VSEBINA – CONTENTS
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Boron-doped hydrogenated amorphous semiconductor MEMS
Z borom dopirani hidrogenirani amorfni polprevodnik MEMS
M. Galindo, C. Zúñiga, R. Palomino-Merino, F. López, W. Calleja, J. de la Hidalga, V. M. Castaño . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3
Wear behaviour of B4C reinforced hybrid aluminum-matrix composites
Vedenje hibridnega kompozita na osnovi aluminija, oja~anega z B4C, pri obrabi
T. Thiyagarajan, R. Subramanian, S. Dharmalingam, N. Radika, A. Gowrisankar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9
Influence of the HIP process on the properties of as-cast Ni-based alloys
Vpliv vro~ega izostatskega stiskanja na lastnosti Ni-zlitin z lito strukturo
J. Malcharcziková, M. Pohludka, V. Michenka, T. ^egan, J. Juøica, M. Kursa. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15
Corrosion of CrN-coated stainless steel in a NaCl solution (w = 3 %)
Korozija nerjavnega jekla s CrN-prevleko v raztopini NaCl (w = 3 %)
I. Kucuk, C. Sarioglu . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19
Preparation and properties of master alloys Nb-Al and Ta-Al for melting and casting of -TiAl intermetallics
Priprava in lastnosti predzlitin Nb-Al in Ta-Al za taljenje in ulivanje intermetalnih zlitin -TiAl
J. Juøica, T. ^egan, K. Skotnicová, D. Petlák, B. Smetana, V. Matìjka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27
Decreasing the carbonitride size and amount in austenitic steel with heat treatment and thermomechanical processing
Zmanj{anje velikosti karbonitridov v avstenitnem jeklu s toplotno obdelavo in termomehansko predelavo
P. Martínek, P. Podaný, J. Nacházel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31
Deep cryogenic treatment of H11 hot-working tool steel
Globoka kriogenska obdelava orodnega jekla H11 za delo v vro~em
P. Suchmann, D. Jandova, J. Niznanska . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37
Numerical prediction of the compound layer growth during the gas nitriding of Fe-M binary alloys
Numeri~no napovedovanje rasti spojinske plasti med plinskim nitriranjem binarnih zlitin Fe-M
R. Kouba, M. Keddam. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43
Antibacterial composite based on nanostructured ZnO mesoscale particles and a poly(vinyl chloride) matrix
Protibakterijski kompozit na osnovi nanostrukturnih delcev ZnO in osnove iz polivinil klorida
J. Sedlák, P. Ba`ant, J. Klofá~, M. Pastorek, I. Kuøitka. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55
Water-soluble cores – verifying development trends
Jedra, topna v vodi – preverjanje smeri razvoja
E. Adámková, P. Jelínek, J. Beòo, F. Mik{ovský . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 61
Mathematical modelling and physical simulation of the hot plastic deformation and recrystallization of steel with
micro-additives
Matemati~no modeliranje in fizikalna simulacija vro~e plasti~ne predelave in rekristalizacije jekla z mikrododatki
E. Kalinowska-Ozgowicz, W. Wajda, W. Ozgowicz . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69
Synthesis of NiTi/Ni-TiO2 composite nanoparticles via ultrasonic spray pyrolysis
Sinteza kompozitnih nanodelcev NiTi/Ni-TiO2 z ultrazvo~no razpr{ilno pirolizo
P. Majeri~, R. Rudolf, I. An`el, J. Bogovi}, S. Stopi}, B. Friedrich . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75
Hydroxyapatite coatings on Cp-Titanium Grade-2 surfaces prepared with plasma spraying
Nanos hidroksiapatita na povr{ino Cp-Titana Grade-2 z nabrizgavanjem s plazmo
R. Rudolf, D. Stamenkovi}, Z. Aleksi}, M. Jenko, I. \ordevi}, A. Todorovi}, V. Jokanovi}, K. T. Rai} . . . . . . . . . . . . . . . . . . . . . . . . . 81
Heat-transfer characteristics of a non-Newtonian Au nanofluid in a cubical enclosure with differentially heated side walls
Zna~ilnosti prenosa toplote nenewtonske Au nanoteko~ine v kockastem ohi{ju z razli~no gretima stranskima stenama
P. Ternik, R. Rudolf, Z. @uni~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 87
Amplitude–frequency response of an aluminium cantilever beam determined with piezoelectric transducers
Amplitