-earth metals, is a promising
way to refine the structure of titanium alloys.7 One of the
most frequently employed alloying elements is
yttrium.8–11
Although it was already shown what microalloying of
the Ti-4.8Al-1.2Mo-2.6V-0.6Cr–0.25Fe alloy with
0.06 % mass fraction of yttrium decreases the oxygen
concentration in - and -solid solutions due to the
formation of Y2O3 particles,4 detailed studies of a weld
structure were not conducted. This is why our study
aimed at investigating the influence of microalloying
titanium alloys with zirconium and yttrium on the
microstructure formation and mechanical properties of
the weld joints of this alloy.
2 EXPERIMENTAL PART
Two-millimeter-thick hot-rolled sheets of an (+)
martensitic titanium alloy alloyed with yttrium (alloy 1)
and zirconium (alloy 2) were used in this study. The
chemical compositions of the alloys are shown in
Table 1. Prior to their use, the sheets were subjected to
vacuum annealing at 750 °C for 1 h followed by arc
welding using non-consumable tungsten electrodes.
Materiali in tehnologije / Materials and technology 51 (2017) 5, 855–859 855
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 669.794:669.296:67.017:66.046.51 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)855(2017)
Table 1: Chemical compositions of the test materials, in mass
fractions (w/%)
Materials
Alloying elements, in mass fractions (w/%)
Ti Al Mo V Cr Fe Y Zr
Alloy 1 Bal 4.80 1.20 2.60 0.60 0.60 0.06 -
Alloy 2 - 0.07
The microstructures of the alloys were analyzed with
light optical (LOM) and transmission electron (TEM)
microscopy on Olympus GX51 and JEM 200C micro-
scopes, respectively. A qualitative phase analysis was
carried out using X-ray diffraction (XRD) with Cu-K
radiation on a DRON-3M powder diffractometer. A
differential thermal analysis (DTA) was performed on a
Du Pont appliance in DTA-1600 crucibles. The average
grain size and grain-size distribution were determined
with an Epiquant optical microscope.
Vickers microhardness was measured at a 1 N load in
accordance with the formula below (ISO 6507) and
multiplied with the standard gravity (g = 9.81) to get the
values in MPa in Equation (1):
HV = 0.1891· (F/d2) (1)
with F being the applied load (Newton) and d being the
average length of the diagonal of the residual indent
(millimeters).
Metallographic samples were prepared according to
the standard procedures. The steps consisted of grinding
with 120–2400 grit SiC paper, polishing with 0.3 μm
colloidal alumina, followed by the final polishing with a
0.05 μm colloidal silica suspension. Kroll’ s reagent with
a composition of 100 mL water + 2 mL HF + 5 mL
HNO3 was used to etch the specimens.12 Thin foils for
the TEM investigation were prepared with electrolytic
polishing in a methanol-containing electrolyte (300 mL
methanol, 175 mL butanol, 30 mL perchloric acid
(70–72 %)) at -30 °C.12
3 RESULTS AND DISCUSSION
With respect to the differences in the microstructure,
weld joints generally consist of three distinct zones, i.e.,
the base metal, the fusion zone and the heat-affected
zone. In the fusion zone (FZ), materials melt during the
welding and crystallize during the subsequent cooling.
The heat-affected zone (HAZ) corresponds to the narrow
transition region between the base metal and the fusion
zone where the material is heated up to sub-transus or
even superheated above the critical temperature. The
latter might obviously cause significant microstructure
changes in comparison with the base metal.
On the present samples, the zone width was
measured and it was around 12 mm and 4.5 mm for the
fusion and heat-affected zones, respectively (Figure 1).
The base-metal microstructure of both alloys is
represented by primary -phase precipitates of a mixed,
globular and lamellar, morphology in a -phase matrix
(Figure 2a). The yttrium-containing alloy 1 is dis-
tinguished by a dispersive oxide precipitation of Y2O3
particles with the average diameter of up to 120 nm
(Figure 2b). These particles were found to be almost
spherical and to precipitate mainly near the grain
boundaries. The TEM micrograph (Figure 2b) shows a
specific 'comet-tail-like' contrast around these particles,
which is attributed to the elastic-stress field caused by
the difference between the thermal-expansion coeffi-
cients of the oxides and the matrix.
The average -phase grain size decreased with the
increasing distance from the FZ (Figure 1) due to the
higher temperatures achieved in the HAZ close to the
A. G. ILLARIONOV et al.: EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON THE STRUCTURE ...
856 Materiali in tehnologije / Materials and technology 51 (2017) 5, 855–859
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Microstructures of base metals: a) alloy 2 (LOM); b) alloy 1
(TEM)
Figure 1: Structures of welded joints: a) alloy 1 with Y; b) alloy 2
with Zr
FZ. The coarsest grains of alloy 2 reached about 400 μm
in diameter, whereas in alloy 1, they did not exceed 330
μm, i.e., the yttrium-containing alloy 1 was characterized
by a 1.2 times finer grain size. A similar relationship was
previously observed in the solution-treated and water-
quenched alloys of the same composition4 and it was
attributed to the presence of the yttrium-oxide particles
in the Y-containing alloy that inhibited the grain-boun-
dary movement.
The X-ray diffraction analysis showed that the vo-
lume fractions of the - and -phases varied significantly
in different parts of the HAZ (Figure 3). The intensity of
-phase peaks in the XRD patterns decreased with the
increasing distance from the FZ, clearly indicating that
the volume fraction of the -phase was significantly
higher in the near-FZ areas. At the same time, the a
lattice parameter changed in the opposite direction.
It can be seen that the (200)-peak intensity ratio for
alloy 1 was considerably higher than that for alloy 2.
This confirms the assumption that yttrium binds oxygen,
thus increasing the stability of the supercooled -phase
upon cooling.4,8–9 Furthermore, such a behaviour resulted
in a higher -phase volume fraction, decreasing the hard-
ness of the alloy. This was confirmed with the micro-
hardness-indentation tests. Alloy 1 had a lower micro-
hardness (2800±50 MPa) than alloy 2 (3300±50 MPa). It
was also found that the -phase morphology along the
HAZ changed from a complex one (Figure 2) to a
lamellar one consisting of thin platelets (Figure 4).
The analysis of the microstructures of the alloys
revealed a noticeable feature of both alloys, i.e., a jagged
shape of the grain boundaries. A similar phenomenon
was seen previously, for instance, in Ni–Mn–In-based
alloys after thermal cycling13 and near-alpha titanium
alloys14 and it was attributed to the formation of marten-
site crystals, which grow towards the grain boundaries
curving them. Figure 4a shows that such a shape coin-
cides closely with the lamellar -phase precipitations on
both sides of the grain boundaries. Thin -lamellae were
formed during the cooling from the single-phase
-region. The stress field next to the growing -lamellae
could interact with the grain boundaries curving them.
The same process occurred in the neighbouring grains
resulting in the jag formation. Since this process can
occur in preferentially oriented grains, the jags were
found only on the corresponding parts of the grains. On
the whole, we suppose that the jagged shape of the grain
boundaries is associated with the interaction of the
boundaries with the products of the -transformation
occurring upon the cooling.
To summarize the above, it can be concluded that the
welding thermal cycle results in significant changes of
the grain structure, phase composition and morphology
of the phases. Yttrium binds oxygen forming oxide
particles that inhibit the grain growth and improve the
-phase stability. In turn, microalloying the alloy with
zirconium does not result in the formation of additional
phases and leads to the solution hardening of the alloy.
The average grain sizes for the weld joints of the
samples of alloys 1 and 2 did not show any significant
difference and was about 480–600 μm (Figure 1). The
fine structure of the alloys mainly consists of the
-phase, with dislocations perpendicular to the lamellae
(Figure 5b).
The -lamellae in the FZs of the alloys tend to form
small colonies that may cross differently oriented
lamellae (Figure 5c). Interlamellar / spacings in
A. G. ILLARIONOV et al.: EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON THE STRUCTURE ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 855–859 857
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: -phase lattice parameter (a) and (200) XRD-peak inten-
sity ratio (IHAZ /IFZ) as a function of distance from the fusion zone
of a welded joint
Figure 4: Microstructures of heat-affected zones of weld joints:
a) alloy 1, b) alloy 2
Table 1: Chemical compositions of the test materials, in mass
fractions (w/%)
Materials
Alloying elements, in mass fractions (w/%)
Ti Al Mo V Cr Fe Y Zr
Alloy 1 Bal 4.80 1.20 2.60 0.60 0.60 0.06 -
Alloy 2 - 0.07
The microstructures of the alloys were analyzed with
light optical (LOM) and transmission electron (TEM)
microscopy on Olympus GX51 and JEM 200C micro-
scopes, respectively. A qualitative phase analysis was
carried out using X-ray diffraction (XRD) with Cu-K
radiation on a DRON-3M powder diffractometer. A
differential thermal analysis (DTA) was performed on a
Du Pont appliance in DTA-1600 crucibles. The average
grain size and grain-size distribution were determined
with an Epiquant optical microscope.
Vickers microhardness was measured at a 1 N load in
accordance with the formula below (ISO 6507) and
multiplied with the standard gravity (g = 9.81) to get the
values in MPa in Equation (1):
HV = 0.1891· (F/d2) (1)
with F being the applied load (Newton) and d being the
average length of the diagonal of the residual indent
(millimeters).
Metallographic samples were prepared according to
the standard procedures. The steps consisted of grinding
with 120–2400 grit SiC paper, polishing with 0.3 μm
colloidal alumina, followed by the final polishing with a
0.05 μm colloidal silica suspension. Kroll’ s reagent with
a composition of 100 mL water + 2 mL HF + 5 mL
HNO3 was used to etch the specimens.12 Thin foils for
the TEM investigation were prepared with electrolytic
polishing in a methanol-containing electrolyte (300 mL
methanol, 175 mL butanol, 30 mL perchloric acid
(70–72 %)) at -30 °C.12
3 RESULTS AND DISCUSSION
With respect to the differences in the microstructure,
weld joints generally consist of three distinct zones, i.e.,
the base metal, the fusion zone and the heat-affected
zone. In the fusion zone (FZ), materials melt during the
welding and crystallize during the subsequent cooling.
The heat-affected zone (HAZ) corresponds to the narrow
transition region between the base metal and the fusion
zone where the material is heated up to sub-transus or
even superheated above the critical temperature. The
latter might obviously cause significant microstructure
changes in comparison with the base metal.
On the present samples, the zone width was
measured and it was around 12 mm and 4.5 mm for the
fusion and heat-affected zones, respectively (Figure 1).
The base-metal microstructure of both alloys is
represented by primary -phase precipitates of a mixed,
globular and lamellar, morphology in a -phase matrix
(Figure 2a). The yttrium-containing alloy 1 is dis-
tinguished by a dispersive oxide precipitation of Y2O3
particles with the average diameter of up to 120 nm
(Figure 2b). These particles were found to be almost
spherical and to precipitate mainly near the grain
boundaries. The TEM micrograph (Figure 2b) shows a
specific 'comet-tail-like' contrast around these particles,
which is attributed to the elastic-stress field caused by
the difference between the thermal-expansion coeffi-
cients of the oxides and the matrix.
The average -phase grain size decreased with the
increasing distance from the FZ (Figure 1) due to the
higher temperatures achieved in the HAZ close to the
A. G. ILLARIONOV et al.: EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON THE STRUCTURE ...
856 Materiali in tehnologije / Materials and technology 51 (2017) 5, 855–859
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Microstructures of base metals: a) alloy 2 (LOM); b) alloy 1
(TEM)
Figure 1: Structures of welded joints: a) alloy 1 with Y; b) alloy 2
with Zr
FZ. The coarsest grains of alloy 2 reached about 400 μm
in diameter, whereas in alloy 1, they did not exceed 330
μm, i.e., the yttrium-containing alloy 1 was characterized
by a 1.2 times finer grain size. A similar relationship was
previously observed in the solution-treated and water-
quenched alloys of the same composition4 and it was
attributed to the presence of the yttrium-oxide particles
in the Y-containing alloy that inhibited the grain-boun-
dary movement.
The X-ray diffraction analysis showed that the vo-
lume fractions of the - and -phases varied significantly
in different parts of the HAZ (Figure 3). The intensity of
-phase peaks in the XRD patterns decreased with the
increasing distance from the FZ, clearly indicating that
the volume fraction of the -phase was significantly
higher in the near-FZ areas. At the same time, the a
lattice parameter changed in the opposite direction.
It can be seen that the (200)-peak intensity ratio for
alloy 1 was considerably higher than that for alloy 2.
This confirms the assumption that yttrium binds oxygen,
thus increasing the stability of the supercooled -phase
upon cooling.4,8–9 Furthermore, such a behaviour resulted
in a higher -phase volume fraction, decreasing the hard-
ness of the alloy. This was confirmed with the micro-
hardness-indentation tests. Alloy 1 had a lower micro-
hardness (2800±50 MPa) than alloy 2 (3300±50 MPa). It
was also found that the -phase morphology along the
HAZ changed from a complex one (Figure 2) to a
lamellar one consisting of thin platelets (Figure 4).
The analysis of the microstructures of the alloys
revealed a noticeable feature of both alloys, i.e., a jagged
shape of the grain boundaries. A similar phenomenon
was seen previously, for instance, in Ni–Mn–In-based
alloys after thermal cycling13 and near-alpha titanium
alloys14 and it was attributed to the formation of marten-
site crystals, which grow towards the grain boundaries
curving them. Figure 4a shows that such a shape coin-
cides closely with the lamellar -phase precipitations on
both sides of the grain boundaries. Thin -lamellae were
formed during the cooling from the single-phase
-region. The stress field next to the growing -lamellae
could interact with the grain boundaries curving them.
The same process occurred in the neighbouring grains
resulting in the jag formation. Since this process can
occur in preferentially oriented grains, the jags were
found only on the corresponding parts of the grains. On
the whole, we suppose that the jagged shape of the grain
boundaries is associated with the interaction of the
boundaries with the products of the -transformation
occurring upon the cooling.
To summarize the above, it can be concluded that the
welding thermal cycle results in significant changes of
the grain structure, phase composition and morphology
of the phases. Yttrium binds oxygen forming oxide
particles that inhibit the grain growth and improve the
-phase stability. In turn, microalloying the alloy with
zirconium does not result in the formation of additional
phases and leads to the solution hardening of the alloy.
The average grain sizes for the weld joints of the
samples of alloys 1 and 2 did not show any significant
difference and was about 480–600 μm (Figure 1). The
fine structure of the alloys mainly consists of the
-phase, with dislocations perpendicular to the lamellae
(Figure 5b).
The -lamellae in the FZs of the alloys tend to form
small colonies that may cross differently oriented
lamellae (Figure 5c). Interlamellar / spacings in
A. G. ILLARIONOV et al.: EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON THE STRUCTURE ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 855–859 857
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: -phase lattice parameter (a) and (200) XRD-peak inten-
sity ratio (IHAZ /IFZ) as a function of distance from the fusion zone
of a welded joint
Figure 4: Microstructures of heat-affected zones of weld joints:
a) alloy 1, b) alloy 2
alloy 2 have a complex structure and high dislocation
density (Figure 5d). -phase interlayers are relatively
wide (up to 200 nm) and non-uniform throughout the
bulk of the alloy. Secondary -phase precipitations were
observed in the interlamellar spacings, both along and
across the  interlayers as well as at the /-interphase
boundaries (Figure 5d).
Generally, the coarse -platelets with a small dislo-
cation density prevail in both yttrium- and zirconium-
containing alloys (Figures 5e, 5f). However, alloy 1 is
distinguished by a more conventional structure of inter-
lamellar spacing. Although the thin lenticular platelets of
the secondary -phase were still present, no highly
dispersive precipitates were found on the interphase
boundaries and within the -layers. It should be noted
that no Y2O3 particles were detected in the FZ. This
might be connected with the intense overheating of the
alloy caused by welding that resulted in a dissociation of
the oxide.
A. G. ILLARIONOV et al.: EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON THE STRUCTURE ...
858 Materiali in tehnologije / Materials and technology 51 (2017) 5, 855–859
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Thin structures of the weld joints of alloys 1 and 2: a), c), d) alloy 1; b), e), f) alloy 2
Figure 6: DTA curves for continuous cooling (20 °C/min cooling rate)
from 1200 °C: a) alloy 1, b) alloy 2
In order to model the welding thermal cycle, a num-
ber of DTA experiments involving continuous heating up
to 1200 °C and the subsequent cooling to 20 °C/min
were carried out. Continuous-cooling DTA curves
allowed us to determine the temperature ranges of the
-transformation (Figure 6).
A comparative analysis of the curves showed that the
-phase decomposition process in alloy 2 consisted of
three distinctive stages, whereas only two exothermal
effects were found for alloy 1. Stages I and II (at
920–750 and 700–550 °C, respectively) were attributed
to the formation of coarse primary -colonies with a low
defect density and secondary -platelets within the
-matrix. Stage III (500–400 °C) was assigned to the
precipitation of the highly dispersed -precipitates
mentioned above (Figure 5 d).
The formation of the -phase upon supercooling
below the transus temperature of the titanium alloys is
initiated by the crystallographic defects of the structure
such as grain boundaries, dislocations, etc. Yttrium binds
oxygen and thus refines the alloy. The products of the
-phase decomposition were quite free of precipitates.
On the other hand, the precipitation of the secondary
-platelets was possible in both alloys studied during
supercooling of up to about 600 °C. At lower tem-
peratures, the -phase was stabilized in alloy 1, whereas
its further decomposition occurred in alloy 2 resulting in
the formation of highly dispersed -precipitates. We
believe that all these differences in the behaviour of the
alloys are associated with different -phase stability
values, which are primarily defined by the concentration
of impurities in the alloys. Specifically, the interstitial
elements in alloy 1, mainly oxygen, are bound to yttrium
in the form of oxides. Zirconium has a lower affinity for
oxygen and thus oxygen atoms remain in the solid
solution leading to a more complete -phase decom-
position upon the cooling (Figure 2) resulting in a lower
-phase lattice parameter. With an increase in the volume
fraction of the body-centred cubic -phase, the strength
of the titanium alloys usually decreases.12 This is why
the microhardness of Y-containing alloy 1 was found to
be lower (3300±100 MPa for alloy 2) and 3000±100 MPa
for alloy 1).
4 CONCLUSIONS
The influence of the microalloying of titanium alloy
Ti–4.8À1–1.2Ìî–2.6V–0.6Ñr–0.25Fe with yttrium and
zirconium on the structure and microhardness was
considered. It was found that the alloying with yttrium
has the most noticeable effect on the structure of a weld
joint and the phase transformation of the alloy.
A yttrium addition refines the alloy by means of
binding with oxygen atoms and forming Y2O3 particles.
These particles pin the grain boundaries in both the
heat-affected and fusion zones, inhibiting the grain
growth during a welding thermal cycle. In addition, due
to a lower concentration of oxygen in the -phase,
yttrium decreases the -transus temperature of the alloy
leading to the stabilization of the supercooled -phase.
Due to a higher volume fraction of the "soft" -phase
in the yttrium-containing alloy, its microhardness was
shown to be lower than that of the alloy with zirconium.
Acknowledgment
We hereby acknowledge the support of the Ministry
of Science and Education of the Russian Federation, in
accordance with the decree of the Government of 9 April
2010, No. 218, project No. 03.G25.31.0234.
5 REFERENCES
1 A I. Khorev, Development of structural titanium alloys for
components and sections in aerospace technology, 10 (2010), 13–23,
doi:10.1080/09507116.2010.486188
2 A. I. Khorev, Complex alloying and heat treatment of titanium
alloys, 6 (2008), 364–368, doi:10.1080/09507110802288312
3 M. J. Donachie, Titanium, A Technical Guide, 2nd ed., ASM Interna-
tional, 2000
4 A. G. Illarionov, A. A. Popov, S. M. Illarionova, Effect of Micro-
alloying, with Rem Inclusively, on the Structure, Phase Composition
and Properties of ( + )-Titanium Alloy, Metal Science and Heat
Treatment, 57 (2016), 719–725, doi:10.1007/s11041-016-9948-0
5 E. W. Collings, The physical metallurgy of titanium alloys, American
Society for Metals, 1984, 261
6 Rios PR. Overview No. 62: A theory for grain boundary pinning by
particles, Acta Metall., 35 (1987) 12, 2805–2814
7 A. I. Khorev, Theory and practices of microalloying of near- and
+ titanium alloys with REE, zirconium, hafnium and rhenium,
Techologii machinostroeniya, 1 (2015), 5–10
8 H. Wu, Y. Han, X. Chen, Effects of Yttrium on Mechanical Pro-
perties and Microstructures of Ti-Si Eutectic Alloy, Chinese Journal
of Aeronautics, 18 (2005), 171–174, doi:10.1016/S1000-9361(11)
60324-5
9 R. P. Kolli, A. A. Herzing, S. Ankem, Characterization of yttrium-
rich precipitates in a titanium alloy weld, Materials Characterization,
122 (2016), 30–35, doi:10.1016/j.matchar.2016.10.014
10 B. Poorganji, A. Kazahari, T. Narushima, C. Ouchi, T. Furuhara,
Effect of yttrium addition on grain growth of ,  and + titanium
alloys, Journal of Physics: Conference Series, 240 (2010), 12170,
doi:10.1088/1742-6596/240/1/012170
11 W. F. Cui, C. M. Liu, L. Zhou, G. Z. Luo, Characteristics of
microstructures and second-phase particles in Y-bearing Ti-1100
alloy, Materials Science and Engineering: A, 323 (2002), 192–197,
doi:10.1016/S0921-5093(01)01362-4
12 G. Lütjering, J. C. Williams, Titanium, 2nd ed., Springer, Berlin,
2007, 442
13 Y. V. Kaletina, E. D. Efimova, E. G. Gerasimov, A. Y. Kaletin, Effect
of thermal cycling on structure and properties of Ni–Mn–In-based
alloys, Technical Physics, 12 (2016), 1894–1897
14 H. M. Flower, P. R. Swann, D. R. F. West, The effect of Si, Zr, Al
and Mo on the structure and strength of Ti martensite, Journal of
Materials Science, 8 (1972), 929–938
A. G. ILLARIONOV et al.: EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON THE STRUCTURE ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 855–859 859
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
alloy 2 have a complex structure and high dislocation
density (Figure 5d). -phase interlayers are relatively
wide (up to 200 nm) and non-uniform throughout the
bulk of the alloy. Secondary -phase precipitations were
observed in the interlamellar spacings, both along and
across the  interlayers as well as at the /-interphase
boundaries (Figure 5d).
Generally, the coarse -platelets with a small dislo-
cation density prevail in both yttrium- and zirconium-
containing alloys (Figures 5e, 5f). However, alloy 1 is
distinguished by a more conventional structure of inter-
lamellar spacing. Although the thin lenticular platelets of
the secondary -phase were still present, no highly
dispersive precipitates were found on the interphase
boundaries and within the -layers. It should be noted
that no Y2O3 particles were detected in the FZ. This
might be connected with the intense overheating of the
alloy caused by welding that resulted in a dissociation of
the oxide.
A. G. ILLARIONOV et al.: EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON THE STRUCTURE ...
858 Materiali in tehnologije / Materials and technology 51 (2017) 5, 855–859
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Thin structures of the weld joints of alloys 1 and 2: a), c), d) alloy 1; b), e), f) alloy 2
Figure 6: DTA curves for continuous cooling (20 °C/min cooling rate)
from 1200 °C: a) alloy 1, b) alloy 2
In order to model the welding thermal cycle, a num-
ber of DTA experiments involving continuous heating up
to 1200 °C and the subsequent cooling to 20 °C/min
were carried out. Continuous-cooling DTA curves
allowed us to determine the temperature ranges of the
-transformation (Figure 6).
A comparative analysis of the curves showed that the
-phase decomposition process in alloy 2 consisted of
three distinctive stages, whereas only two exothermal
effects were found for alloy 1. Stages I and II (at
920–750 and 700–550 °C, respectively) were attributed
to the formation of coarse primary -colonies with a low
defect density and secondary -platelets within the
-matrix. Stage III (500–400 °C) was assigned to the
precipitation of the highly dispersed -precipitates
mentioned above (Figure 5 d).
The formation of the -phase upon supercooling
below the transus temperature of the titanium alloys is
initiated by the crystallographic defects of the structure
such as grain boundaries, dislocations, etc. Yttrium binds
oxygen and thus refines the alloy. The products of the
-phase decomposition were quite free of precipitates.
On the other hand, the precipitation of the secondary
-platelets was possible in both alloys studied during
supercooling of up to about 600 °C. At lower tem-
peratures, the -phase was stabilized in alloy 1, whereas
its further decomposition occurred in alloy 2 resulting in
the formation of highly dispersed -precipitates. We
believe that all these differences in the behaviour of the
alloys are associated with different -phase stability
values, which are primarily defined by the concentration
of impurities in the alloys. Specifically, the interstitial
elements in alloy 1, mainly oxygen, are bound to yttrium
in the form of oxides. Zirconium has a lower affinity for
oxygen and thus oxygen atoms remain in the solid
solution leading to a more complete -phase decom-
position upon the cooling (Figure 2) resulting in a lower
-phase lattice parameter. With an increase in the volume
fraction of the body-centred cubic -phase, the strength
of the titanium alloys usually decreases.12 This is why
the microhardness of Y-containing alloy 1 was found to
be lower (3300±100 MPa for alloy 2) and 3000±100 MPa
for alloy 1).
4 CONCLUSIONS
The influence of the microalloying of titanium alloy
Ti–4.8À1–1.2Ìî–2.6V–0.6Ñr–0.25Fe with yttrium and
zirconium on the structure and microhardness was
considered. It was found that the alloying with yttrium
has the most noticeable effect on the structure of a weld
joint and the phase transformation of the alloy.
A yttrium addition refines the alloy by means of
binding with oxygen atoms and forming Y2O3 particles.
These particles pin the grain boundaries in both the
heat-affected and fusion zones, inhibiting the grain
growth during a welding thermal cycle. In addition, due
to a lower concentration of oxygen in the -phase,
yttrium decreases the -transus temperature of the alloy
leading to the stabilization of the supercooled -phase.
Due to a higher volume fraction of the "soft" -phase
in the yttrium-containing alloy, its microhardness was
shown to be lower than that of the alloy with zirconium.
Acknowledgment
We hereby acknowledge the support of the Ministry
of Science and Education of the Russian Federation, in
accordance with the decree of the Government of 9 April
2010, No. 218, project No. 03.G25.31.0234.
5 REFERENCES
1 A I. Khorev, Development of structural titanium alloys for
components and sections in aerospace technology, 10 (2010), 13–23,
doi:10.1080/09507116.2010.486188
2 A. I. Khorev, Complex alloying and heat treatment of titanium
alloys, 6 (2008), 364–368, doi:10.1080/09507110802288312
3 M. J. Donachie, Titanium, A Technical Guide, 2nd ed., ASM Interna-
tional, 2000
4 A. G. Illarionov, A. A. Popov, S. M. Illarionova, Effect of Micro-
alloying, with Rem Inclusively, on the Structure, Phase Composition
and Properties of ( + )-Titanium Alloy, Metal Science and Heat
Treatment, 57 (2016), 719–725, doi:10.1007/s11041-016-9948-0
5 E. W. Collings, The physical metallurgy of titanium alloys, American
Society for Metals, 1984, 261
6 Rios PR. Overview No. 62: A theory for grain boundary pinning by
particles, Acta Metall., 35 (1987) 12, 2805–2814
7 A. I. Khorev, Theory and practices of microalloying of near- and
+ titanium alloys with REE, zirconium, hafnium and rhenium,
Techologii machinostroeniya, 1 (2015), 5–10
8 H. Wu, Y. Han, X. Chen, Effects of Yttrium on Mechanical Pro-
perties and Microstructures of Ti-Si Eutectic Alloy, Chinese Journal
of Aeronautics, 18 (2005), 171–174, doi:10.1016/S1000-9361(11)
60324-5
9 R. P. Kolli, A. A. Herzing, S. Ankem, Characterization of yttrium-
rich precipitates in a titanium alloy weld, Materials Characterization,
122 (2016), 30–35, doi:10.1016/j.matchar.2016.10.014
10 B. Poorganji, A. Kazahari, T. Narushima, C. Ouchi, T. Furuhara,
Effect of yttrium addition on grain growth of ,  and + titanium
alloys, Journal of Physics: Conference Series, 240 (2010), 12170,
doi:10.1088/1742-6596/240/1/012170
11 W. F. Cui, C. M. Liu, L. Zhou, G. Z. Luo, Characteristics of
microstructures and second-phase particles in Y-bearing Ti-1100
alloy, Materials Science and Engineering: A, 323 (2002), 192–197,
doi:10.1016/S0921-5093(01)01362-4
12 G. Lütjering, J. C. Williams, Titanium, 2nd ed., Springer, Berlin,
2007, 442
13 Y. V. Kaletina, E. D. Efimova, E. G. Gerasimov, A. Y. Kaletin, Effect
of thermal cycling on structure and properties of Ni–Mn–In-based
alloys, Technical Physics, 12 (2016), 1894–1897
14 H. M. Flower, P. R. Swann, D. R. F. West, The effect of Si, Zr, Al
and Mo on the structure and strength of Ti martensite, Journal of
Materials Science, 8 (1972), 929–938
A. G. ILLARIONOV et al.: EFFECT OF YTTRIUM AND ZIRCONIUM MICROALLOYING ON THE STRUCTURE ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 855–859 859
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
861–869
MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS
OF Al/Cu FRICTION-STIR SPOT WELDS
RAZVOJ MIKROSTRUKTURE IN STATISTI^NA ANALIZA
VRTILNO-TORNIH TO^KASTIH ZVAROV Al/Cu
Mukuna Patrick Mubiayi1, Esther Titilayo Akinlabi1,
Mamookho Elizabeth Makhatha2
1University of Johannesburg, Department of Mechanical Engineering Science, Auckland Park Kingsway Campus,
2006 Johannesburg, South Africa
2University of Johannesburg, Department of Metallurgy, School of Mining, Metallurgy and Chemical Engineering, Doornfontein Campus,
2028 Doornfontein, South Africa
patrickmubiayi@gmail.com
Prejem rokopisa – received: 2016-11-18; sprejem za objavo – accepted for publication: 2017-03-16
doi:10.17222/mit.2016.320
In this paper, friction-stir spot welding (FSSW) is performed on 3 mm thick AA1060 and C11000 using different process
parameters and tool geometries. The microstructure- and the microhardness-profile analyses were conducted and the probability
distribution function (PDF) of the obtained microhardness values was determined. Optical images showed a good material
mixing in most of the spot welds produced, whereas the energy-dispersive-spectroscopy (EDS) analysis showed the presence of
intermetallic compounds. Microhardness results revealed that process parameters and tool geometries have significant effects on
the distribution of microhardness values in different locations of the produced spot welds. Furthermore, goodness-of-fit values
showed that most of the R2 values ranged between 0.8842 and 0.9999, which indicated how well the model fits with the
experimental data. On the other hand, the residuals comprised positive and negative runs which also indicated the existence of a
certain correlation with the experimentation.
Keywords: aluminium, copper, friction-stir spot welding, statistical analysis
V prispevku je predstavljeno vrtilno-torno to~kovno varjenje (angl. FSSW) na 3 mm debeli plo~evini AA1060 in C11000 z
uporabo razli~nih procesnih parametrov in razli~no geometrijo orodij. Izvedene so bile analize profila mikrostrukture in
mikrotrdote in dolo~ena je bila porazdelitvena funkcija verjetnosti za dobljene vrednosti mikrotrdote. Posnetki so pokazali
dobro me{anje materiala na ve~ini izdelanih to~kovnih zvarov, medtem ko je EDS-analiza pokazala prisotnost intermetalnih
spojin. Meritve mikrotrdote so pokazale, da imajo procesni parametri in geometrija orodij pomemben vpliv na porazdelitev
mikrotrdote glede na razli~ne lokacije izdelanih to~kovnih zvarov. Nadalje je serija opazovanj oz. ocena pokazala, da je ve~ina
R2 vrednosti rangiranih med 0,8842 in 0,9999. To potrjuje, da se model dobro ujema z eksperimentalnimi podatki. Po drugi
strani pa razlika med pozitivnimi in negativnimi cikli ka`e na obstoj dolo~ene korelacije s preizku{anjem.
Klju~ne besede: aluminij, baker, vrtilno-torno to~kovno varjenje, statisti~na analiza
1 INTRODUCTION
Friction-stir welding (FSW) is a fairly new solid-state
joining technique created and patented by The Welding
Institute (TWI) in 1991 for butt and lap welding of
ferrous and non-ferrous metals and plastics.1 Friction-stir
spot welding (FSSW) is a novel variant of linear
friction-stir welding (FSW) used for spot-welding
applications.2 The FSSW process uses a non-consumable
rotating tool which is plunged into the workpieces to be
welded. Before attaining the selected plunge depth, the
rotating tool is held at the same position for a fixed time,
which is defined as the dwell time. Consequently, the
rotating tool is withdrawn from the welded joint, leaving
a solid-phase friction-stir spot weld behind. Throughout
the FSSW process, the tool penetration depth and the
Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869 861
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.1:67.017:669.71:621.791 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)861(2017)
Figure 1: Schematic diagram of the friction-stir spot-welding process3
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
861–869
MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS
OF Al/Cu FRICTION-STIR SPOT WELDS
RAZVOJ MIKROSTRUKTURE IN STATISTI^NA ANALIZA
VRTILNO-TORNIH TO^KASTIH ZVAROV Al/Cu
Mukuna Patrick Mubiayi1, Esther Titilayo Akinlabi1,
Mamookho Elizabeth Makhatha2
1University of Johannesburg, Department of Mechanical Engineering Science, Auckland Park Kingsway Campus,
2006 Johannesburg, South Africa
2University of Johannesburg, Department of Metallurgy, School of Mining, Metallurgy and Chemical Engineering, Doornfontein Campus,
2028 Doornfontein, South Africa
patrickmubiayi@gmail.com
Prejem rokopisa – received: 2016-11-18; sprejem za objavo – accepted for publication: 2017-03-16
doi:10.17222/mit.2016.320
In this paper, friction-stir spot welding (FSSW) is performed on 3 mm thick AA1060 and C11000 using different process
parameters and tool geometries. The microstructure- and the microhardness-profile analyses were conducted and the probability
distribution function (PDF) of the obtained microhardness values was determined. Optical images showed a good material
mixing in most of the spot welds produced, whereas the energy-dispersive-spectroscopy (EDS) analysis showed the presence of
intermetallic compounds. Microhardness results revealed that process parameters and tool geometries have significant effects on
the distribution of microhardness values in different locations of the produced spot welds. Furthermore, goodness-of-fit values
showed that most of the R2 values ranged between 0.8842 and 0.9999, which indicated how well the model fits with the
experimental data. On the other hand, the residuals comprised positive and negative runs which also indicated the existence of a
certain correlation with the experimentation.
Keywords: aluminium, copper, friction-stir spot welding, statistical analysis
V prispevku je predstavljeno vrtilno-torno to~kovno varjenje (angl. FSSW) na 3 mm debeli plo~evini AA1060 in C11000 z
uporabo razli~nih procesnih parametrov in razli~no geometrijo orodij. Izvedene so bile analize profila mikrostrukture in
mikrotrdote in dolo~ena je bila porazdelitvena funkcija verjetnosti za dobljene vrednosti mikrotrdote. Posnetki so pokazali
dobro me{anje materiala na ve~ini izdelanih to~kovnih zvarov, medtem ko je EDS-analiza pokazala prisotnost intermetalnih
spojin. Meritve mikrotrdote so pokazale, da imajo procesni parametri in geometrija orodij pomemben vpliv na porazdelitev
mikrotrdote glede na razli~ne lokacije izdelanih to~kovnih zvarov. Nadalje je serija opazovanj oz. ocena pokazala, da je ve~ina
R2 vrednosti rangiranih med 0,8842 in 0,9999. To potrjuje, da se model dobro ujema z eksperimentalnimi podatki. Po drugi
strani pa razlika med pozitivnimi in negativnimi cikli ka`e na obstoj dolo~ene korelacije s preizku{anjem.
Klju~ne besede: aluminij, baker, vrtilno-torno to~kovno varjenje, statisti~na analiza
1 INTRODUCTION
Friction-stir welding (FSW) is a fairly new solid-state
joining technique created and patented by The Welding
Institute (TWI) in 1991 for butt and lap welding of
ferrous and non-ferrous metals and plastics.1 Friction-stir
spot welding (FSSW) is a novel variant of linear
friction-stir welding (FSW) used for spot-welding
applications.2 The FSSW process uses a non-consumable
rotating tool which is plunged into the workpieces to be
welded. Before attaining the selected plunge depth, the
rotating tool is held at the same position for a fixed time,
which is defined as the dwell time. Consequently, the
rotating tool is withdrawn from the welded joint, leaving
a solid-phase friction-stir spot weld behind. Throughout
the FSSW process, the tool penetration depth and the
Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869 861
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.1:67.017:669.71:621.791 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)861(2017)
Figure 1: Schematic diagram of the friction-stir spot-welding process3
dwell time fundamentally determine the heat generation,
material plasticization around the tool’s pin, weld
geometry and, hence, the mechanical properties of the
welded joint.2 Figure 1 depicts a schematic illustration
of the FSSW technique.
It should be noted that the FSSW process uses a
non-consumable tool, which is similar to the FSW tool.4
The shoulder produces the frictional heat, whereas the
pin creates the material flow between the work sheets.2,3
In addition to the tool, the other parameters involved in
FSSW include the tool rotation speed, the tool plunge
depth and the tool dwell time. These parameters have an
effect on the strength, the surface texture of the welded
joints2 and the existence of weld defects.3
A typical cross-section of a friction-stir spot weld
displays five different structures including the parent
material (PM), the heat-affected zone (HAZ), the
thermomechanically affected zone (TMAZ), the stir zone
(SZ) and the hook.2
Friction-stir welding (FSW) and friction-stir spot
welding (FSSW) of aluminium and copper have thus far
not been fully investigated due to the huge difference
between their melting temperatures and the high chemi-
cal affinity of both materials which facilitate the forma-
tion of brittle intermetallic Al/Cu phases.6–18 Vickers
hardness measurement is a common technique used to
characterise the hardness of materials and it was reported
that the presence of intermetallics affects the hardness
values of the produced FS welds and FSS welds.11,17 On
the other hand, statistical analyses of Al/Cu friction-stir
welds and spot welds have not been well researched. S.
Akinlabi and E. T. Akinlabi19 conducted statistical ana-
lyses on the data obtained from dissimilar friction-stir
butt welds of aluminium (AA5754) and copper (C11000)
to understand the link between the process parameters
and the properties of the resulting welds. They concluded
that the downward vertical force has a significant effect
on the ultimate tensile strength (UTS) of the produced
welds. A robust relationship between the electrical resis-
tivity and the heat input into the welds was also ob-
served.19
Statistical analyses evaluate the effects of process
parameters on the properties of the produced spot welds
and establish the relationships amongst the process
parameters.
Therefore, in the current study, an effort was made to
further understand the relationships between the process
parameters and the microhardness profiles of FSSW
welds between copper and aluminium using the proba-
bility distribution function (PDF). Furthermore, the
microstructure and the chemical analyses of the pro-
duced spot welds were also studied.
2 MATERIALS AND WELDING PARAMETERS
In this study, AA1060 and C11000 base materials
with dimensions of 3 mm thickness, 600 mm length and
120 mm width were friction-stir spot lap welded. The
chemical compositions of the two parent materials were
determined using a spectrometer. The chemical compo-
sition of the aluminium sheet is as follows: 0.058 %
mass fraction of Si, 0.481 % mass fraction of Fe,
0.011 % mass fraction of Ga, 0.05 % mass fraction of
other elements and the rest is Al. The chemical com-
position of the copper sheet is: 0.137 % mass fraction of
Zn, <0.1 % mass fractions of Pb, 0.02 % mass fraction of
Ni, 0.023 % mass fraction of Al, 0.012 % mass fraction
of Co, 0.077 % mass fraction of B, 0.036 % mass
fraction of Sb, 0.043 % mass fraction of Nb, <0.492 %
mass fraction of other elements and the rest is Cu.
The sheets were friction-stir spot welded in a 30 mm
overlap configuration. The spot welds were produced at
the eNtsa of Nelson Mandela Metropolitan University,
Port Elizabeth, South Africa using an MTS PDS I-Stir.
The tool material used was H13 tool steel hardened to
50–52 HRC with a 4 mm tool pin, 5 mm tool diameter
and 15 mm tool shoulder. The friction-stir spot welds
were produced at rotational speeds of 800 min–1 and
1200 min–1, the tool-shoulder-plunge depths employed
were 0.5 mm and 1 mm at a constant dwell time of 10 s.
The two different tool profiles used in the current study
were flat pin/flat shoulder and conical pin/concave
shoulder tool, designated as FPS and CCS, respectively.
The produced welds were designated as XX_XX_XX
with the first part describing the tool geometry, the
second part indicating the rotational speed and the third
part indicating the shoulder plunge depth.
The weld samples were sectioned using wire elec-
trical discharge machining (WEDM), grinded and
polished, mounted and prepared, using the ASTM
standard metallographic procedure and ASTM Standard
E3-11.20
A solution of FeCl3 (10g) + HCl (6 mL) + ethanol
(C2H5OH) (20 mL) + H2O (80 mL) was used to etch the
copper side of the spot welds, while the aluminium side
was etched with H2O (190 mL) + HNO3 (5 mL) + HCl
(10 mL) + HF (2 mL). The microstructure of the spot
welds was studied using an optical microscope (Olympus
BX51M) equipped with the Stream software. Scanning
electron microscopy combined with energy dispersive
spectroscopy (SEM/EDS) was used to further examine
the microstructure and the chemical analyses, respec-
tively. A TESCAN equipped with Oxford Instruments
X-Max was used for the SEM/EDS analyses. The
Vickers microhardness was determined using a dia-
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
862 Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Representation of a spot weld with dashed lines illustrating
the location of the microhardness-profile measurements
mond-pyramid-indenter EMCO Test DuraScan tester.
Two different locations on the spot welds were used,
viz., the top and the bottom, as illustrated in Figure 2.
The microhardness measurements were carried out from
the keyhole for all the different parameters and tool
geometries, to find the probability density function
(PDF) of each one. The Matlab 2014 software program
was used to determine the PDF.
3 RESULTS AND DISCUSSION
The microstructure of the friction-stir spot welds was
studied using an optical microscope and the results are
depicted in Figures 3 and 4.
Figure 3 illustrates the microstructures of the spot
welds produced at: a) 800 min–1 and b) 1200 min–1, a
1 mm shoulder plunge depth using a flat pin and a flat
shoulder tool. On the other hand, Figure 4 shows the
microstructure of the spot weld produced at: a) 800 min–1
and b) 1200 min–1, a 0.5 mm shoulder plunge depth
using a conical pin and a concave shoulder tool.
It can be seen in Figure 3a that there are a copper
ring and a mixture of Al/Cu particles in the stir zone.
There is no palpable welding defect in the weld and
copper is disseminated in this zone with different shapes.
In the upper part of the joint, a large bulk of copper with
irregular shapes can be observed (Figure 3a). The tool
pin was inserted in the aluminium plate and the copper
ring extruded upward from the lower copper plate into
the aluminium plate was observed. This was in agree-
ment with the previous work.15
Additionally, the intermixing of copper and alumi-
nium was not homogenous for different spot welds and
different microstructures were formed in different
regions of the welds. It was reported that the FSW of
dissimilar materials is different from that of similar
materials due to the formation of a complex, intercalated
vortex-like and related flow pattern.21
In Figure 4b, a good interlaced structure can be seen.
This is formed by aluminium and copper, thereby
indicating that the two plates are bonded firmly in this
region, which is composed of a lamellar structure of
copper particles with a streamlined shape of aluminium
strips. In this region, a few disseminated copper particles
were also observed.
The energy-dispersive-spectrometry analyses at the
selected points in the stir zone were recorded using SEM
micrographs (Table 1). Intermetallic compounds were
found in most of the produced welds. Two intermetallic
compounds, viz., Al2Cu and Al3Cu4 were found in the
weld produced at 1200 min–1, 0.5 mm shoulder plunge
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869 863
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Optical microscope images showing the macrostructures of
the joints at: a) 800 min–1 and b) 1200 min–1, 0.5 mm shoulder plunge
depth using a conical pin and concave shoulder tool
Figure 3: Optical microscope images showing the macrostructures of
the joints at: a) 800 min–1 and b) 1200 min–1, 1 mm shoulder plunge
depth using a flat pin and flat shoulder tool
dwell time fundamentally determine the heat generation,
material plasticization around the tool’s pin, weld
geometry and, hence, the mechanical properties of the
welded joint.2 Figure 1 depicts a schematic illustration
of the FSSW technique.
It should be noted that the FSSW process uses a
non-consumable tool, which is similar to the FSW tool.4
The shoulder produces the frictional heat, whereas the
pin creates the material flow between the work sheets.2,3
In addition to the tool, the other parameters involved in
FSSW include the tool rotation speed, the tool plunge
depth and the tool dwell time. These parameters have an
effect on the strength, the surface texture of the welded
joints2 and the existence of weld defects.3
A typical cross-section of a friction-stir spot weld
displays five different structures including the parent
material (PM), the heat-affected zone (HAZ), the
thermomechanically affected zone (TMAZ), the stir zone
(SZ) and the hook.2
Friction-stir welding (FSW) and friction-stir spot
welding (FSSW) of aluminium and copper have thus far
not been fully investigated due to the huge difference
between their melting temperatures and the high chemi-
cal affinity of both materials which facilitate the forma-
tion of brittle intermetallic Al/Cu phases.6–18 Vickers
hardness measurement is a common technique used to
characterise the hardness of materials and it was reported
that the presence of intermetallics affects the hardness
values of the produced FS welds and FSS welds.11,17 On
the other hand, statistical analyses of Al/Cu friction-stir
welds and spot welds have not been well researched. S.
Akinlabi and E. T. Akinlabi19 conducted statistical ana-
lyses on the data obtained from dissimilar friction-stir
butt welds of aluminium (AA5754) and copper (C11000)
to understand the link between the process parameters
and the properties of the resulting welds. They concluded
that the downward vertical force has a significant effect
on the ultimate tensile strength (UTS) of the produced
welds. A robust relationship between the electrical resis-
tivity and the heat input into the welds was also ob-
served.19
Statistical analyses evaluate the effects of process
parameters on the properties of the produced spot welds
and establish the relationships amongst the process
parameters.
Therefore, in the current study, an effort was made to
further understand the relationships between the process
parameters and the microhardness profiles of FSSW
welds between copper and aluminium using the proba-
bility distribution function (PDF). Furthermore, the
microstructure and the chemical analyses of the pro-
duced spot welds were also studied.
2 MATERIALS AND WELDING PARAMETERS
In this study, AA1060 and C11000 base materials
with dimensions of 3 mm thickness, 600 mm length and
120 mm width were friction-stir spot lap welded. The
chemical compositions of the two parent materials were
determined using a spectrometer. The chemical compo-
sition of the aluminium sheet is as follows: 0.058 %
mass fraction of Si, 0.481 % mass fraction of Fe,
0.011 % mass fraction of Ga, 0.05 % mass fraction of
other elements and the rest is Al. The chemical com-
position of the copper sheet is: 0.137 % mass fraction of
Zn, <0.1 % mass fractions of Pb, 0.02 % mass fraction of
Ni, 0.023 % mass fraction of Al, 0.012 % mass fraction
of Co, 0.077 % mass fraction of B, 0.036 % mass
fraction of Sb, 0.043 % mass fraction of Nb, <0.492 %
mass fraction of other elements and the rest is Cu.
The sheets were friction-stir spot welded in a 30 mm
overlap configuration. The spot welds were produced at
the eNtsa of Nelson Mandela Metropolitan University,
Port Elizabeth, South Africa using an MTS PDS I-Stir.
The tool material used was H13 tool steel hardened to
50–52 HRC with a 4 mm tool pin, 5 mm tool diameter
and 15 mm tool shoulder. The friction-stir spot welds
were produced at rotational speeds of 800 min–1 and
1200 min–1, the tool-shoulder-plunge depths employed
were 0.5 mm and 1 mm at a constant dwell time of 10 s.
The two different tool profiles used in the current study
were flat pin/flat shoulder and conical pin/concave
shoulder tool, designated as FPS and CCS, respectively.
The produced welds were designated as XX_XX_XX
with the first part describing the tool geometry, the
second part indicating the rotational speed and the third
part indicating the shoulder plunge depth.
The weld samples were sectioned using wire elec-
trical discharge machining (WEDM), grinded and
polished, mounted and prepared, using the ASTM
standard metallographic procedure and ASTM Standard
E3-11.20
A solution of FeCl3 (10g) + HCl (6 mL) + ethanol
(C2H5OH) (20 mL) + H2O (80 mL) was used to etch the
copper side of the spot welds, while the aluminium side
was etched with H2O (190 mL) + HNO3 (5 mL) + HCl
(10 mL) + HF (2 mL). The microstructure of the spot
welds was studied using an optical microscope (Olympus
BX51M) equipped with the Stream software. Scanning
electron microscopy combined with energy dispersive
spectroscopy (SEM/EDS) was used to further examine
the microstructure and the chemical analyses, respec-
tively. A TESCAN equipped with Oxford Instruments
X-Max was used for the SEM/EDS analyses. The
Vickers microhardness was determined using a dia-
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
862 Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Representation of a spot weld with dashed lines illustrating
the location of the microhardness-profile measurements
mond-pyramid-indenter EMCO Test DuraScan tester.
Two different locations on the spot welds were used,
viz., the top and the bottom, as illustrated in Figure 2.
The microhardness measurements were carried out from
the keyhole for all the different parameters and tool
geometries, to find the probability density function
(PDF) of each one. The Matlab 2014 software program
was used to determine the PDF.
3 RESULTS AND DISCUSSION
The microstructure of the friction-stir spot welds was
studied using an optical microscope and the results are
depicted in Figures 3 and 4.
Figure 3 illustrates the microstructures of the spot
welds produced at: a) 800 min–1 and b) 1200 min–1, a
1 mm shoulder plunge depth using a flat pin and a flat
shoulder tool. On the other hand, Figure 4 shows the
microstructure of the spot weld produced at: a) 800 min–1
and b) 1200 min–1, a 0.5 mm shoulder plunge depth
using a conical pin and a concave shoulder tool.
It can be seen in Figure 3a that there are a copper
ring and a mixture of Al/Cu particles in the stir zone.
There is no palpable welding defect in the weld and
copper is disseminated in this zone with different shapes.
In the upper part of the joint, a large bulk of copper with
irregular shapes can be observed (Figure 3a). The tool
pin was inserted in the aluminium plate and the copper
ring extruded upward from the lower copper plate into
the aluminium plate was observed. This was in agree-
ment with the previous work.15
Additionally, the intermixing of copper and alumi-
nium was not homogenous for different spot welds and
different microstructures were formed in different
regions of the welds. It was reported that the FSW of
dissimilar materials is different from that of similar
materials due to the formation of a complex, intercalated
vortex-like and related flow pattern.21
In Figure 4b, a good interlaced structure can be seen.
This is formed by aluminium and copper, thereby
indicating that the two plates are bonded firmly in this
region, which is composed of a lamellar structure of
copper particles with a streamlined shape of aluminium
strips. In this region, a few disseminated copper particles
were also observed.
The energy-dispersive-spectrometry analyses at the
selected points in the stir zone were recorded using SEM
micrographs (Table 1). Intermetallic compounds were
found in most of the produced welds. Two intermetallic
compounds, viz., Al2Cu and Al3Cu4 were found in the
weld produced at 1200 min–1, 0.5 mm shoulder plunge
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869 863
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 4: Optical microscope images showing the macrostructures of
the joints at: a) 800 min–1 and b) 1200 min–1, 0.5 mm shoulder plunge
depth using a conical pin and concave shoulder tool
Figure 3: Optical microscope images showing the macrostructures of
the joints at: a) 800 min–1 and b) 1200 min–1, 1 mm shoulder plunge
depth using a flat pin and flat shoulder tool
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
864 Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Table 1 displays SEM micrographs and the EDS analysis of selected points on the produced friction-spot welds: a) (FFS_1200_0.5), b)
(FFS_1200_1), c) (CCS_1200_0.5) and d) (CCS_1200_1).
Point
Composition Intermetallic
compoundAl Cu
1
2
3
4
5
54.94
48.5
4.34
24.55
41.98
45.06
51.5
95.66
75.45
58.02
-
Al2Cu
-
Al3Cu4
Al2Cu
Point
Composition Intermetallic
compoundAl Cu
1
2
3
4
5
1.46
17.71
36.27
94.53
16.98
98.54
82.29
63.73
5.07
83.02
-
Al4Cu9
AlCu
-
Al4Cu9
Point
Composition Intermetallic
compoundAl Cu
1
2
3
4
5
3.1
7.82
49.49
29.22
13.77
96.9
92.18
50.27
70.78
86.23
-
-
Al2Cu
AlCu
AlCu3
Point
Composition Intermetallic
compoundAl Cu
1
2
3
4
5
2.26
1.33
92.77
37.35
9.15
97.74
98.67
6.97
62.65
90.85
-
-
-
AlCu
-
depth, whereas the weld produced at 1200 min–1 and 1
mm shoulder plunge depth contained Al4Cu9 and AlCu
intermetallic compounds. These intermetallics were
found in the welds produced using a flat pin and flat
shoulder tool. On the other hand, the spot welds pro-
duced using a conical pin and concave shoulder con-
tained Al2Cu, AlCu3 and AlCu intermetallic compounds.
These intermetallic compounds were found in the weld
produced at 1200 min–1 and 0.5 mm shoulder plunge
depth. Only the AlCu intermetallic was found in the
weld produced at 1200 min–1 and 1 mm shoulder plunge
depth, but the concentration of this intermetallic com-
pound was relatively small. It was reported that the
presence of intermetallic compounds could affect the
microhardness profile.17
Figure 5 (a, b, c and d) shows the microhardness
values of the spot welds produced using a flat pin and a
flat shoulder tool, or a conical pin and a concave
shoulder at different process parameters. The micro-
hardness values of the parent materials were in the range
of 86.7–96.3 HV for Cu while for Al, the range was
between 34.6–40.3 HV. In all the samples, high micro-
hardness values were recorded at the top, in the region
close to the keyhole.
It was reported that all the mechanical tests are
subject to large statistical variations, which should be
evaluated.22
The probability distribution function (PDF) of the
Vickers hardness was reported in the literature to
correspond to the Gaussian (or normal) distribution23 and
log-normal distribution.24 A. M. Hassan et al.25 studied
the significance of the process parameters of friction-stir
welding of aluminium-matrix composites to set the
optimum level for each of these parameters and to
further predict which responses are affected when using
analyses of variance (ANOVA).25 The present study used
the Matlab 2014a statistical toolbox to analyse the pro-
bability density function (PDF) of the obtained micro-
hardness results. This was done to understand how
different parameters and tool geometries affect the
probability of obtaining specific microhardness values.
A probability density function (PDF) is a function
that defines the relative possibility for a random variable
to take on a given value. The probability of the random
variable falling within a particular range [a, b] of values
is given with a finite integral of the PDF within that
range [a, b], Equation (1):
p x f x e
b
a x
b
a
( ) ( , , )
( )
= =∫ ∫
−
−μ
 
 

 

1
2
2
2
π
(1)
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869 865
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Microhardness distributions along the welds produced using different tools and process parameters: a) flat pin/flat shoulder (FPS), top;
b) bottom; c) conical pin/concave shoulder (CCS), top; d) bottom
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
864 Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Table 1 displays SEM micrographs and the EDS analysis of selected points on the produced friction-spot welds: a) (FFS_1200_0.5), b)
(FFS_1200_1), c) (CCS_1200_0.5) and d) (CCS_1200_1).
Point
Composition Intermetallic
compoundAl Cu
1
2
3
4
5
54.94
48.5
4.34
24.55
41.98
45.06
51.5
95.66
75.45
58.02
-
Al2Cu
-
Al3Cu4
Al2Cu
Point
Composition Intermetallic
compoundAl Cu
1
2
3
4
5
1.46
17.71
36.27
94.53
16.98
98.54
82.29
63.73
5.07
83.02
-
Al4Cu9
AlCu
-
Al4Cu9
Point
Composition Intermetallic
compoundAl Cu
1
2
3
4
5
3.1
7.82
49.49
29.22
13.77
96.9
92.18
50.27
70.78
86.23
-
-
Al2Cu
AlCu
AlCu3
Point
Composition Intermetallic
compoundAl Cu
1
2
3
4
5
2.26
1.33
92.77
37.35
9.15
97.74
98.67
6.97
62.65
90.85
-
-
-
AlCu
-
depth, whereas the weld produced at 1200 min–1 and 1
mm shoulder plunge depth contained Al4Cu9 and AlCu
intermetallic compounds. These intermetallics were
found in the welds produced using a flat pin and flat
shoulder tool. On the other hand, the spot welds pro-
duced using a conical pin and concave shoulder con-
tained Al2Cu, AlCu3 and AlCu intermetallic compounds.
These intermetallic compounds were found in the weld
produced at 1200 min–1 and 0.5 mm shoulder plunge
depth. Only the AlCu intermetallic was found in the
weld produced at 1200 min–1 and 1 mm shoulder plunge
depth, but the concentration of this intermetallic com-
pound was relatively small. It was reported that the
presence of intermetallic compounds could affect the
microhardness profile.17
Figure 5 (a, b, c and d) shows the microhardness
values of the spot welds produced using a flat pin and a
flat shoulder tool, or a conical pin and a concave
shoulder at different process parameters. The micro-
hardness values of the parent materials were in the range
of 86.7–96.3 HV for Cu while for Al, the range was
between 34.6–40.3 HV. In all the samples, high micro-
hardness values were recorded at the top, in the region
close to the keyhole.
It was reported that all the mechanical tests are
subject to large statistical variations, which should be
evaluated.22
The probability distribution function (PDF) of the
Vickers hardness was reported in the literature to
correspond to the Gaussian (or normal) distribution23 and
log-normal distribution.24 A. M. Hassan et al.25 studied
the significance of the process parameters of friction-stir
welding of aluminium-matrix composites to set the
optimum level for each of these parameters and to
further predict which responses are affected when using
analyses of variance (ANOVA).25 The present study used
the Matlab 2014a statistical toolbox to analyse the pro-
bability density function (PDF) of the obtained micro-
hardness results. This was done to understand how
different parameters and tool geometries affect the
probability of obtaining specific microhardness values.
A probability density function (PDF) is a function
that defines the relative possibility for a random variable
to take on a given value. The probability of the random
variable falling within a particular range [a, b] of values
is given with a finite integral of the PDF within that
range [a, b], Equation (1):
p x f x e
b
a x
b
a
( ) ( , , )
( )
= =∫ ∫
−
−μ
 
 

 

1
2
2
2
π
(1)
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869 865
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 5: Microhardness distributions along the welds produced using different tools and process parameters: a) flat pin/flat shoulder (FPS), top;
b) bottom; c) conical pin/concave shoulder (CCS), top; d) bottom
where  is the standard deviation, 2 is the variance and
μ is the mean.
It is given by the area under the density function,
nonetheless above the horizontal axis and in between the
lowest and highest values of the range. The probability
density function is a non-negative value and its integral
over the entire space is equal to one.
PDF histograms of the microhardness and their fits
for the parent materials, namely, aluminium and copper,
are depicted in Figures 6a and 7b, respectively. The
PDFs of the top and bottom hardness measurements
were investigated, Figures 7a and 7b depict the PDF
histograms of the microhardness for the weld produced
at 800 min–1, 0.5 mm shoulder plunge depth, using a flat
pin and a flat shoulder tool.
It can be seen that the probability of having the
microhardness values between 40 and 45 HV is high in
the histogram, as shown in Figure 7a, which represents
the microhardness measured at the top of the spot weld.
This corresponds mostly to the microhardness of the
aluminium parent material, whereas the possibility of
getting high microhardness values between 50 and 60
HV is low.
On the other hand, the PDF of the bottom measu-
rement (Figure 7b) shows that there is a high possibility
of getting microhardness values between 85 and 95 HV.
This corresponds to the microhardness of the copper-
parent sheets, while the microhardness values between
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
866 Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Depicting the microhardness PDF histograms of the parent
materials: a) aluminium, b) copper
Figure 8: PDF histogram of the microhardness of CCS_800_0.5 spot
weld: a) top and b) bottom
Figure 7: PDF histograms of the microhardness of FPS_800_0.5 spot
weld: a) top and b) bottom
100 HV and 105 HV are likely to show a lower possi-
bility of being obtained when using the same process
parameters as those used in this research work. The
possibility of having higher microhardness values
compared to the values of the parent materials in the two
different sheets (copper and aluminium) was observed to
be due to the presence of a mixture of copper and
aluminium in the vicinity of the keyhole. Additionally,
Figure 8 depicts a PDF histogram of the microhardness
measurements ((a) top and (b) bottom) for the weld
produced at 800 min–1, 0.5 mm shoulder plunge depth,
using a conical pin and a concave shoulder.
The results show that there is a higher possibility for
obtaining microhardness values between 30 50 HV and
50 HV, whereas the possibility for high values between
120 and 130 HV is low (Figure 8a). The trend for the
bottom area (Figure 8b) is similar to the one discussed
above for the PDF of the bottom microhardness obtained
using a flat pin and a flat shoulder.
Moreover, when the rotational speed of 1200 min–1 is
increased, the possibility of getting high microhardness
values ranging between 100–110 HV and 90–100 HV
increases for the top area of the spot weld produced
using a conical pin and a concave shoulder, and for the
top and bottom microhardness values, respectively. It can
be seen that the rotational speed and the tool geometry
may influence the possibility of different probability
distributions.
The model shows that, in order to get the probability
in a specific region, the integral of the PDF for the region
of interest need to be computed. The PDF found in the
current research work is a normal distribution (called a
Gaussian distribution as well). In order to get any pro-
bability, we can compute the finite integral of the normal
distribution equation (1).
The goodness of fit and the residuals were also
analysed. The results show that most of the R2 values
range between 0.8842 and 0.9999, which is an indication
of how well the model fits with the experimental data
shown in Table 2.
Table 2: R2 and adjusted R2 of the welds produced using a flat pin/flat
shoulder tool and conical pin/concave shoulder tool for the micro-
hardness measured on the top and at the bottom
Sample ID R2 square R2 square
FPS_800_0.5 0.9896 0.8842
FPS_800_1 0.9996 0.9999
FPS_1200_0.5 0.9999 0.9937
FPS_1200_1 0.9993 0.9298
CCS_800_0.5 0.9924 0.9818
CCS_800_1 0.9976 0.9424
CCS_1200_0.5 0.9893 0.987
CCS_1200_1 0.9997 0.9866
Figures 9a to 9b depict the goodness of fit and the
residuals for the weld produced at 1200 min–1, 1 mm
shoulder plunge depth, using a flat pin and a flat
shoulder (the top-microhardness measurement). In
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869 867
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 9: a) goodness of fit and b) the residuals for the spot weld
produced at 1200 min–1, 1 mm shoulder plunge depth, using a flat pin
and a flat shoulder (top-microhardness measurements)
Figure 10: a) goodness of fit and b) the residuals for the spot weld
produced at 1200 min–1, 1 mm shoulder plunge depth, using a conical
pin and a concave shoulder (top-microhardness measurements)
where  is the standard deviation, 2 is the variance and
μ is the mean.
It is given by the area under the density function,
nonetheless above the horizontal axis and in between the
lowest and highest values of the range. The probability
density function is a non-negative value and its integral
over the entire space is equal to one.
PDF histograms of the microhardness and their fits
for the parent materials, namely, aluminium and copper,
are depicted in Figures 6a and 7b, respectively. The
PDFs of the top and bottom hardness measurements
were investigated, Figures 7a and 7b depict the PDF
histograms of the microhardness for the weld produced
at 800 min–1, 0.5 mm shoulder plunge depth, using a flat
pin and a flat shoulder tool.
It can be seen that the probability of having the
microhardness values between 40 and 45 HV is high in
the histogram, as shown in Figure 7a, which represents
the microhardness measured at the top of the spot weld.
This corresponds mostly to the microhardness of the
aluminium parent material, whereas the possibility of
getting high microhardness values between 50 and 60
HV is low.
On the other hand, the PDF of the bottom measu-
rement (Figure 7b) shows that there is a high possibility
of getting microhardness values between 85 and 95 HV.
This corresponds to the microhardness of the copper-
parent sheets, while the microhardness values between
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
866 Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Depicting the microhardness PDF histograms of the parent
materials: a) aluminium, b) copper
Figure 8: PDF histogram of the microhardness of CCS_800_0.5 spot
weld: a) top and b) bottom
Figure 7: PDF histograms of the microhardness of FPS_800_0.5 spot
weld: a) top and b) bottom
100 HV and 105 HV are likely to show a lower possi-
bility of being obtained when using the same process
parameters as those used in this research work. The
possibility of having higher microhardness values
compared to the values of the parent materials in the two
different sheets (copper and aluminium) was observed to
be due to the presence of a mixture of copper and
aluminium in the vicinity of the keyhole. Additionally,
Figure 8 depicts a PDF histogram of the microhardness
measurements ((a) top and (b) bottom) for the weld
produced at 800 min–1, 0.5 mm shoulder plunge depth,
using a conical pin and a concave shoulder.
The results show that there is a higher possibility for
obtaining microhardness values between 30 50 HV and
50 HV, whereas the possibility for high values between
120 and 130 HV is low (Figure 8a). The trend for the
bottom area (Figure 8b) is similar to the one discussed
above for the PDF of the bottom microhardness obtained
using a flat pin and a flat shoulder.
Moreover, when the rotational speed of 1200 min–1 is
increased, the possibility of getting high microhardness
values ranging between 100–110 HV and 90–100 HV
increases for the top area of the spot weld produced
using a conical pin and a concave shoulder, and for the
top and bottom microhardness values, respectively. It can
be seen that the rotational speed and the tool geometry
may influence the possibility of different probability
distributions.
The model shows that, in order to get the probability
in a specific region, the integral of the PDF for the region
of interest need to be computed. The PDF found in the
current research work is a normal distribution (called a
Gaussian distribution as well). In order to get any pro-
bability, we can compute the finite integral of the normal
distribution equation (1).
The goodness of fit and the residuals were also
analysed. The results show that most of the R2 values
range between 0.8842 and 0.9999, which is an indication
of how well the model fits with the experimental data
shown in Table 2.
Table 2: R2 and adjusted R2 of the welds produced using a flat pin/flat
shoulder tool and conical pin/concave shoulder tool for the micro-
hardness measured on the top and at the bottom
Sample ID R2 square R2 square
FPS_800_0.5 0.9896 0.8842
FPS_800_1 0.9996 0.9999
FPS_1200_0.5 0.9999 0.9937
FPS_1200_1 0.9993 0.9298
CCS_800_0.5 0.9924 0.9818
CCS_800_1 0.9976 0.9424
CCS_1200_0.5 0.9893 0.987
CCS_1200_1 0.9997 0.9866
Figures 9a to 9b depict the goodness of fit and the
residuals for the weld produced at 1200 min–1, 1 mm
shoulder plunge depth, using a flat pin and a flat
shoulder (the top-microhardness measurement). In
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869 867
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 9: a) goodness of fit and b) the residuals for the spot weld
produced at 1200 min–1, 1 mm shoulder plunge depth, using a flat pin
and a flat shoulder (top-microhardness measurements)
Figure 10: a) goodness of fit and b) the residuals for the spot weld
produced at 1200 min–1, 1 mm shoulder plunge depth, using a conical
pin and a concave shoulder (top-microhardness measurements)
addition, Figures 10a to 10b depict the goodness of fit
and the residuals for the weld produced at 1200 min–1, 1
mm shoulder plunge depth, using a conical pin and a
concave shoulder (the top-microhardness measurement).
Besides the R2 values, the residual analysis was also
employed in the study in order to check the adequacy of
the models. Figures 9b and 10b show the residual plots
for the microhardness values obtained at the top, using a
flat pin/flat shoulder and a conical pin/concave shoulder,
using the 1200 min–1 speed and 1 mm shoulder plunge
depth, respectively. It was reported that the tendencies to
have runs of positive and negative residuals indicate the
existence of a certain correlation with the experimen-
tation.26
Tables 3 and 4 present the standard deviation, the
variance and the mean obtained from the statistical
analyses of the measured microhardness values for the
two different positions, namely, the top and bottom of
the spot welds produced with different tools and using
different process parameters.
Table 3: Mean, variance, mu and sigma of the spot samples produced
using a flat pin/flat shoulder tool and a conical pin/concave shoulder
tool for the microhardness taken on the top
Sample ID Mean 2 μ 
FPS_800_0.5 46.543 82.8442 46.5429 9.102
FPS_800_1 75.099 4994.97 4.00167 0.796
FPS_1200_0.5 45.121 441.806 3.71115 0.443
FPS_1200_1 51.731 672.845 3.83391 0.474
CCS_800_0.5 44.575 409.802 3.70341 0.433
CCS_800_1 55.887 813.763 3.90757 0.481
CCS_1200_0.5 45.453 482.889 3.71166 0.458
CCS_1200_1 40.942 248.04 3.64315 0.371
Table 4: Mean, variance, mu and sigma of the spot samples produced
using a flat pin/flat shoulder tool and a conical pin/concave shoulder
tool for the microhardness taken at the bottom
Sample ID Mean 2 μ 
FPS_800_0.5 92.234 32.7196 4.52241 0.062
FPS_800_1 84.7 119.974 84.7 10.95
FPS_1200_0.5 84.7 119.974 84.7 10.95
FPS_1200_1 78.429 58.0637 78.4286 7.62
CCS_800_0.5 89.986 14.2859 89.9857 3.78
CCS_800_1 88.564 102.307 88.5643 10.11
CCS_1200_0.5 86.252 89.6604 4.45128 0.109
CCS_1200_1 82.043 54.8519 4.40318 0.09
It should be noted that the mean is equal to μ if the
distribution is normal. It can be seen that in some cases
in the current work μ and the mean have different values,
which shows that the distributions in some of the
analyses were not normal. The standard deviation 
should be close to zero, but in the current work, the value
of the standard deviation is not close to zero in some of
the cases. This shows that the microhardness values are
not close to the expected values and this could be due to
the microhardness values measured in different locations
of the weld samples, far apart from each other.
This was further suspected due to the presence of
intermetallics, which could have been the cause of the
high microhardness values since intermetallics are
invariably hard and brittle.
Each figure contains residuals versus the distance
(the distance from the keyhole), taking into account the
data and the constant variance of the residuals. In the
plot of residuals versus distance, it is shown that the
models are adequate to predict the responses in an
acceptable manner.
4 CONCLUSIONS
According to the presented results, some conclusions
can be drawn:
• The microstructure of the produced spot welds shows
a good material mixing, the presence of a copper ring
and a mixture of Al/Cu particles present in the stir
zone.
• The EDS analyses of the produced friction-stir spot
welds exhibited the presence of intermetallic com-
pounds, which are known to affect the microhard-
ness.
• The microhardness values obtained at the top were
high for all the samples and this was in the region
close to the spot-weld keyhole. In addition, all the
microhardness values obtained at the bottom of the
samples, in the region close to the keyhole, have
lower values, which were close to the average value
of the copper base material. This occurred for all the
spot welds produced using a conical pin and concave
shoulder.
• The probability-density-function (PDF) histograms
of the microhardness results revealed that the process
parameters and the tool geometries have significant
effects on the distribution of the microhardness
values in different locations of the produced spot
welds.
• Additionally, goodness-of-fit values were also
analysed and these showed that most of the R2 values
ranged between 0.8842 and 0.9999, which is an
indication of how well the model fits with the
produced experimental data.
Acknowledgements
The authors wish to acknowledge the financial
support of the University of Johannesburg and the
assistance from Mr. Riaan Brown (eNtsa at Nelson
Mandela Metropolitan University).
5 REFERENCES
1 W. M. Thomas, E. D. Nicholas, J. C. Needham, M. G. Murch, P.
Temple-Smith, C. J. Dawes, International Patent No. PCT/GB92/
02203, GB patent application No. 9125978.8, 1991
2 H. Badarinarayan, Fundamentals of Friction Stir Spot Welding, PhD
Thesis, 2009, Missouri
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
868 Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
3 W. Yuan, R. S. Mishra, B. Carlson, R. Verma, R. K. Mishra, Material
flow and microstructural evolution during friction stir spot welding
of AZ31 magnesium alloy, Materials Science and Engineering, A
543 (2012), 200–209
4 J. Podr`aj, B. Jerman, D. Klob~ar, Welding defects at friction stir
welding, Metalurgija, 54 (2015) 2, 387–389
5 J. M. Timothy, Friction Stir Welding of Commercially Available
Superplastic Aluminium, PhD thesis, 2008, Department of Engi-
neering and Design, Brunel University, Brunel
6 J. Ouyang, E. Yarrapareddy, R. Kovacevic, Microstructural evolution
in the friction stir welded 6061 aluminum alloy (T6-temper condi-
tion) to copper, Journal of Materials Processing Technology, 172
(2006), 110–122
7 P. Liu, Q. Shi, W. Wang, X. Wang, Z. Zhang, Microstructure and
XRD analysis of FSW joints for copper T2/aluminium 5A06
dissimilar materials, Materials Letters, 62 (2008), 4106–4108
8 P. Xue, B. L. Xiao, D. R. Ni, Z. Y. Ma, Enhanced mechanical pro-
perties of friction stir welded dissimilar Al–Cu joint by intermetallic
compounds, Materials Science and Engineering A, 527 (2010),
5723–5727
9 A. Esmaeili, M. K. Besharati Givi, H. R. Zareie Rajani, A metallur-
gical and mechanical study on dissimilar friction stir welding of
aluminum 1050 to brass (CuZn30), Materials Science and Engi-
neering, A 528 (2011), 7093–7102
10 A. Abdollah-Zadeh, T. Saeid, B. Sazgari, Microstructural and mecha-
nical properties of friction stir welded aluminum/copper lap joints,
Journal of Alloys and Compounds, 460 (2008), 535–538
11 E. T. Akinlabi, Characterisation of Dissimilar Friction Stir Welds
between 5754 Aluminium Alloy and C11000 Copper, D-Tech thesis,
Nelson Mandela Metropolitan University, South Africa, 2010
12 M. N. Avettand-Fenoel, R. Taillard, G. Ji, D. Goran, Multiscale
Study of Interfacial Intermetallic Compounds in a Dissimilar Al
6082-T6/Cu Friction-Stir Weld, Metallurgical and Materials
Transactions A, (2012), 4655–4666
13 M. F. X Muthu, V. Jayabalan, Tool Travel Speed Effects on the
Microstructure of Friction Stir Welded Aluminium – Copper Joints,
Journal of Materials Processing Technology, 217 (2015), 105–113
14 M. P. Mubiayi, E. T Akinlabi, Friction Stir Spot Welding between
Copper and Aluminium: Microstructural Evolution, Proceedings of
the International MultiConference of Engineers and Computer
Scientists, Vol II, IMECS 2015, March 18–20, 2015, Hong Kong
15 R. Heideman, C. Johnson, S. Kou, Metallurgical analysis of Al/Cu
friction stir spot welding, Science and Technology of Welding and
Joining, 15 (2010) 7, 597–604
16 U. Özdemir, S. Sayer, Ç. Yeni, Bornova-Izmir, Effect of Pin Pene-
tration Depth on the Mechanical Properties of Friction Stir Spot
Welded Aluminum and Copper, Materials Testing in Joining Techno-
logy, 54 (2012) 4, 233–239
17 M. Shiraly, M. Shamanian, M. R. Toroghinejad, M. Ahmadi Jazani,
Effect of Tool Rotation Rate on Microstructure and Mechanical
Behavior of Friction Stir Spot-Welded Al/Cu Composite, Journal of
Materials Engineering and Performance, 23 (2014) 2, 413–420
18 M. P. Mubiayi, E. T. Akinlabi, Evolving properties of friction stir
spot welds between AA1060 and commercially pure copper C11000,
Transactions of Nonferrous Metals Society of China, 26 (2016),
1852–1862
19 E. T Akinlabi, S. A. Akinlabi, Friction stir welding of dissimilar
materials – statistical analysis of the weld data, Proceedings of the
International MultiConference of Engineers and Computer Scientists,
2012, 1368–1373
20 ASTM Standard E3-11, Standard guide for preparation of
metallographic specimens, ASTM International, West Consho-
hocken, PA, 2011, doi: 10.1520/E0003-11, www.astm.org
21 L. E Murr, A Review of FSW Research on Dissimilar Metal and
Alloy Systems, Journal of Materials Engineering and Performance, 8
(2010)19, 1071–108
22 J. M. Schneider, M. Bigerelle, A. Iost, Statistical analysis of the
Vickers hardness, Materials Science and Engineering, A 262 (1999),
256–263
23 A. L. Yurkov, N. V. Jhuravleva, E. S. Lukin, Kinetic microhardness
measurements of sialon-based ceramics, Journal of Materials
Science, 29 (1994), 6551–6560
24 I. Y. Yanchev, E. P. Trifonova, Analysis of microhardness data in Tlx
ln l-x Se, Journal of Materials Science, 30 (1995), 5576–5580
25 A. M. Hassan, M. Almomani, T. Qasim, A. Ghaithan, Statistical
analysis of some mechanical properties of friction stir welded
aluminium matrix composite, Int. J. Experimental Design and
Process Optimisation, 3 (2012)1, 91–109
26 K. Palanikumar, R. Karthikeyan, Optimal machining conditions for
turning of particulate metal matrix composites using Taguchi and
response surface methodologies, Machining Science and Technology,
10 (2006) 4, 417–433
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869 869
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
addition, Figures 10a to 10b depict the goodness of fit
and the residuals for the weld produced at 1200 min–1, 1
mm shoulder plunge depth, using a conical pin and a
concave shoulder (the top-microhardness measurement).
Besides the R2 values, the residual analysis was also
employed in the study in order to check the adequacy of
the models. Figures 9b and 10b show the residual plots
for the microhardness values obtained at the top, using a
flat pin/flat shoulder and a conical pin/concave shoulder,
using the 1200 min–1 speed and 1 mm shoulder plunge
depth, respectively. It was reported that the tendencies to
have runs of positive and negative residuals indicate the
existence of a certain correlation with the experimen-
tation.26
Tables 3 and 4 present the standard deviation, the
variance and the mean obtained from the statistical
analyses of the measured microhardness values for the
two different positions, namely, the top and bottom of
the spot welds produced with different tools and using
different process parameters.
Table 3: Mean, variance, mu and sigma of the spot samples produced
using a flat pin/flat shoulder tool and a conical pin/concave shoulder
tool for the microhardness taken on the top
Sample ID Mean 2 μ 
FPS_800_0.5 46.543 82.8442 46.5429 9.102
FPS_800_1 75.099 4994.97 4.00167 0.796
FPS_1200_0.5 45.121 441.806 3.71115 0.443
FPS_1200_1 51.731 672.845 3.83391 0.474
CCS_800_0.5 44.575 409.802 3.70341 0.433
CCS_800_1 55.887 813.763 3.90757 0.481
CCS_1200_0.5 45.453 482.889 3.71166 0.458
CCS_1200_1 40.942 248.04 3.64315 0.371
Table 4: Mean, variance, mu and sigma of the spot samples produced
using a flat pin/flat shoulder tool and a conical pin/concave shoulder
tool for the microhardness taken at the bottom
Sample ID Mean 2 μ 
FPS_800_0.5 92.234 32.7196 4.52241 0.062
FPS_800_1 84.7 119.974 84.7 10.95
FPS_1200_0.5 84.7 119.974 84.7 10.95
FPS_1200_1 78.429 58.0637 78.4286 7.62
CCS_800_0.5 89.986 14.2859 89.9857 3.78
CCS_800_1 88.564 102.307 88.5643 10.11
CCS_1200_0.5 86.252 89.6604 4.45128 0.109
CCS_1200_1 82.043 54.8519 4.40318 0.09
It should be noted that the mean is equal to μ if the
distribution is normal. It can be seen that in some cases
in the current work μ and the mean have different values,
which shows that the distributions in some of the
analyses were not normal. The standard deviation 
should be close to zero, but in the current work, the value
of the standard deviation is not close to zero in some of
the cases. This shows that the microhardness values are
not close to the expected values and this could be due to
the microhardness values measured in different locations
of the weld samples, far apart from each other.
This was further suspected due to the presence of
intermetallics, which could have been the cause of the
high microhardness values since intermetallics are
invariably hard and brittle.
Each figure contains residuals versus the distance
(the distance from the keyhole), taking into account the
data and the constant variance of the residuals. In the
plot of residuals versus distance, it is shown that the
models are adequate to predict the responses in an
acceptable manner.
4 CONCLUSIONS
According to the presented results, some conclusions
can be drawn:
• The microstructure of the produced spot welds shows
a good material mixing, the presence of a copper ring
and a mixture of Al/Cu particles present in the stir
zone.
• The EDS analyses of the produced friction-stir spot
welds exhibited the presence of intermetallic com-
pounds, which are known to affect the microhard-
ness.
• The microhardness values obtained at the top were
high for all the samples and this was in the region
close to the spot-weld keyhole. In addition, all the
microhardness values obtained at the bottom of the
samples, in the region close to the keyhole, have
lower values, which were close to the average value
of the copper base material. This occurred for all the
spot welds produced using a conical pin and concave
shoulder.
• The probability-density-function (PDF) histograms
of the microhardness results revealed that the process
parameters and the tool geometries have significant
effects on the distribution of the microhardness
values in different locations of the produced spot
welds.
• Additionally, goodness-of-fit values were also
analysed and these showed that most of the R2 values
ranged between 0.8842 and 0.9999, which is an
indication of how well the model fits with the
produced experimental data.
Acknowledgements
The authors wish to acknowledge the financial
support of the University of Johannesburg and the
assistance from Mr. Riaan Brown (eNtsa at Nelson
Mandela Metropolitan University).
5 REFERENCES
1 W. M. Thomas, E. D. Nicholas, J. C. Needham, M. G. Murch, P.
Temple-Smith, C. J. Dawes, International Patent No. PCT/GB92/
02203, GB patent application No. 9125978.8, 1991
2 H. Badarinarayan, Fundamentals of Friction Stir Spot Welding, PhD
Thesis, 2009, Missouri
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
868 Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
3 W. Yuan, R. S. Mishra, B. Carlson, R. Verma, R. K. Mishra, Material
flow and microstructural evolution during friction stir spot welding
of AZ31 magnesium alloy, Materials Science and Engineering, A
543 (2012), 200–209
4 J. Podr`aj, B. Jerman, D. Klob~ar, Welding defects at friction stir
welding, Metalurgija, 54 (2015) 2, 387–389
5 J. M. Timothy, Friction Stir Welding of Commercially Available
Superplastic Aluminium, PhD thesis, 2008, Department of Engi-
neering and Design, Brunel University, Brunel
6 J. Ouyang, E. Yarrapareddy, R. Kovacevic, Microstructural evolution
in the friction stir welded 6061 aluminum alloy (T6-temper condi-
tion) to copper, Journal of Materials Processing Technology, 172
(2006), 110–122
7 P. Liu, Q. Shi, W. Wang, X. Wang, Z. Zhang, Microstructure and
XRD analysis of FSW joints for copper T2/aluminium 5A06
dissimilar materials, Materials Letters, 62 (2008), 4106–4108
8 P. Xue, B. L. Xiao, D. R. Ni, Z. Y. Ma, Enhanced mechanical pro-
perties of friction stir welded dissimilar Al–Cu joint by intermetallic
compounds, Materials Science and Engineering A, 527 (2010),
5723–5727
9 A. Esmaeili, M. K. Besharati Givi, H. R. Zareie Rajani, A metallur-
gical and mechanical study on dissimilar friction stir welding of
aluminum 1050 to brass (CuZn30), Materials Science and Engi-
neering, A 528 (2011), 7093–7102
10 A. Abdollah-Zadeh, T. Saeid, B. Sazgari, Microstructural and mecha-
nical properties of friction stir welded aluminum/copper lap joints,
Journal of Alloys and Compounds, 460 (2008), 535–538
11 E. T. Akinlabi, Characterisation of Dissimilar Friction Stir Welds
between 5754 Aluminium Alloy and C11000 Copper, D-Tech thesis,
Nelson Mandela Metropolitan University, South Africa, 2010
12 M. N. Avettand-Fenoel, R. Taillard, G. Ji, D. Goran, Multiscale
Study of Interfacial Intermetallic Compounds in a Dissimilar Al
6082-T6/Cu Friction-Stir Weld, Metallurgical and Materials
Transactions A, (2012), 4655–4666
13 M. F. X Muthu, V. Jayabalan, Tool Travel Speed Effects on the
Microstructure of Friction Stir Welded Aluminium – Copper Joints,
Journal of Materials Processing Technology, 217 (2015), 105–113
14 M. P. Mubiayi, E. T Akinlabi, Friction Stir Spot Welding between
Copper and Aluminium: Microstructural Evolution, Proceedings of
the International MultiConference of Engineers and Computer
Scientists, Vol II, IMECS 2015, March 18–20, 2015, Hong Kong
15 R. Heideman, C. Johnson, S. Kou, Metallurgical analysis of Al/Cu
friction stir spot welding, Science and Technology of Welding and
Joining, 15 (2010) 7, 597–604
16 U. Özdemir, S. Sayer, Ç. Yeni, Bornova-Izmir, Effect of Pin Pene-
tration Depth on the Mechanical Properties of Friction Stir Spot
Welded Aluminum and Copper, Materials Testing in Joining Techno-
logy, 54 (2012) 4, 233–239
17 M. Shiraly, M. Shamanian, M. R. Toroghinejad, M. Ahmadi Jazani,
Effect of Tool Rotation Rate on Microstructure and Mechanical
Behavior of Friction Stir Spot-Welded Al/Cu Composite, Journal of
Materials Engineering and Performance, 23 (2014) 2, 413–420
18 M. P. Mubiayi, E. T. Akinlabi, Evolving properties of friction stir
spot welds between AA1060 and commercially pure copper C11000,
Transactions of Nonferrous Metals Society of China, 26 (2016),
1852–1862
19 E. T Akinlabi, S. A. Akinlabi, Friction stir welding of dissimilar
materials – statistical analysis of the weld data, Proceedings of the
International MultiConference of Engineers and Computer Scientists,
2012, 1368–1373
20 ASTM Standard E3-11, Standard guide for preparation of
metallographic specimens, ASTM International, West Consho-
hocken, PA, 2011, doi: 10.1520/E0003-11, www.astm.org
21 L. E Murr, A Review of FSW Research on Dissimilar Metal and
Alloy Systems, Journal of Materials Engineering and Performance, 8
(2010)19, 1071–108
22 J. M. Schneider, M. Bigerelle, A. Iost, Statistical analysis of the
Vickers hardness, Materials Science and Engineering, A 262 (1999),
256–263
23 A. L. Yurkov, N. V. Jhuravleva, E. S. Lukin, Kinetic microhardness
measurements of sialon-based ceramics, Journal of Materials
Science, 29 (1994), 6551–6560
24 I. Y. Yanchev, E. P. Trifonova, Analysis of microhardness data in Tlx
ln l-x Se, Journal of Materials Science, 30 (1995), 5576–5580
25 A. M. Hassan, M. Almomani, T. Qasim, A. Ghaithan, Statistical
analysis of some mechanical properties of friction stir welded
aluminium matrix composite, Int. J. Experimental Design and
Process Optimisation, 3 (2012)1, 91–109
26 K. Palanikumar, R. Karthikeyan, Optimal machining conditions for
turning of particulate metal matrix composites using Taguchi and
response surface methodologies, Machining Science and Technology,
10 (2006) 4, 417–433
M. P. MUBIAYI et al.: MICROSTRUCTURE EVOLUTION AND STATISTICAL ANALYSIS OF Al/Cu ...
Materiali in tehnologije / Materials and technology 51 (2017) 5, 861–869 869
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
871–878
SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE
USED FOR RESIN TEETH
SINTEZA PMMA/ZnO NANODELCEV KOMPOZITOV ZA
IZDELAVO ZOB IZ UMETNIH SMOL
Danica Popovi}1, Rajko Bobovnik2, Silvester Bolka2, Miroslav Vukadinovi}1,
Vojkan Lazi}1, Rebeka Rudolf3,4
1University of Belgrade, School of Dental Medicine, Dr. Suboti}a 8, 11000 Belgrade, Serbia
2Faculty of Polymer Technology, Ozare 19, 2380 Slovenj Gradec, Slovenia
3Zlatarna Celje d.o.o., Kersnikova ulica 19, 3000 Celje, Slovenia
4University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia
d.popovic984@gmail.com
Prejem rokopisa – received: 2017-02-24; sprejem za objavo – accepted for publication: 2017-03-31
doi:10.17222/mit.2017.025
Wear resistance is one of the most important physical properties of the artificial teeth used in acrylic dentures. The goal of this
research was to synthesize a new composite material made of matrix Poly-(methyl methacrylate)-PMMA with different
percentages (2 % and 3 % of volume fractions) of zinc-oxide nanoparticles (ZnO NPs) as reinforcing elements, to improve its
mechanical properties. The dynamic mechanical behaviour of this composite was studied through the DMA method in
comparison to the pure PMMA supported by the characterization of their microstructures. Then the wear resistance was
analysed on the samples, which were prepared in the form of teeth. In this context their vertical height loss was measured after
100,000 chewing cycles on a chewing simulator, before and after the artificial thermal ageing. Investigations showed that the
PMMA/ZnO NP composites dampened the vibrations better than the pure PMMA, which could be assigned to the homogenous
distribution of ZnO NPs in the PMMA matrix. It was found that the mean vertical height loss for the pure PMMA teeth was
significantly higher (more than 4 times) compared to composite teeth made with ZnO NPs. Introducing the thermal artificial
ageing led to the finding that there was no effect on the height loss by the composite material with 3 % of volume fractions of
ZnO NPs. Based on this it was concluded that PMMA/ZnO NPs composites showed improved in-vitro wear resistance
compared to acrylic-resin denture teeth, so this new composite material should be preferred when occlusal stability is
considered to be of high priority.
Keywords: poly-(methyl methacrylate)–PMMA, zinc-oxide nanoparticles, composite, resin teeth
Odpornost proti obrabi je ena izmed najbolj pomembnih fizikalnih lastnosti umetnih zob pri akrilnih protezah. Cilj te raziskave
je bil sintetiziranje novega kompozitnega materiala, izdelanega iz matrice poli-(metil meta akrilata) – PMMA z razli~nimi
volumskimi odstotki (2 % in 3 %) nanodelcev cinkovega oksida (ZnO NPs), uporabljenimi kot oja~itveni element za izbolj{anje
mehanskih lastnosti materiala. Dinami~no mehansko obna{anje teh kompozitov smo raziskali z metodo DMA v primerjavi s
~istim PMMA, ter s karakterizacijo njihovih mikrostruktur. Nato je bila analizirana odpornost proti obrabi na vzorcih, ki so bili
pripravljeni v obliki zob. Izmerjena je bila njihova navpi~na izguba vi{ine po 100.000 `ve~ilnih ciklih na `ve~ilnem simulatorju,
pred in po umetnem toplotnem staranju. Preiskave so pokazale, da so PMMA/ZnO NPs kompoziti bolje bla`ili vibracije kot ~isti
PMMA, kar lahko pripi{emo homogeni porazdelitvi ZnO NPs v PMMA matrici. Ugotovili smo, da je bila povpre~na navpi~na
izguba vi{ina za ~isti PMMA zob znatno vi{ja (ve~ kot 4×) v primerjavi s kompozitnimi zobmi z ZnO NPs. Umetno toplotno
staranje ni imelo u~inka na izgubo navpi~ne vi{ine zoba iz kompozitnega materiala s 3 volumskimi % ZnO NPs. Na podlagi tega
smo sklenili, da imajo kompoziti PMMA/ZnO NPs izbolj{ano in vitro odpornost proti obrabi, v primerjavi z zobno protezo iz
akrilne smole, zato naj bi imel ta novi kompozitni material prednost pri uporabi, kadar je potrebna okluzalna stabilnost.
Klju~ne besede: poli-(metil meta akrilat)-PMMA, nanodelci cinkovega oksida, kompoziti, zobje iz umetnih smol
1 INTRODUCTION
Teeth loss in the human population is associated with
numerous problems such as chewing difficulties,
alteration of speech and facial expressions, which all
leads to de-socialization. According to the Oral
Health-Healthy People 2010: Objectives for Improving
Health 26 % of the US population and 33 % of the
Central Europe population aged between 65 years and 74
years are toothless; but some dramatic data and a wide
variation in edentulism prevalence among adults aged 50
and above was found in the following countries: 48.3 %
in Ireland, Malaysia 56.6 % and Netherlands 65.4 %,
which indicates an international problem.1–3
The problems of toothless people in the region of the
Balkan peninsula and South-Eastern Europe are usually
solved with a complete or partial acrylic denture, which
is the easiest and the cheapest solution, because more
than 90 % of these patients are in a bad financial
situation and incapable of buying dental implants and
hybrid bridges above them. With an acrylic denture the
wear resistance represents one of the most important
physical properties because, in the case that the material
has excessive wear, that might cause loss of the vertical
dimension of occlusion (VDO) which is associated with
decreased occlusal forces, loss of masticatory efficiency,
improper tooth relationship, and fatigue of masticatory
muscles.4 Recently, a study examined how VDO can
Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878 871
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.3:616.314-77:678.6:67.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)871(2017)
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
871–878
SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE
USED FOR RESIN TEETH
SINTEZA PMMA/ZnO NANODELCEV KOMPOZITOV ZA
IZDELAVO ZOB IZ UMETNIH SMOL
Danica Popovi}1, Rajko Bobovnik2, Silvester Bolka2, Miroslav Vukadinovi}1,
Vojkan Lazi}1, Rebeka Rudolf3,4
1University of Belgrade, School of Dental Medicine, Dr. Suboti}a 8, 11000 Belgrade, Serbia
2Faculty of Polymer Technology, Ozare 19, 2380 Slovenj Gradec, Slovenia
3Zlatarna Celje d.o.o., Kersnikova ulica 19, 3000 Celje, Slovenia
4University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia
d.popovic984@gmail.com
Prejem rokopisa – received: 2017-02-24; sprejem za objavo – accepted for publication: 2017-03-31
doi:10.17222/mit.2017.025
Wear resistance is one of the most important physical properties of the artificial teeth used in acrylic dentures. The goal of this
research was to synthesize a new composite material made of matrix Poly-(methyl methacrylate)-PMMA with different
percentages (2 % and 3 % of volume fractions) of zinc-oxide nanoparticles (ZnO NPs) as reinforcing elements, to improve its
mechanical properties. The dynamic mechanical behaviour of this composite was studied through the DMA method in
comparison to the pure PMMA supported by the characterization of their microstructures. Then the wear resistance was
analysed on the samples, which were prepared in the form of teeth. In this context their vertical height loss was measured after
100,000 chewing cycles on a chewing simulator, before and after the artificial thermal ageing. Investigations showed that the
PMMA/ZnO NP composites dampened the vibrations better than the pure PMMA, which could be assigned to the homogenous
distribution of ZnO NPs in the PMMA matrix. It was found that the mean vertical height loss for the pure PMMA teeth was
significantly higher (more than 4 times) compared to composite teeth made with ZnO NPs. Introducing the thermal artificial
ageing led to the finding that there was no effect on the height loss by the composite material with 3 % of volume fractions of
ZnO NPs. Based on this it was concluded that PMMA/ZnO NPs composites showed improved in-vitro wear resistance
compared to acrylic-resin denture teeth, so this new composite material should be preferred when occlusal stability is
considered to be of high priority.
Keywords: poly-(methyl methacrylate)–PMMA, zinc-oxide nanoparticles, composite, resin teeth
Odpornost proti obrabi je ena izmed najbolj pomembnih fizikalnih lastnosti umetnih zob pri akrilnih protezah. Cilj te raziskave
je bil sintetiziranje novega kompozitnega materiala, izdelanega iz matrice poli-(metil meta akrilata) – PMMA z razli~nimi
volumskimi odstotki (2 % in 3 %) nanodelcev cinkovega oksida (ZnO NPs), uporabljenimi kot oja~itveni element za izbolj{anje
mehanskih lastnosti materiala. Dinami~no mehansko obna{anje teh kompozitov smo raziskali z metodo DMA v primerjavi s
~istim PMMA, ter s karakterizacijo njihovih mikrostruktur. Nato je bila analizirana odpornost proti obrabi na vzorcih, ki so bili
pripravljeni v obliki zob. Izmerjena je bila njihova navpi~na izguba vi{ine po 100.000 `ve~ilnih ciklih na `ve~ilnem simulatorju,
pred in po umetnem toplotnem staranju. Preiskave so pokazale, da so PMMA/ZnO NPs kompoziti bolje bla`ili vibracije kot ~isti
PMMA, kar lahko pripi{emo homogeni porazdelitvi ZnO NPs v PMMA matrici. Ugotovili smo, da je bila povpre~na navpi~na
izguba vi{ina za ~isti PMMA zob znatno vi{ja (ve~ kot 4×) v primerjavi s kompozitnimi zobmi z ZnO NPs. Umetno toplotno
staranje ni imelo u~inka na izgubo navpi~ne vi{ine zoba iz kompozitnega materiala s 3 volumskimi % ZnO NPs. Na podlagi tega
smo sklenili, da imajo kompoziti PMMA/ZnO NPs izbolj{ano in vitro odpornost proti obrabi, v primerjavi z zobno protezo iz
akrilne smole, zato naj bi imel ta novi kompozitni material prednost pri uporabi, kadar je potrebna okluzalna stabilnost.
Klju~ne besede: poli-(metil meta akrilat)-PMMA, nanodelci cinkovega oksida, kompoziti, zobje iz umetnih smol
1 INTRODUCTION
Teeth loss in the human population is associated with
numerous problems such as chewing difficulties,
alteration of speech and facial expressions, which all
leads to de-socialization. According to the Oral
Health-Healthy People 2010: Objectives for Improving
Health 26 % of the US population and 33 % of the
Central Europe population aged between 65 years and 74
years are toothless; but some dramatic data and a wide
variation in edentulism prevalence among adults aged 50
and above was found in the following countries: 48.3 %
in Ireland, Malaysia 56.6 % and Netherlands 65.4 %,
which indicates an international problem.1–3
The problems of toothless people in the region of the
Balkan peninsula and South-Eastern Europe are usually
solved with a complete or partial acrylic denture, which
is the easiest and the cheapest solution, because more
than 90 % of these patients are in a bad financial
situation and incapable of buying dental implants and
hybrid bridges above them. With an acrylic denture the
wear resistance represents one of the most important
physical properties because, in the case that the material
has excessive wear, that might cause loss of the vertical
dimension of occlusion (VDO) which is associated with
decreased occlusal forces, loss of masticatory efficiency,
improper tooth relationship, and fatigue of masticatory
muscles.4 Recently, a study examined how VDO can
Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878 871
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
UDK 620.3:616.314-77:678.6:67.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(5)871(2017)
affect brain function in complete denture wearers, and
also measured occlusal forces when there is a change in
VDO in which electroencephalograms were used.5 Many
authors have pointed out the necessity of creating
complete new dentures, or relining of the denture base,
which is not a simple procedure and presents financial
pressure for the patients.6 So, finding the new combi-
nation of materials with improved mechanical properties,
which include better wear resistance, would be helpful
for many reasons for those patients.
Many basic researches based on polymer materials
represent the challenge. One of these are made of
supramolecular polymers as a new high-tech material.7,8
Over the past two decades a great effort has been made
by the community of researchers for the synthesis and
utilization of nanomaterials in the field of medical
materials.9 Using fine nanoparticles as reinforcing
elements could improve the functional properties of
conventional materials such as PMMA, which has been
used for dentures since 1937. The ZnO NPs are small
objects that behave as an entire unit regarding their
transport and properties. The introduction of ZnO NPs
represents a new approach in dentistry by the production
of a denture, because they could decrease the level of
residual monomer and, besides this, they also have anti-
microbial activity. This was shown in recent studies.10,11
Namely, ZnO NPs are environmentally friendly materials
in the nanometre size range, and could be synthesized in
a wide range of particle shapes and structures.12–14 ZnO
in bulk form is a largely inert, white powder compound
that has a very broad application, from the chemical
industry through cosmetic and medical products. In
dentistry they have an important place as a medicament
for healing injures after periodontal surgical treatment, in
the cementation of temporary, as well as definitive,
crowns etc. ZnO is also a biocompatible material that
exhibits antimicrobial properties against gram-negative
as well as gram-positive bacteria.15 A recent study also
investigated, proved and determined the minimal
inhibitory concentration of ZnO NPs against Candida
Albicans.16
Nanostructured materials with unique and fascinating
properties motivate scientists tremendously to explore
and understand their formation and growth processes and
following the critical volume in several studies we tried
to find the right percentage of ZnO NPs in a PMMA
matrix which could advance the properties of this new
composite, because increasing the percentage of added
NPs could downgrade some of them.17,18
According to the presented state-of-the-art study, the
aim of this research was:
(i) To synthesize a composite PMMA/ZnO NPs material,
to test its mechanical properties, and evaluate the
developed microstructures;
(ii) To investigate the wear resistance of resin teeth in a
chewing simulator with two combinations, before and
after thermal artificial ageing. This investigation
included not only the classical testing samples, but
also the real model of the shapes of first upper molars
(testing resin teeth) and showed possible practical
applications in dentistry.
2 EXPERIMENTAL PART
2.1 Materials for composite PMMA/ZnO NPs
Commercially available ZnO nanopowder (purity
99.99 %, density 5.606 g/cm3, insoluble in water, for-
mula weight 81.37, melting point 1975 °C) was pur-
chased from Interdent (Belgrade, Serbia). The average
size of the ZnO NPs was 30 nm and they were in the
form of a sphere. The typical shape and morphology of
characteristic ZnO NPs are shown in Figure 1, which
was made on a transmission electron microscope (TEM).
PMMA matrix (polymethylmethacrylate powder with
benzoyl-peroxide, pigments and for a liquid-methyl-
methacrylate HQ 60, ethylene glycoldimethacrylate,
methylmethacrylate composite) was distributed by
Galenika (Belgrade, Serbia). This material was used for
Biogal® acrylic teeth (Galenika, Belgrade, Serbia) and
has already passed through the applied ISO Standards.
2.2 Measurement of ZnO nanoparticles’ size distri-
bution and zeta potential in water and MMA
Because of the frequent problem with agglomeration
among nanoparticles, especially ZnO NPs in water, we
used the DLS technique to confirm or deny this fact
indirectly and compare them with ZnO NPs in suspen-
sion with MMA. Initially, ZnO NPs powder (0.05 g) with
water and MMA monomer (15 mL by volume in both
cases) was mixed through magnetic stirring for the time
duration of 45 min. The size distributions and surface
charge of the nanoparticles were determined with a
Malvern Zetasizer Nano ZS (Malvern Instruments Ltd.,
U.K.) by the DLS technique. The suitable parameters for
ZnO NPs Absorption: 0.1 and Refractive index: 2.0.
were chosen according to 19.
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
872 Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Micrograph of shape and morphology of characteristic ZnO
NPs (TEM, 10 000×)
2.3 Preparation of samples
Samples were made in two shapes: Lamina (dimen-
sion: 45 mm × 3 mm × 7 mm) for characterization of the
new composite and then in the form of the first upper
molar for direct potential use in dentistry.
The ZnO NPs in the PMMA matrix were mixed with
PMMA powder and then free radical chain poly-
merization was performed of MMA monomer in bulk.
The new composite was synthesized at the selection of
the appropriate technological regime, which started in
the moment of adding monomer liquid PMMA/ZnO NPs
powder for 1 h and 45 min at 100 °C. The procedure for
both samples followed by shaping the samples firstly in
wax, then duplicated and changed in to the pure PMMA,
PMMA filled with 2 % of volume fractions and 3 % of
volume fractions of ZnO NPs. The volume % were
chosen according to our previous experiences and based
on 7,16,18. All samples in the form of the first upper molar
were identical in form, plan parallel and highly polished.
They were prepared for a wear test in a chewing
simulator. Schematic presentations of the composite
PMMA/ZnO NPs synthesis and their corresponding
testing are shown in Figure 2. A list of samples’ names,
their mixing ratio PMMA/ZnO NPs and technical
descriptions are given in Table 1. PMMA1 was chosen
for a comparison with the new composites.
2.4 Characterization
Thermo-mechanical properties were examined using
a Perkin Elmer DMA 8000 dynamic mechanical anal-
yser. The TT_DMA software, Version 14310, was used
for the evaluation of the results. The viscoelastic proper-
ties of the samples were analysed by recording the
storage modulus (E’), loss modulus (E’’) and loss factor
(tan ) as a function of temperature. The height of the
peak of the loss factor determined the damping beha-
viour – with decreasing peak height the damping beha-
viour was increasing. For the analyses of test samples,
the DMA instrument was operated in the dual cantilever
mode. The viscoelastic analyses were carried out on
samples with dimensions of approximately 42 mm ×
5 mm × 2 mm. The samples were heated at 2 °C/min
from room temperature to 180 °C under an air atmo-
sphere. A frequency of 10 Hz and amplitude of 20 μm
were used.
A flash differential scanning calorimeter (Flash DSC)
was used for the identification of the glass transition and
reorganisation of the polymers for all the samples. Flash
DSC works by ultra heating and cooling rates that induce
physical transitions and chemical processes. This charac-
terization was very important for dentistry practice
because it allowed us to follow the level of polymeri-
zation in the residual monomer, which could produce an
allergic reaction by the use of the new composite for the
resin teeth in removable dentures in the mouths of
patients.
2.5 Determination of wear resistance
The wear tests of the samples in the shape of upper
first molar were performed in a chewing simulator
CS-4.2 economy line (SD Mechatronics, Germany). The
simulator has two identical sample chambers and two
stepper motors which allow computer-controlled vertical
and horizontal movements. The masticatory cycle in this
research consisted of three phases: contact with a vertical
load of 5 kg, horizontal sliding of 0.4 mm, vertical
sliding of 0.2 mm and separating the teeth and machine
antagonist. In our investigation, two types of tests were
made: before and after thermal artificial ageing including
all samples from Table 1. The machine was set at
100.000 cycles, which simulated chewing over a period
of one year. The loss of substance was measured with an
electronic digital calliper – Globathronics GmbH (Ger-
many). Accelerated thermal ageing was carried out by
immersing the samples in a water bath at a temperature
of 70 °C.20,21
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878 873
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Table 1: Composition, mixing ratio and forms of samples for different testing
Samples/
abbreviation
Composition Mixing ratio Form forcharacterization Form for wear test
PMMA1 Pure PMMA 15 mL MMA + 23.4 gPMMA Lamina Tooth model
PMMA2 PMMA/2vol%ZnO NPs 15 mL + 22.93 g PMMA+ 0,47 g ZnO NPs Lamina Tooth model
PMMA3 PMMA/3vol%ZnO NPs 15 mL + 22.7 g PMMA+ 0,7 g ZnO NPs Lamina Tooth model
Figure 2: Schematic presentation of synthesis composite PMMA/ZnO
NPs and samples for testing
affect brain function in complete denture wearers, and
also measured occlusal forces when there is a change in
VDO in which electroencephalograms were used.5 Many
authors have pointed out the necessity of creating
complete new dentures, or relining of the denture base,
which is not a simple procedure and presents financial
pressure for the patients.6 So, finding the new combi-
nation of materials with improved mechanical properties,
which include better wear resistance, would be helpful
for many reasons for those patients.
Many basic researches based on polymer materials
represent the challenge. One of these are made of
supramolecular polymers as a new high-tech material.7,8
Over the past two decades a great effort has been made
by the community of researchers for the synthesis and
utilization of nanomaterials in the field of medical
materials.9 Using fine nanoparticles as reinforcing
elements could improve the functional properties of
conventional materials such as PMMA, which has been
used for dentures since 1937. The ZnO NPs are small
objects that behave as an entire unit regarding their
transport and properties. The introduction of ZnO NPs
represents a new approach in dentistry by the production
of a denture, because they could decrease the level of
residual monomer and, besides this, they also have anti-
microbial activity. This was shown in recent studies.10,11
Namely, ZnO NPs are environmentally friendly materials
in the nanometre size range, and could be synthesized in
a wide range of particle shapes and structures.12–14 ZnO
in bulk form is a largely inert, white powder compound
that has a very broad application, from the chemical
industry through cosmetic and medical products. In
dentistry they have an important place as a medicament
for healing injures after periodontal surgical treatment, in
the cementation of temporary, as well as definitive,
crowns etc. ZnO is also a biocompatible material that
exhibits antimicrobial properties against gram-negative
as well as gram-positive bacteria.15 A recent study also
investigated, proved and determined the minimal
inhibitory concentration of ZnO NPs against Candida
Albicans.16
Nanostructured materials with unique and fascinating
properties motivate scientists tremendously to explore
and understand their formation and growth processes and
following the critical volume in several studies we tried
to find the right percentage of ZnO NPs in a PMMA
matrix which could advance the properties of this new
composite, because increasing the percentage of added
NPs could downgrade some of them.17,18
According to the presented state-of-the-art study, the
aim of this research was:
(i) To synthesize a composite PMMA/ZnO NPs material,
to test its mechanical properties, and evaluate the
developed microstructures;
(ii) To investigate the wear resistance of resin teeth in a
chewing simulator with two combinations, before and
after thermal artificial ageing. This investigation
included not only the classical testing samples, but
also the real model of the shapes of first upper molars
(testing resin teeth) and showed possible practical
applications in dentistry.
2 EXPERIMENTAL PART
2.1 Materials for composite PMMA/ZnO NPs
Commercially available ZnO nanopowder (purity
99.99 %, density 5.606 g/cm3, insoluble in water, for-
mula weight 81.37, melting point 1975 °C) was pur-
chased from Interdent (Belgrade, Serbia). The average
size of the ZnO NPs was 30 nm and they were in the
form of a sphere. The typical shape and morphology of
characteristic ZnO NPs are shown in Figure 1, which
was made on a transmission electron microscope (TEM).
PMMA matrix (polymethylmethacrylate powder with
benzoyl-peroxide, pigments and for a liquid-methyl-
methacrylate HQ 60, ethylene glycoldimethacrylate,
methylmethacrylate composite) was distributed by
Galenika (Belgrade, Serbia). This material was used for
Biogal® acrylic teeth (Galenika, Belgrade, Serbia) and
has already passed through the applied ISO Standards.
2.2 Measurement of ZnO nanoparticles’ size distri-
bution and zeta potential in water and MMA
Because of the frequent problem with agglomeration
among nanoparticles, especially ZnO NPs in water, we
used the DLS technique to confirm or deny this fact
indirectly and compare them with ZnO NPs in suspen-
sion with MMA. Initially, ZnO NPs powder (0.05 g) with
water and MMA monomer (15 mL by volume in both
cases) was mixed through magnetic stirring for the time
duration of 45 min. The size distributions and surface
charge of the nanoparticles were determined with a
Malvern Zetasizer Nano ZS (Malvern Instruments Ltd.,
U.K.) by the DLS technique. The suitable parameters for
ZnO NPs Absorption: 0.1 and Refractive index: 2.0.
were chosen according to 19.
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
872 Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 1: Micrograph of shape and morphology of characteristic ZnO
NPs (TEM, 10 000×)
2.3 Preparation of samples
Samples were made in two shapes: Lamina (dimen-
sion: 45 mm × 3 mm × 7 mm) for characterization of the
new composite and then in the form of the first upper
molar for direct potential use in dentistry.
The ZnO NPs in the PMMA matrix were mixed with
PMMA powder and then free radical chain poly-
merization was performed of MMA monomer in bulk.
The new composite was synthesized at the selection of
the appropriate technological regime, which started in
the moment of adding monomer liquid PMMA/ZnO NPs
powder for 1 h and 45 min at 100 °C. The procedure for
both samples followed by shaping the samples firstly in
wax, then duplicated and changed in to the pure PMMA,
PMMA filled with 2 % of volume fractions and 3 % of
volume fractions of ZnO NPs. The volume % were
chosen according to our previous experiences and based
on 7,16,18. All samples in the form of the first upper molar
were identical in form, plan parallel and highly polished.
They were prepared for a wear test in a chewing
simulator. Schematic presentations of the composite
PMMA/ZnO NPs synthesis and their corresponding
testing are shown in Figure 2. A list of samples’ names,
their mixing ratio PMMA/ZnO NPs and technical
descriptions are given in Table 1. PMMA1 was chosen
for a comparison with the new composites.
2.4 Characterization
Thermo-mechanical properties were examined using
a Perkin Elmer DMA 8000 dynamic mechanical anal-
yser. The TT_DMA software, Version 14310, was used
for the evaluation of the results. The viscoelastic proper-
ties of the samples were analysed by recording the
storage modulus (E’), loss modulus (E’’) and loss factor
(tan ) as a function of temperature. The height of the
peak of the loss factor determined the damping beha-
viour – with decreasing peak height the damping beha-
viour was increasing. For the analyses of test samples,
the DMA instrument was operated in the dual cantilever
mode. The viscoelastic analyses were carried out on
samples with dimensions of approximately 42 mm ×
5 mm × 2 mm. The samples were heated at 2 °C/min
from room temperature to 180 °C under an air atmo-
sphere. A frequency of 10 Hz and amplitude of 20 μm
were used.
A flash differential scanning calorimeter (Flash DSC)
was used for the identification of the glass transition and
reorganisation of the polymers for all the samples. Flash
DSC works by ultra heating and cooling rates that induce
physical transitions and chemical processes. This charac-
terization was very important for dentistry practice
because it allowed us to follow the level of polymeri-
zation in the residual monomer, which could produce an
allergic reaction by the use of the new composite for the
resin teeth in removable dentures in the mouths of
patients.
2.5 Determination of wear resistance
The wear tests of the samples in the shape of upper
first molar were performed in a chewing simulator
CS-4.2 economy line (SD Mechatronics, Germany). The
simulator has two identical sample chambers and two
stepper motors which allow computer-controlled vertical
and horizontal movements. The masticatory cycle in this
research consisted of three phases: contact with a vertical
load of 5 kg, horizontal sliding of 0.4 mm, vertical
sliding of 0.2 mm and separating the teeth and machine
antagonist. In our investigation, two types of tests were
made: before and after thermal artificial ageing including
all samples from Table 1. The machine was set at
100.000 cycles, which simulated chewing over a period
of one year. The loss of substance was measured with an
electronic digital calliper – Globathronics GmbH (Ger-
many). Accelerated thermal ageing was carried out by
immersing the samples in a water bath at a temperature
of 70 °C.20,21
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878 873
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Table 1: Composition, mixing ratio and forms of samples for different testing
Samples/
abbreviation
Composition Mixing ratio Form forcharacterization Form for wear test
PMMA1 Pure PMMA 15 mL MMA + 23.4 gPMMA Lamina Tooth model
PMMA2 PMMA/2vol%ZnO NPs 15 mL + 22.93 g PMMA+ 0,47 g ZnO NPs Lamina Tooth model
PMMA3 PMMA/3vol%ZnO NPs 15 mL + 22.7 g PMMA+ 0,7 g ZnO NPs Lamina Tooth model
Figure 2: Schematic presentation of synthesis composite PMMA/ZnO
NPs and samples for testing
2.6 Microstructure analysis
Microstructural characterizations of all the samples
were carried out with scanning electron microscopy
(SEM-Sirion 400 NC and Quanta 200 3D). The
specimens, without previous preparation, were broken in
an atmosphere of N2 to minimise the influence of
metallographic preparation regarding changes in the
structure of the composite. After breaking, the samples
were sprayed with Au (Jeol JSM 8310 appliance), which
enables the observation of the non-conductive surfaces of
the samples with an electron beam. The samples were
positioned into the chamber of the microscope and
observations were performed with an accelerating
voltage of 15 kV.
2.7 Statistical analysis
The statistical analysis included all the parameters
involved in the laboratory tests and it was performed in
the statistical package SPSS® 17.0. Multivariate linear
regression analysis was used to determine the predictors
of differences between the analysed groups of
samples. The threshold value for accepting the working
hypothesis was set at p < 0.05.
3 RESULTS AND DISCUSSION
3.1 Macroscopic view
A macroscopic view of both samples’ shapes
confirmed the homogenous structure of PMMA and
PMMA/ZnO NPs. There was only an acceptable diffe-
rence in colour. Adding ZnO NPs to the mixture with
PMMA gave a brighter colour as the percentage of
nanoparticles grew. This could be solved in the future by
the addition of proper pigment, which results in a similar
colour as in the teeth or oral mucosa, which are
replicated in resin teeth and denture bases.
3.2 Size distribution and zeta-potential by DLS
The statement of the homogeneous distribution of
ZnO NPs in samples was also confirmed by DLS
measurement of the results, which are shown in Table 2.
In ZnO NPs with MMA suspension, the average particle
size was 2.21±0.10 μm, while 42.1 % of the particles
were in the size range of 21.04 nm, whereas in ZnO NPs
with water suspension, the average particle size was
4.01±0.4 μm, while 38.1 % of the particles were in the
size range of 68.06 nm. Therefore, the average size of
the particles was greater in ZnO NPs with water
suspension with more presence of larger-sized clusters of
particles. Distribution of ZnO NPs in MMA suspension
was homogenous, which showed additionally the
insolubility of ZnO NPs in water. ZnO NPs with MMA
had a sharp and narrow intensity vs. size curve, as com-
pared to ZnO NPs with water; thereby, it can be stated
that there is a higher degree of agglomeration in the ZnO
NPs with water suspension (Figure 3). The present study
showed how it is possible to predict, simulate and
characterize the composite material, which is aimed for
use in dentistry. The DLS technique could confirm
indirectly that ZnO NPs in an MMA suspension give a
smaller agglomeration, which was also proven through
the calculated large hydrodynamic diameters. ZnO NPs
which have a zeta-potential between minus 30 mV and
plus 30 mV show a tendency for coagulation. Therefore,
ZnO NPs suspensions with water were unstable solutions
and showed the tendency of agglomerate, while adding
MMA in the ZnO NPs did not show much difference in
the stability of the suspension as compared with the
water. The samples in our work were made by mixing
first ZnO NPs nanopowder and PMMA powder matrix in
vacuum, at room temperature for 2 h then adding MMA,
which reduced the aggregation among ZnO NPs. This
may allow us to prepare homogenous nanocomposites
with organic matrices without additional surface
modification. The chosen average ZnO NPs size showed
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
874 Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Number vs. size distribution curve of ZnO NPs with MMA
and water
Table 2: Composition and the type of mixing used in the preparation of the solvent with the DLS measurements
Exp.
No. Composition Methodology
DLS Measurement
Run of
measurement Average size range (ìm) Zeta potential (mV)
1 0.05 g ZnONPs+ 15 mL H20
Magnetic stirring
(45 min)
I 3.50
4.01 ± 0.40
–3.45
–3.38 ± 0.81II 4.08 –2.35
III 4.48 –4.35
2 0.05 g ZnONPs+ 15 mL MMA
Magnetic stirring
(45 min)
I 2.28
2.21 ± 0.10
–3.40
–3.48 ± 0.07II 2.28 –3.57
III 2.06 –3.49
that is of vital importance for the final properties of the
new developed composite PMMA/ZnO NPs since it
affected the UV-vis absorption and thermal stability.
During the synthesis it was found that smaller ZnO NPs
showed a higher degree of agglomeration, which is in
accordance with the literature.22–24
3.3 Determination of properties
The results obtained by the DMA (Figure 4) showed
the following results. The difference in the elastic
modulus at low temperatures which may be due to a
slight decrease in the cross linking of PMMA by adding
ZnO NPs. At elevated temperatures, the incidence of
chain transfer increased (at termination). With chain
transfer, the propagating polymer radical reacts with
another molecule by proton abstraction and this leads to
branching and cross-linking.
Partly, the E modulus increases due to the addition of
ZnO NPs, but this increase is minimal. Because of the
absence of good interactions between the ZnO NPs and
the PMMA matrix, the loss factor increases with the
amount of addition of ZnO NPs in the PMMA matrix.
ZnO NPs added to the PMMA increased the E modulus
and reduced the level of cross-linking. At the same time,
the poor interaction between the filler and the matrix
results in a less elastic response, which means that the
composite of PMMA/ZnO NPs dampen the vibrations
better than pure PMMA. With the increasing of the loss
factor (tan delta) the damping of the material also
increases. In our case, the height of the loss factor peak
of the sample PMMA3 was the highest; therefore, the
damping of the vibrations of the sample PMMA3 is the
highest (Figure 5).
Characterization using the DMA method showed that
changes in the vibration-damping behaviour of the new
composite could mean that at the gingival a more gentle
feeling appears, which leads to this material being more
friendly to oral mucosa due to vibration damping. This
should also be proven through practice. Composite
PMMA/ZnO NPs would require longer curing times in
comparison with the pure PMMA. Based on this finding
it is expected to obtain a better material with the surface
modification of ZnO NPs where good interactions
between filler and matrix can be formed.
From the Flash DSC measurements (Figure 6), the
lowest glass-transition temperature was measured in the
PMMA3 sample (129.6 °C), followed by PMMA2
(131.0 °C) and the highest glass-transition temperature
was in the pure PMMA1 (131.1 °C). This fact is also
confirmed by the position of the peaks of the loss factor
(tan ), where the lowest glass-transition temperature
was measured in the PMMA3 (146.1 °C), followed by
PMMA2 (147.2°C) and the highest glass-transition
temperature in the pure PMMA1 (147.8 °C). In this
context it was found that the degree of cross-linking was
reduced by the addition of ZnO NPs. With increasing
glass-transition temperature the loss factor decreased,
which indicates that the material PMMA3 (tan  =
1.2003) has the best vibration-damping behaviour,
followed by PMMA2 (tan  = 1.1856). The pure
PMMA1 (tan  = 1.1839) showed the most elastic
response. From the height of tan  it can be concluded
that the 3 % addition of ZnO NPs in PMMA is needed to
obtain a significant change in the vibration-damping
behaviour of the composite. Despite the small differen-
ces measured, as well on DMA (E modulus, tan ) as on
Flash DSC (glass-transition temperatures), all the results
had the same tendency.
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878 875
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Results of flash DSC analysis for all samples, heating
1.000 °C/s (diagram heat flow (mW) – T (°C))
Figure 4: DMA results- diagram E – modulus (N/mm2) – T (°C) for
all samples
Figure 5: DMA results-loss factor vs. temperature for all samples
2.6 Microstructure analysis
Microstructural characterizations of all the samples
were carried out with scanning electron microscopy
(SEM-Sirion 400 NC and Quanta 200 3D). The
specimens, without previous preparation, were broken in
an atmosphere of N2 to minimise the influence of
metallographic preparation regarding changes in the
structure of the composite. After breaking, the samples
were sprayed with Au (Jeol JSM 8310 appliance), which
enables the observation of the non-conductive surfaces of
the samples with an electron beam. The samples were
positioned into the chamber of the microscope and
observations were performed with an accelerating
voltage of 15 kV.
2.7 Statistical analysis
The statistical analysis included all the parameters
involved in the laboratory tests and it was performed in
the statistical package SPSS® 17.0. Multivariate linear
regression analysis was used to determine the predictors
of differences between the analysed groups of
samples. The threshold value for accepting the working
hypothesis was set at p < 0.05.
3 RESULTS AND DISCUSSION
3.1 Macroscopic view
A macroscopic view of both samples’ shapes
confirmed the homogenous structure of PMMA and
PMMA/ZnO NPs. There was only an acceptable diffe-
rence in colour. Adding ZnO NPs to the mixture with
PMMA gave a brighter colour as the percentage of
nanoparticles grew. This could be solved in the future by
the addition of proper pigment, which results in a similar
colour as in the teeth or oral mucosa, which are
replicated in resin teeth and denture bases.
3.2 Size distribution and zeta-potential by DLS
The statement of the homogeneous distribution of
ZnO NPs in samples was also confirmed by DLS
measurement of the results, which are shown in Table 2.
In ZnO NPs with MMA suspension, the average particle
size was 2.21±0.10 μm, while 42.1 % of the particles
were in the size range of 21.04 nm, whereas in ZnO NPs
with water suspension, the average particle size was
4.01±0.4 μm, while 38.1 % of the particles were in the
size range of 68.06 nm. Therefore, the average size of
the particles was greater in ZnO NPs with water
suspension with more presence of larger-sized clusters of
particles. Distribution of ZnO NPs in MMA suspension
was homogenous, which showed additionally the
insolubility of ZnO NPs in water. ZnO NPs with MMA
had a sharp and narrow intensity vs. size curve, as com-
pared to ZnO NPs with water; thereby, it can be stated
that there is a higher degree of agglomeration in the ZnO
NPs with water suspension (Figure 3). The present study
showed how it is possible to predict, simulate and
characterize the composite material, which is aimed for
use in dentistry. The DLS technique could confirm
indirectly that ZnO NPs in an MMA suspension give a
smaller agglomeration, which was also proven through
the calculated large hydrodynamic diameters. ZnO NPs
which have a zeta-potential between minus 30 mV and
plus 30 mV show a tendency for coagulation. Therefore,
ZnO NPs suspensions with water were unstable solutions
and showed the tendency of agglomerate, while adding
MMA in the ZnO NPs did not show much difference in
the stability of the suspension as compared with the
water. The samples in our work were made by mixing
first ZnO NPs nanopowder and PMMA powder matrix in
vacuum, at room temperature for 2 h then adding MMA,
which reduced the aggregation among ZnO NPs. This
may allow us to prepare homogenous nanocomposites
with organic matrices without additional surface
modification. The chosen average ZnO NPs size showed
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
874 Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Number vs. size distribution curve of ZnO NPs with MMA
and water
Table 2: Composition and the type of mixing used in the preparation of the solvent with the DLS measurements
Exp.
No. Composition Methodology
DLS Measurement
Run of
measurement Average size range (ìm) Zeta potential (mV)
1 0.05 g ZnONPs+ 15 mL H20
Magnetic stirring
(45 min)
I 3.50
4.01 ± 0.40
–3.45
–3.38 ± 0.81II 4.08 –2.35
III 4.48 –4.35
2 0.05 g ZnONPs+ 15 mL MMA
Magnetic stirring
(45 min)
I 2.28
2.21 ± 0.10
–3.40
–3.48 ± 0.07II 2.28 –3.57
III 2.06 –3.49
that is of vital importance for the final properties of the
new developed composite PMMA/ZnO NPs since it
affected the UV-vis absorption and thermal stability.
During the synthesis it was found that smaller ZnO NPs
showed a higher degree of agglomeration, which is in
accordance with the literature.22–24
3.3 Determination of properties
The results obtained by the DMA (Figure 4) showed
the following results. The difference in the elastic
modulus at low temperatures which may be due to a
slight decrease in the cross linking of PMMA by adding
ZnO NPs. At elevated temperatures, the incidence of
chain transfer increased (at termination). With chain
transfer, the propagating polymer radical reacts with
another molecule by proton abstraction and this leads to
branching and cross-linking.
Partly, the E modulus increases due to the addition of
ZnO NPs, but this increase is minimal. Because of the
absence of good interactions between the ZnO NPs and
the PMMA matrix, the loss factor increases with the
amount of addition of ZnO NPs in the PMMA matrix.
ZnO NPs added to the PMMA increased the E modulus
and reduced the level of cross-linking. At the same time,
the poor interaction between the filler and the matrix
results in a less elastic response, which means that the
composite of PMMA/ZnO NPs dampen the vibrations
better than pure PMMA. With the increasing of the loss
factor (tan delta) the damping of the material also
increases. In our case, the height of the loss factor peak
of the sample PMMA3 was the highest; therefore, the
damping of the vibrations of the sample PMMA3 is the
highest (Figure 5).
Characterization using the DMA method showed that
changes in the vibration-damping behaviour of the new
composite could mean that at the gingival a more gentle
feeling appears, which leads to this material being more
friendly to oral mucosa due to vibration damping. This
should also be proven through practice. Composite
PMMA/ZnO NPs would require longer curing times in
comparison with the pure PMMA. Based on this finding
it is expected to obtain a better material with the surface
modification of ZnO NPs where good interactions
between filler and matrix can be formed.
From the Flash DSC measurements (Figure 6), the
lowest glass-transition temperature was measured in the
PMMA3 sample (129.6 °C), followed by PMMA2
(131.0 °C) and the highest glass-transition temperature
was in the pure PMMA1 (131.1 °C). This fact is also
confirmed by the position of the peaks of the loss factor
(tan ), where the lowest glass-transition temperature
was measured in the PMMA3 (146.1 °C), followed by
PMMA2 (147.2°C) and the highest glass-transition
temperature in the pure PMMA1 (147.8 °C). In this
context it was found that the degree of cross-linking was
reduced by the addition of ZnO NPs. With increasing
glass-transition temperature the loss factor decreased,
which indicates that the material PMMA3 (tan  =
1.2003) has the best vibration-damping behaviour,
followed by PMMA2 (tan  = 1.1856). The pure
PMMA1 (tan  = 1.1839) showed the most elastic
response. From the height of tan  it can be concluded
that the 3 % addition of ZnO NPs in PMMA is needed to
obtain a significant change in the vibration-damping
behaviour of the composite. Despite the small differen-
ces measured, as well on DMA (E modulus, tan ) as on
Flash DSC (glass-transition temperatures), all the results
had the same tendency.
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878 875
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 6: Results of flash DSC analysis for all samples, heating
1.000 °C/s (diagram heat flow (mW) – T (°C))
Figure 4: DMA results- diagram E – modulus (N/mm2) – T (°C) for
all samples
Figure 5: DMA results-loss factor vs. temperature for all samples
The difference in the E modulus at low temperatures
may have been due to the slight decrease in the cross-
linking of the PMMA by adding ZnO NPs. The degree of
cross linking by adding ZnO NPs decreased, as shown
by the measurements on the Flash DSC and the level of
the peak of tan . This results, at temperatures up to
about 45 °C, in a minimum E modulus in the PMMA3
and a maximum E modulus in the PMMA2 (the differen-
ce between the glassy transition of PMMA2 and pure
PMMA1 was minimal). Partly, the E modulus increased
due to the addition of ZnO NPs, but this increase was
minimal. Because of the absence of good interactions
between the ZnO NPs and the PMMA matrix, the loss
factor increased with the amount of addition of ZnO NPs
in the PMMA matrix.
3.4 Wear resistance
The mean vertical dimension of each tooth sample
was measured (from the top of the field where the wear
test was performed to the bottom) before the artificial
chewing action and after. All the samples were subjected
to artificial thermal ageing (the samples were in warm
water at 70 °C for one month) and, after that, to mecha-
nical ageing on the chewing simulator. The results of
mean height loss of the control (PMMA1) and tested
group (PMMA2, PMMA3) are shown in Figure 7.
Recently, a study showed that the median vertical wear
of polymer denture teeth, made by different materials,
has been reported to be above 0.2 mm after 2 years of
observation and, in over 50 % of these, variability of
wear. This could be attributed to specific patient factors
such as biting force, nutrition habits and other unknown
factors. Gender differences were also found in the spatial
and temporal parameters of masticatory movement path
and rhythm.25–27 On the other hand, we used the same
material and improved these mechanical properties by
adding a different percentage of ZnO NPs; we also
checked the height loss of the material only after one
year of mechanical ageing on a chewing simulator.
The tested PMMA2/3 samples showed a higher wear
resistance than the tooth sample PMMA1 by 4 times.
Artificial thermal ageing had an effect on the pure
PMMA1 and PMMA2, but there was no effect in the
group with samples made of PMMA3. There was no
statistically significant difference in the loss of height
between samples made of PMMA2 with PMMA3.
Occlusal wear values of the samples made of pure
PMMA1 and PMMA2 after thermal artificial ageing
were 2-times as big compared to the samples made from
PMMA3. Wear resistance of restorative materials under
clinical conditions is a rather complicated phenomenon
compared to other mechanical and physical properties of
materials.25 Furthermore, the cause of thermal stability
and its effect on the wear resistance of both PMMA2 and
PMMA3 samples could be linked significantly. Long-
term polymerization at 70 °C, after polymerization at the
usual 100 °C for 1 h and 45 min can be recommended,
because of its positive effect on the wear resistance of
PMMA1 and PMMA2 after artificial ageing (70 °C, one
month). PMMA3 had lower height loss before, as well as
after, thermal artificial ageing compared to the other
samples.
3.5 Microstructure
The microstructure of the fracture surface of the
newly developed PMMA2 is shown in micrographs
(Figure 8), where the visible fracture is brittle. Detailed
observation by higher magnification revealed that ZnO
NPs (coloured in white) are distributed homogeneously
through the PMMA matrix with some small evidence of
ZnO NPs’ agglomeration (about 1 μm) and signs of
de-polymerized dark fields around them. This led to the
deteriorated final properties of the composite. In the
future, it is necessary to avoid the de-polymerization pro-
cess by the adequate preparation of ZnO NPs’ surface
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
876 Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: Micrographs of characteristic surface fracture in PMMA2
Figure 7: Loss of height of resin teeth after wear test on chewing
simulator before and after thermal artificial aging
modification.23 On the other hand, from the literature it is
known that homogeneously dispersed ZnO NPs in the
PMMA matrix can be explained by the theory of the
different kinds of integration by grafting copolymer
chains.28,29
4 CONCLUSIONS
Based on the methodology applied in this study, and
by considering the obtained results, the following con-
clusions can be drawn:
• The newly developed composite of PMMA/ZnO NPs
has a better vibrations damping effect than pure
PMMA1. This could lead to the gingival’s more
gentle feeling of appearance.
• Composite resin denture PMMA teeth reinforced by
the ZnO NPs showed better wear resistance by about
4 times compared to the pure PMMA1.
• The microstructure of PMMA/ZnO NPs consists of a
PMMA matrix and homogenously distributed ZnO
NPs.
• Combinations of different characterization techniques
in the designing of polymer composites reinforced by
nanoparticles with in vitro chewing simulation enable
the determination of functional composite’s behavior
in dentistry.
Conflicts of interests
There are no conflicts of interests to declare.
Acknowledgment
This study was supported by the research project
"Development of PMMA composite enriched/enhanced
with nanoparticles (Ag, ZnO) and biocompatibility
testing", number of project 451-03-2802-IP Type 1/144
and by the Infrastructure Programmes I0-0046 and
I0-0029 financed by the Slovenian Agency ARRS.
Thanks to the Ministry of Education, Science and Sport,
Republic of Slovenia (Programme MARTINA, OP20.
00369), which enabled the research with co-financing.
Special acknowledgements go to Mohammed Shariq
for work on DLS.
Note:
The responsible translator for the English language is
mag. Shelagh Hedges, Faculty of Mechanical Engineer-
ing, University of Maribor, Slovenia.
Abbreviations:
PMMA – Poly-(Methyl-Methacrylate)(powder)
MMA – Methyl-Methacrylate(liquid)
ZnO NPs – Zinc-Oxide Nanoparticles
DMA – Dynamic mechanical analysis
PMMA2 – PMMA + 2 vol. %ZnO NPs
PMMA3 – PMMA + 3 vol. %ZnO NPs
Flash DSC – Flash differential scanning calorimeter
DLS – Dynamic light scattering
TEM – Transmission Electron Microscope
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people.gov/Publications, 21–18 to 21–19, 25.11.2008
2 F. Müller, M. Naharro, G. E. Carlsson, What are prevalence and
incidence of tooth loss in the adult and elderly population in
Europe?, Clin. Oral Implants Res., 18 (2007) 3, 2–14, doi:10.1111/
j.1600-0501.2007.01459.x
3 D. A. Felton, Edentulism and comorbid factors, J. Prosthodont., 18
(2009) 2, 86–88, doi:10.1111/j.1532-849X.2009.00437.x
4 J. Zeng, Y. Sato, C. Ohkubo, T. Hosoi, In vitro wear resistance of
three types of composite resin denture teeth, J. Prosthet. Dent., 94
(2005) 5, 453–457, doi:10.1016/j.prosdent.2005.08.010
5 M. Risa, Y. Yoshikazu, M. Masakazu , O. Chikahiro, The influence
of vertical dimension of occlusion changes on the electroencephalo-
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aureus to ZnO/PVC nanocomposites, Int. J. Nanomed., 8 (2013),
1177–1184, doi:10.2147/IJN.S42010
9 I. Bilecka, M. Niederberger, New developments in the nonaqueous
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10 K. Memarzadeh, M. Vargas, J. Huang, J. Fan, R. P. Allaker, Nano
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11 M. Bitenc, Z. Crnjak Orel, Synthesis and characterization of
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12 Y. Li, J. Z. Zhang, Hydrogen generation from photoelectron chemical
water splitting based on nanomaterials, Laser Photonics Rev., 4
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13 C. Buzea, I. Pacheco, K. Robbie , Nanomaterials and nanoparticles:
Sources and toxicity, Biointerphases, 2 (2007) 4, 17–71, doi:10.1116/
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14 G. R. Patzke, Y. Zhou, R. Kontic, F. Conrad, Oxide nanomaterials:
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16 M. Cierech, J. Wojnarowicz, D. Szmigiel, B. B¹czkowski, A. M.
Grudniak, K. I. Wolska, Preparation and characterization of
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17 P. Patil, G. Gaikwad, D. R. Patil, J. Naik, Synthesis of 1-D ZnOnano-
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18 M. Demir, M. Memesa, P. Castignolles, G. Wegner, PMMA/Zinc
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D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878 877
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
The difference in the E modulus at low temperatures
may have been due to the slight decrease in the cross-
linking of the PMMA by adding ZnO NPs. The degree of
cross linking by adding ZnO NPs decreased, as shown
by the measurements on the Flash DSC and the level of
the peak of tan . This results, at temperatures up to
about 45 °C, in a minimum E modulus in the PMMA3
and a maximum E modulus in the PMMA2 (the differen-
ce between the glassy transition of PMMA2 and pure
PMMA1 was minimal). Partly, the E modulus increased
due to the addition of ZnO NPs, but this increase was
minimal. Because of the absence of good interactions
between the ZnO NPs and the PMMA matrix, the loss
factor increased with the amount of addition of ZnO NPs
in the PMMA matrix.
3.4 Wear resistance
The mean vertical dimension of each tooth sample
was measured (from the top of the field where the wear
test was performed to the bottom) before the artificial
chewing action and after. All the samples were subjected
to artificial thermal ageing (the samples were in warm
water at 70 °C for one month) and, after that, to mecha-
nical ageing on the chewing simulator. The results of
mean height loss of the control (PMMA1) and tested
group (PMMA2, PMMA3) are shown in Figure 7.
Recently, a study showed that the median vertical wear
of polymer denture teeth, made by different materials,
has been reported to be above 0.2 mm after 2 years of
observation and, in over 50 % of these, variability of
wear. This could be attributed to specific patient factors
such as biting force, nutrition habits and other unknown
factors. Gender differences were also found in the spatial
and temporal parameters of masticatory movement path
and rhythm.25–27 On the other hand, we used the same
material and improved these mechanical properties by
adding a different percentage of ZnO NPs; we also
checked the height loss of the material only after one
year of mechanical ageing on a chewing simulator.
The tested PMMA2/3 samples showed a higher wear
resistance than the tooth sample PMMA1 by 4 times.
Artificial thermal ageing had an effect on the pure
PMMA1 and PMMA2, but there was no effect in the
group with samples made of PMMA3. There was no
statistically significant difference in the loss of height
between samples made of PMMA2 with PMMA3.
Occlusal wear values of the samples made of pure
PMMA1 and PMMA2 after thermal artificial ageing
were 2-times as big compared to the samples made from
PMMA3. Wear resistance of restorative materials under
clinical conditions is a rather complicated phenomenon
compared to other mechanical and physical properties of
materials.25 Furthermore, the cause of thermal stability
and its effect on the wear resistance of both PMMA2 and
PMMA3 samples could be linked significantly. Long-
term polymerization at 70 °C, after polymerization at the
usual 100 °C for 1 h and 45 min can be recommended,
because of its positive effect on the wear resistance of
PMMA1 and PMMA2 after artificial ageing (70 °C, one
month). PMMA3 had lower height loss before, as well as
after, thermal artificial ageing compared to the other
samples.
3.5 Microstructure
The microstructure of the fracture surface of the
newly developed PMMA2 is shown in micrographs
(Figure 8), where the visible fracture is brittle. Detailed
observation by higher magnification revealed that ZnO
NPs (coloured in white) are distributed homogeneously
through the PMMA matrix with some small evidence of
ZnO NPs’ agglomeration (about 1 μm) and signs of
de-polymerized dark fields around them. This led to the
deteriorated final properties of the composite. In the
future, it is necessary to avoid the de-polymerization pro-
cess by the adequate preparation of ZnO NPs’ surface
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
876 Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 8: Micrographs of characteristic surface fracture in PMMA2
Figure 7: Loss of height of resin teeth after wear test on chewing
simulator before and after thermal artificial aging
modification.23 On the other hand, from the literature it is
known that homogeneously dispersed ZnO NPs in the
PMMA matrix can be explained by the theory of the
different kinds of integration by grafting copolymer
chains.28,29
4 CONCLUSIONS
Based on the methodology applied in this study, and
by considering the obtained results, the following con-
clusions can be drawn:
• The newly developed composite of PMMA/ZnO NPs
has a better vibrations damping effect than pure
PMMA1. This could lead to the gingival’s more
gentle feeling of appearance.
• Composite resin denture PMMA teeth reinforced by
the ZnO NPs showed better wear resistance by about
4 times compared to the pure PMMA1.
• The microstructure of PMMA/ZnO NPs consists of a
PMMA matrix and homogenously distributed ZnO
NPs.
• Combinations of different characterization techniques
in the designing of polymer composites reinforced by
nanoparticles with in vitro chewing simulation enable
the determination of functional composite’s behavior
in dentistry.
Conflicts of interests
There are no conflicts of interests to declare.
Acknowledgment
This study was supported by the research project
"Development of PMMA composite enriched/enhanced
with nanoparticles (Ag, ZnO) and biocompatibility
testing", number of project 451-03-2802-IP Type 1/144
and by the Infrastructure Programmes I0-0046 and
I0-0029 financed by the Slovenian Agency ARRS.
Thanks to the Ministry of Education, Science and Sport,
Republic of Slovenia (Programme MARTINA, OP20.
00369), which enabled the research with co-financing.
Special acknowledgements go to Mohammed Shariq
for work on DLS.
Note:
The responsible translator for the English language is
mag. Shelagh Hedges, Faculty of Mechanical Engineer-
ing, University of Maribor, Slovenia.
Abbreviations:
PMMA – Poly-(Methyl-Methacrylate)(powder)
MMA – Methyl-Methacrylate(liquid)
ZnO NPs – Zinc-Oxide Nanoparticles
DMA – Dynamic mechanical analysis
PMMA2 – PMMA + 2 vol. %ZnO NPs
PMMA3 – PMMA + 3 vol. %ZnO NPs
Flash DSC – Flash differential scanning calorimeter
DLS – Dynamic light scattering
TEM – Transmission Electron Microscope
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cedia 9 (2014) 13–17, doi:10.1016/j.apcbee.2014.01.003
20 A. Mahomed, D. W. L. Hukins, S. N. Kukureka, Effect of accelerated
aging on the viscoelastic properties of Elast-Eon™: A polyurethane
with soft poly(dimethylsiloxane) and poly(hexamethylene oxide)
segments, Mat. Sci. Eng. C, 30 (2010) 8, 1298–1303, doi:10.1016/
j.msec.2010.07.014
21 P. Spasojevic, B. Adnadjevic, S. Velickovic, J. Jovanovic, Influence
of microwave heating on the polymerization kinetics and application
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(2011) 6, 3598–3606, doi:10.1002/app.33041
22 C. Feldmann, Polyol mediated synthesis of oxide particle suspen-
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2193–2196, doi:10.1016/S1359-6462(01)00902-2
23 R. Y. Hong, J. Z. Qian, J. X. Cao, Synthesis and characterization of
PMMA grafted ZnO nanoparticles, Powder Technol., 163 (2006),
160–168, doi:10.1016/j.powtec.2006.01.015
24 A. An`lovar, A. Crnjak, M. @igon, Poly(methyl methacrylate) com-
posites prepared by in situ polymerization using organophillicnano-
to-submicrometer zinc oxide particles, Eur. Polym. J., 46 (2010) 6,
1216–1224, doi:10.1016/j.eurpolymj.2010.03.010
25 M. Ghazal, B. Yang, K. Ludwig, M. Kern, Two-body wear of resin
and ceramic denture teeth in comparison to human enamel, Dent.
Mater., 24 (2008) 4, 502–507, doi:10.1016/j.dental.2007.04.012
26 S. D. Heintze, G. Zellweger, S. Sbicego, V. Rousson, C. Muñoz-
Viveros, T. Stober, Wear of two denture teeth materials in vivo-
2-year results, Dent. Mater., 29 (2013) 9, 191–204, doi:10.1016/
j.dental.2013.04.012
27 T. Kyoko, S. Hiroshi, Gender differences in masticatory movement
path and rhythm in dentate adults, J. Prosthodont. Res., 58 (2014) 4,
237–242, doi:10.1016/j.jpor.2014.06.001
28 E. Thang, G. Cheng, X. Ma, Preparation of nano-ZnO/PMMA
composite particles via grafting of the copolymer onto the surface of
zinc oxide nanoparticle, Powder Technol., 161 (2006) 3, 209–214,
doi:10.1016/j.powtec.2005.10.007
29 A. Anzlovar, Z. Crnjak Orel, M. Zigon, Sub micrometer and
nano-ZnO as filler in PMMA materials, Mater. Tehnol., 45 (2011) 3,
269–274
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
878 Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
ERRATUM
In Materiali in tehnologije/Materials and Technology 50 (2016) 5, 797-804, doi: 10.17222/mit.2015.307
in the article entitled:
Measurement of bio impedance on an isolated rat sciatic nerve elicited with specific current
stimulating pulses – Meritev bioimpedance na izoliranem `ivcu ischiadicusu podgane vzbujene s
posebnimi tokovnimi stimulacijskimi impulzi
written by Janez Rozman1,3, Monika C. @u`ek2, Robert Frange`2, Samo Ribari~3
1Center for Implantable Technology and Sensors, ITIS d. o. o., Lepi pot 11, 1000 Ljubljana, Slovenia
2Institute of Physiology, Pharmacology and Toxicology, Veterinary Faculty, University of Ljubljana, Gerbi~eva 60,
1000 Ljubljana, Slovenia
3 Institute of Pathophysiology, Zalo{ka 4, Medical Faculty, University of Ljubljana, Republic of Slovenia
two additional authors are missing. The correct authorship of this article is:
Robert Brajkovi~1, Dejan Kri`aj1, Janez Rozman2,3, Monika C. @u`ek4, Robert Frange`4, Samo Ribari~3
1Faculty of Electrical Engineering, University of Ljubljana, Tr`a{ka 25, 1000 Ljubljana,
2Center for Implantable Technology and Sensors, ITIS d. o. o. Ljubljana, Lepi pot 11, 1000 Ljubljana,
3Institute of Pathophysiology, Medical Faculty, University of Ljubljana, Vrazov trg 2, 1000 Ljubljana,
4Institute of Physiology, Pharmacology and Toxicology, Veterinary Faculty, University of Ljubljana, Gerbi~eva 60,
1000 Ljubljana, Republic of Slovenia
MIT Editorial
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
19 M. Roman, Particle Size and Zeta Potential of ZnO, APCBEE Pro-
cedia 9 (2014) 13–17, doi:10.1016/j.apcbee.2014.01.003
20 A. Mahomed, D. W. L. Hukins, S. N. Kukureka, Effect of accelerated
aging on the viscoelastic properties of Elast-Eon™: A polyurethane
with soft poly(dimethylsiloxane) and poly(hexamethylene oxide)
segments, Mat. Sci. Eng. C, 30 (2010) 8, 1298–1303, doi:10.1016/
j.msec.2010.07.014
21 P. Spasojevic, B. Adnadjevic, S. Velickovic, J. Jovanovic, Influence
of microwave heating on the polymerization kinetics and application
properties of the PMMA dental materials, J. Appl. Polym. Sci., 119
(2011) 6, 3598–3606, doi:10.1002/app.33041
22 C. Feldmann, Polyol mediated synthesis of oxide particle suspen-
sions and their application, Scripta Mater., 44 (2001) 8–9,
2193–2196, doi:10.1016/S1359-6462(01)00902-2
23 R. Y. Hong, J. Z. Qian, J. X. Cao, Synthesis and characterization of
PMMA grafted ZnO nanoparticles, Powder Technol., 163 (2006),
160–168, doi:10.1016/j.powtec.2006.01.015
24 A. An`lovar, A. Crnjak, M. @igon, Poly(methyl methacrylate) com-
posites prepared by in situ polymerization using organophillicnano-
to-submicrometer zinc oxide particles, Eur. Polym. J., 46 (2010) 6,
1216–1224, doi:10.1016/j.eurpolymj.2010.03.010
25 M. Ghazal, B. Yang, K. Ludwig, M. Kern, Two-body wear of resin
and ceramic denture teeth in comparison to human enamel, Dent.
Mater., 24 (2008) 4, 502–507, doi:10.1016/j.dental.2007.04.012
26 S. D. Heintze, G. Zellweger, S. Sbicego, V. Rousson, C. Muñoz-
Viveros, T. Stober, Wear of two denture teeth materials in vivo-
2-year results, Dent. Mater., 29 (2013) 9, 191–204, doi:10.1016/
j.dental.2013.04.012
27 T. Kyoko, S. Hiroshi, Gender differences in masticatory movement
path and rhythm in dentate adults, J. Prosthodont. Res., 58 (2014) 4,
237–242, doi:10.1016/j.jpor.2014.06.001
28 E. Thang, G. Cheng, X. Ma, Preparation of nano-ZnO/PMMA
composite particles via grafting of the copolymer onto the surface of
zinc oxide nanoparticle, Powder Technol., 161 (2006) 3, 209–214,
doi:10.1016/j.powtec.2005.10.007
29 A. Anzlovar, Z. Crnjak Orel, M. Zigon, Sub micrometer and
nano-ZnO as filler in PMMA materials, Mater. Tehnol., 45 (2011) 3,
269–274
D. POPOVI] et al.: SYNTHESIS OF PMMA/ZnO NANOPARTICLES COMPOSITE USED FOR RESIN TEETH
878 Materiali in tehnologije / Materials and technology 51 (2017) 5, 871–878
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
ERRATUM
In Materiali in tehnologije/Materials and Technology 50 (2016) 5, 797-804, doi: 10.17222/mit.2015.307
in the article entitled:
Measurement of bio impedance on an isolated rat sciatic nerve elicited with specific current
stimulating pulses – Meritev bioimpedance na izoliranem `ivcu ischiadicusu podgane vzbujene s
posebnimi tokovnimi stimulacijskimi impulzi
written by Janez Rozman1,3, Monika C. @u`ek2, Robert Frange`2, Samo Ribari~3
1Center for Implantable Technology and Sensors, ITIS d. o. o., Lepi pot 11, 1000 Ljubljana, Slovenia
2Institute of Physiology, Pharmacology and Toxicology, Veterinary Faculty, University of Ljubljana, Gerbi~eva 60,
1000 Ljubljana, Slovenia
3 Institute of Pathophysiology, Zalo{ka 4, Medical Faculty, University of Ljubljana, Republic of Slovenia
two additional authors are missing. The correct authorship of this article is:
Robert Brajkovi~1, Dejan Kri`aj1, Janez Rozman2,3, Monika C. @u`ek4, Robert Frange`4, Samo Ribari~3
1Faculty of Electrical Engineering, University of Ljubljana, Tr`a{ka 25, 1000 Ljubljana,
2Center for Implantable Technology and Sensors, ITIS d. o. o. Ljubljana, Lepi pot 11, 1000 Ljubljana,
3Institute of Pathophysiology, Medical Faculty, University of Ljubljana, Vrazov trg 2, 1000 Ljubljana,
4Institute of Physiology, Pharmacology and Toxicology, Veterinary Faculty, University of Ljubljana, Gerbi~eva 60,
1000 Ljubljana, Republic of Slovenia
MIT Editorial
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2 ARTICLE PROOFS
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requested to return the proofs with any corrections
within two days. In the case of a delay the Editor will
postpone publication of the article until the article proof
is received.
3 COPYRIGHT
In addition to the paper, authors must also enclose a
written statement that the paper is original work and has
not been published in this form anywhere else and that it
is not under consideration for publication elsewhere.
On publication, copyright will pass to the publisher.
The Journal of Materials and Technology must be stated
as the source in all later publications.
The Editorial Board of the Journal of Materials and
Technology:
• decide whether to accept a paper for publication;
• obtain professional reviewers for papers and decide
on any proposals to shorten or extend them;
• obtain correct terminology and edit language.
The e-files of papers will be kept in the archives of
the Materials and Technology journal.
4 PUBLICATION FEE
Authors will be asked to pay a publication fee for
each article before its publishing in MIT journal. After
the article has been accepted for publishing, this fee
needs to be paid. The fee is  300 for regular articles,
and  150 for articles, presented at ICMT annual
conference in Portoro`. The additional cost for coloured
printing for one page is  80. These fees do include
value-added tax (VAT).
 
MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY 
Lepi pot 11, 1000 Ljubljana, Slovenia 
Phone: +386 1 4701 860, +386 1 4701 857, Fax: +386 01 4701 939, e-mail: mit@imt.si 
Editor-in-chief: dr. Paul John McGuiness 
----------------------------------------------------------------------------------------------------------------- 
 
 
MIT TEMPLATE 
 
 
TITLE 
(The fewest possible words that adequately describe the contents of the paper in English) 
NASLOV  
(v slovenskem jeziku, for Slovenian authors only) 
 
Author`s Name Author`s Surname1, Author`s Name Author`s Surname2, Author`s Name 
Author`s Surname3* 
1Author`s Affiliation, Address, City, State 
2Author`s Affiliation, Address, City, State  
3Author`s Affiliation, Address, City, State 
*Corresponding author`s e-mail  
 
Abstract (in English) 
The abstract must be viewed as a miniversion of the paper and should not exceed 250 words. It 
must be written in `past tense` , because it refers to work already done. It must state the principal 
objectives and scope of the investigation, describe the methodology employed, summarize the 
results and state the principal conclusions.  
It must not give information or conclusions that are not inculded in the paper and it must not cite 
references to the literature. 
Keywords (in English) 
Up to 3-4 words (abbreviations of the expressions are not allowed, only the common used ones: 
SEM, XRD, XPS, etc.) 
 
References to unpublished or not readily accessible
reports must be avoided.
References must be cited in English.
All references should have a DOI number, if it exists
for the given reference.
Authors should avoid self-citing in the references as
far as possible.
In the list of references, monographs, articles in
journals, journals, contributions to conference pro-
ceedings, patent documents, electronic monographs,
articles and other electronic documents must be cited in
accordance with the following examples:
1. Monographs
H. Ibach, H. Luth, Solid State physics, 2nd ed.,
Springer, Berlin 1991, 245
2. Articles in journals
T. Mauder, J. Stetina, Improvement of the casting of
special steel with a wide solid liquid interface, Mater.
Tehnol., 50 (2015) 1–2, doi:10.17222/mit.2014.122
3. Contributions to conference proceedings, sym-
posiums or conferences
I. Rak, M. Kocak, V. Gliha, N. Gubeljak: Fracture
behaviour of over-matched high strength steel welds
containing soft root layers, Proc. of the 2nd Inter.
Symp. on Mis-Matching of Interfaces and Welds,
Reinsford, 1997, 627–641
4. Contributions in electronic form/online
• Articles: M. P. Wnuk: Principles of fracture
mechanics for space applications, http://mit.imt.si/
Revija/izvodi/kzt996/wnuk.pdf, 30.01.2000
• Other: http://www.imt.si/, 15.03.2016
5. Standards
ISO 15787:2001(E) Technical product documen-
tation, Heat treated ferrous part, Presentation and
indication ISO Committee, Geneve
SPECIAL NOTICE TO AUTHORS
In the past few years we have received an increasing
number of articles that include computer simulations,
with 3D colour images as a result. As the journal charges
80 per page for colour printing, some authors have
decided to convert their colour images to grey-scale
images. We believe that this lessens both the impact and
the clarity of the results. Since the Materials and
Technology journal consistently strives for better quality
presentation and a higher impact factor, we will be
rejecting articles of this type.
2 ARTICLE PROOFS
Authors will receive a set of proofs. They are
requested to return the proofs with any corrections
within two days. In the case of a delay the Editor will
postpone publication of the article until the article proof
is received.
3 COPYRIGHT
In addition to the paper, authors must also enclose a
written statement that the paper is original work and has
not been published in this form anywhere else and that it
is not under consideration for publication elsewhere.
On publication, copyright will pass to the publisher.
The Journal of Materials and Technology must be stated
as the source in all later publications.
The Editorial Board of the Journal of Materials and
Technology:
• decide whether to accept a paper for publication;
• obtain professional reviewers for papers and decide
on any proposals to shorten or extend them;
• obtain correct terminology and edit language.
The e-files of papers will be kept in the archives of
the Materials and Technology journal.
4 PUBLICATION FEE
Authors will be asked to pay a publication fee for
each article before its publishing in MIT journal. After
the article has been accepted for publishing, this fee
needs to be paid. The fee is  300 for regular articles,
and  150 for articles, presented at ICMT annual
conference in Portoro`. The additional cost for coloured
printing for one page is  80. These fees do include
value-added tax (VAT).
 
MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY 
Lepi pot 11, 1000 Ljubljana, Slovenia 
Phone: +386 1 4701 860, +386 1 4701 857, Fax: +386 01 4701 939, e-mail: mit@imt.si 
Editor-in-chief: dr. Paul John McGuiness 
----------------------------------------------------------------------------------------------------------------- 
 
 
MIT TEMPLATE 
 
 
TITLE 
(The fewest possible words that adequately describe the contents of the paper in English) 
NASLOV  
(v slovenskem jeziku, for Slovenian authors only) 
 
Author`s Name Author`s Surname1, Author`s Name Author`s Surname2, Author`s Name 
Author`s Surname3* 
1Author`s Affiliation, Address, City, State 
2Author`s Affiliation, Address, City, State  
3Author`s Affiliation, Address, City, State 
*Corresponding author`s e-mail  
 
Abstract (in English) 
The abstract must be viewed as a miniversion of the paper and should not exceed 250 words. It 
must be written in `past tense` , because it refers to work already done. It must state the principal 
objectives and scope of the investigation, describe the methodology employed, summarize the 
results and state the principal conclusions.  
It must not give information or conclusions that are not inculded in the paper and it must not cite 
references to the literature. 
Keywords (in English) 
Up to 3-4 words (abbreviations of the expressions are not allowed, only the common used ones: 
SEM, XRD, XPS, etc.) 
 
MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY 
Lepi pot 11, 1000 Ljubljana, Slovenia 
Phone: +386 1 4701 860, +386 1 4701 857, Fax: +386 01 4701 939, e-mail: mit@imt.si
Editor-in-chief: dr. Paul John McGuiness 
----------------------------------------------------------------------------------------------------------------- 
Povzetek (v slovenskem jeziku, for Slovenian authors only) 
Do 250 besed  
Ključne besede (v slovenskem jeziku, for Slovenian authors only)
Od 3 do 4 besede (okrajšave izrazov niso dovoljene, razen splošno sprejetih: SEM, XRD, XPS, 
itd.) 
1 INTRODUCTION  
The purpose of the introduction is to supply sufficient background information to allow the 
reader to understand and evaluate the results of the present study. It should state briefly and 
clearly the purpose in writing paper, present as clearly as possible the nature and scope of the 
problem investigated, review recent literature to orient the reader, state the method of the 
investigation, and if necessary, the reasons for the choice of a particular method. 
2 EXPERIMENTAL PART
The experimental part must give full details of the experimental aparatus and the methods used in 
obtaining results. It must include accurate details relating to the equipment and materials used, 
incuding quantities, temperatures, times etc. It must provide sufficient information for a collegaue 
in the same field to reproduce the experiment. The data must be presented clearly and concisely, 
using graphs and tables. Figures and Tables must be written bold and not abbreviated throughout 
text. Figures and Tables must also follow the numerical order of appearance in the text. Once 
mentioned, tehy can be repeated. Equations must be written with Equation Editor and must be 
marked on the right-hand side of the text with numbers in round brackets. Throughout the text  
they must not be abbreviated.Physical quantities in the text have to be written as a plain text.
Equation 1                                                                                                                                   (1) 
Equation 2                                                                                                                                   (2) 
3 RESULTS 
Results section should provide an overall picture of the experiments without repeating any of the 
details in the experimental section. The data should be presented clearly and concisely, using 
graphs and tables. Figures and Tables should be written bold and not abbreviated throughout text. 
MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY 
Lepi pot 11, 1000 Ljubljana, Slovenia 
Phone: +386 1 4701 860, +386 1 4701 857, Fax: +386 01 4701 939, e-mail: mit@imt.si
Editor-in-chief: dr. Paul John McGuiness 
----------------------------------------------------------------------------------------------------------------- 
Figures and Tables must also follow the numerical order of appearance in the text. Once 
mentioned they can be repeated.  
4 DISCUSSION 
The purpose of the discussion is to present the principles, relationships, and generalisations 
shown by the results. The discussion must make clear the significance of the results and compare 
the findings with previously published work. It must discuss the results not simply repeat them. 
5 CONCLUSIONS 
The conclusion section should be short. It must present one or more conclusions which have to be 
drawn from the results and the subsequent discussion. Conclusions must be concise and easily 
understood by the reader. 
Acknowledgment 
6 REFERENCES 
All references need a DOI number if it exists for the given reference.  
In the list of references monographs, articles in journals, journals, contributions to conference 
proceedings, patent documents, electronic monographs, articles and other electronic documents 
must be cited in accordance with the following examples: 
1. Monographs:
H. Ibach, H. Luth, Solid State physics, 2nd ed., Springer, Berlin 1991, 245 
2. Articles in journals:
T. Mauder, J. Stetina, Improvement of the casting of special steel with wide solid liquid 
interface, Mater. Tehnol., 50 (2015) 1, doi:10.17222/mit.2014.122 
3. Contributions to conference proceedings, symposiums:
I. Rak, M. Kocak, V. Gliha, N. Gubeljak, Fracture behaviour of over-matched high 
strength steel welds containing soft root layers, Proc. of the 2nd Inter. Symp. on Mis-
Matching of Interfaces and Welds, Reinsford 1997, 627–641
4. Contributions in electronic form/online:
Articles: M. P. Wnuk: Principles of fracture mechanics for space applications, 
http://mit.imt.si/Revija/izvodi/kzt996/wnuk.pdf, 30.01.2000 
Other: http://www.imt.si, 15.03.2016 
5. Standards:  
ISO 15787:2001(E) Technical product documentation, Heat treated ferrous part, 
Presentation and indication ISO Committee, Geneve
MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY 
Lepi pot 11, 1000 Ljubljana, Slovenia 
Phone: +386 1 4701 860, +386 1 4701 857, Fax: +386 01 4701 939, e-mail: mit@imt.si
Editor-in-chief: dr. Paul John McGuiness 
----------------------------------------------------------------------------------------------------------------- 
Povzetek (v slovenskem jeziku, for Slovenian authors only) 
Do 250 besed  
Ključne besede (v slovenskem jeziku, for Slovenian authors only)
Od 3 do 4 besede (okrajšave izrazov niso dovoljene, razen splošno sprejetih: SEM, XRD, XPS, 
itd.) 
1 INTRODUCTION  
The purpose of the introduction is to supply sufficient background information to allow the 
reader to understand and evaluate the results of the present study. It should state briefly and 
clearly the purpose in writing paper, present as clearly as possible the nature and scope of the 
problem investigated, review recent literature to orient the reader, state the method of the 
investigation, and if necessary, the reasons for the choice of a particular method. 
2 EXPERIMENTAL PART
The experimental part must give full details of the experimental aparatus and the methods used in 
obtaining results. It must include accurate details relating to the equipment and materials used, 
incuding quantities, temperatures, times etc. It must provide sufficient information for a collegaue 
in the same field to reproduce the experiment. The data must be presented clearly and concisely, 
using graphs and tables. Figures and Tables must be written bold and not abbreviated throughout 
text. Figures and Tables must also follow the numerical order of appearance in the text. Once 
mentioned, tehy can be repeated. Equations must be written with Equation Editor and must be 
marked on the right-hand side of the text with numbers in round brackets. Throughout the text  
they must not be abbreviated.Physical quantities in the text have to be written as a plain text.
Equation 1                                                                                                                                   (1) 
Equation 2                                                                                                                                   (2) 
3 RESULTS 
Results section should provide an overall picture of the experiments without repeating any of the 
details in the experimental section. The data should be presented clearly and concisely, using 
graphs and tables. Figures and Tables should be written bold and not abbreviated throughout text. 
MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY 
Lepi pot 11, 1000 Ljubljana, Slovenia 
Phone: +386 1 4701 860, +386 1 4701 857, Fax: +386 01 4701 939, e-mail: mit@imt.si
Editor-in-chief: dr. Paul John McGuiness 
----------------------------------------------------------------------------------------------------------------- 
Figures and Tables must also follow the numerical order of appearance in the text. Once 
mentioned they can be repeated.  
4 DISCUSSION 
The purpose of the discussion is to present the principles, relationships, and generalisations 
shown by the results. The discussion must make clear the significance of the results and compare 
the findings with previously published work. It must discuss the results not simply repeat them. 
5 CONCLUSIONS 
The conclusion section should be short. It must present one or more conclusions which have to be 
drawn from the results and the subsequent discussion. Conclusions must be concise and easily 
understood by the reader. 
Acknowledgment 
6 REFERENCES 
All references need a DOI number if it exists for the given reference.  
In the list of references monographs, articles in journals, journals, contributions to conference 
proceedings, patent documents, electronic monographs, articles and other electronic documents 
must be cited in accordance with the following examples: 
1. Monographs:
H. Ibach, H. Luth, Solid State physics, 2nd ed., Springer, Berlin 1991, 245 
2. Articles in journals:
T. Mauder, J. Stetina, Improvement of the casting of special steel with wide solid liquid 
interface, Mater. Tehnol., 50 (2015) 1, doi:10.17222/mit.2014.122 
3. Contributions to conference proceedings, symposiums:
I. Rak, M. Kocak, V. Gliha, N. Gubeljak, Fracture behaviour of over-matched high 
strength steel welds containing soft root layers, Proc. of the 2nd Inter. Symp. on Mis-
Matching of Interfaces and Welds, Reinsford 1997, 627–641
4. Contributions in electronic form/online:
Articles: M. P. Wnuk: Principles of fracture mechanics for space applications, 
http://mit.imt.si/Revija/izvodi/kzt996/wnuk.pdf, 30.01.2000 
Other: http://www.imt.si, 15.03.2016 
5. Standards:  
ISO 15787:2001(E) Technical product documentation, Heat treated ferrous part, 
Presentation and indication ISO Committee, Geneve
MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY 
Lepi pot 11, 1000 Ljubljana, Slovenia 
Phone: +386 1 4701 860, +386 1 4701 857, Fax: +386 01 4701 939, e-mail: mit@imt.si
Editor-in-chief: dr. Paul John McGuiness 
----------------------------------------------------------------------------------------------------------------- 
MIT CHECKLIST 
PLEASE READ CAREFULLY THE STATEMENTS WRITTEN BELOW AND FULFILL IT 
WITH MARKING THE CHECKBOX.  
THIS DOCUMENTATION IS REQUIRED TO PROCESS YOUR SUBMITTED
MANUSCRIPT, ENTITLED: __________________ FOR PUBLICATION IN MATERIALS 
AND TECHNOLOGY JOURNAL. 
Have you included a cover letter that makes clear the novelty of the investigation?  
Have you supplied the names of at least two potential referees? 
Is the manuscript written in accordance with the Instructions for Authors and with the 
MIT template? 
Does the number of figures and tables in your paper not exceed the required number 
according with Instructions for Authors?  
Are the figures and captions positioned in the text? 
Do you accept the terms of the MIT journal policy about the publication fee?  
 
Signature:________________________ 
SPECIAL NOTICE:  
From 1st October 2016 the journal Materials and Technology will be introducing a publication fee. All 
fees listed below, were agreed by the Editorial Board of Materials and Technology during its meeting 
on 15th June 2016: 
Publication fee -regular  €300, 00 
Publication fee – ICM&T conference paper €150, 00 
Additional charge  for colour printing   € 80, 00 /page 
MIT journal subrscription (Slovenia)   €42, 00 
MIT journal subscription (abroad)   €85, 00 
 
MIT (Materials and Technology journal) Editorial Office  
Editor-in-Chief: dr. Paul John McGuiness 
 
Ljubljana, September 2016 
 
OBVESTILO! 
Od 1.10.2016 dalje objava člankov v reviji Materiali in tehnologije (MIT) ni več brezplačna. Objavo 
članka Uredništvo revije MIT zaračunava na podlagi sklepa iz seje Uredniškega odbora revije Materiali 
in tehnologije/Materials and Technology, dne 15. junija 2016, po spodnjem ceniku: 
Objava prispevka  300, 00 € 
Objava konferečnega prispevka (ICM&T) 150, 00 € 
Doplačilo za barvni tisk   80, 00 €/stran 
Naročnina na revijo MIT (Slovenija)   42, 00 € 
Naročnina na revijo MIT (tujina)   85, 00 € 
 
Uredništvo revije Materiali in tehnologije (MIT) 
Glavni in odgovorni urednik: dr. Paul John McGuiness 
 
Ljubljana, september 2016 
MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY 
Lepi pot 11, 1000 Ljubljana, Slovenia 
Phone: +386 1 4701 860, +386 1 4701 857, Fax: +386 01 4701 939, e-mail: mit@imt.si
Editor-in-chief: dr. Paul John McGuiness 
----------------------------------------------------------------------------------------------------------------- 
MIT CHECKLIST 
PLEASE READ CAREFULLY THE STATEMENTS WRITTEN BELOW AND FULFILL IT 
WITH MARKING THE CHECKBOX.  
THIS DOCUMENTATION IS REQUIRED TO PROCESS YOUR SUBMITTED
MANUSCRIPT, ENTITLED: __________________ FOR PUBLICATION IN MATERIALS 
AND TECHNOLOGY JOURNAL. 
Have you included a cover letter that makes clear the novelty of the investigation?  
Have you supplied the names of at least two potential referees? 
Is the manuscript written in accordance with the Instructions for Authors and with the 
MIT template? 
Does the number of figures and tables in your paper not exceed the required number 
according with Instructions for Authors?  
Are the figures and captions positioned in the text? 
Do you accept the terms of the MIT journal policy about the publication fee?  
 
Signature:________________________ 
SPECIAL NOTICE:  
From 1st October 2016 the journal Materials and Technology will be introducing a publication fee. All 
fees listed below, were agreed by the Editorial Board of Materials and Technology during its meeting 
on 15th June 2016: 
Publication fee -regular  €300, 00 
Publication fee – ICM&T conference paper €150, 00 
Additional charge  for colour printing   € 80, 00 /page 
MIT journal subrscription (Slovenia)   €42, 00 
MIT journal subscription (abroad)   €85, 00 
 
MIT (Materials and Technology journal) Editorial Office  
Editor-in-Chief: dr. Paul John McGuiness 
 
Ljubljana, September 2016 
 
OBVESTILO! 
Od 1.10.2016 dalje objava člankov v reviji Materiali in tehnologije (MIT) ni več brezplačna. Objavo 
članka Uredništvo revije MIT zaračunava na podlagi sklepa iz seje Uredniškega odbora revije Materiali 
in tehnologije/Materials and Technology, dne 15. junija 2016, po spodnjem ceniku: 
Objava prispevka  300, 00 € 
Objava konferečnega prispevka (ICM&T) 150, 00 € 
Doplačilo za barvni tisk   80, 00 €/stran 
Naročnina na revijo MIT (Slovenija)   42, 00 € 
Naročnina na revijo MIT (tujina)   85, 00 € 
 
Uredništvo revije Materiali in tehnologije (MIT) 
Glavni in odgovorni urednik: dr. Paul John McGuiness 
 
Ljubljana, september 2016 

E L E C T R O N I C
A   C   C   E   S   S
http ://mit. imt. s i
ISSN 1580-2949
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ISSN: 1580-2949
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2017
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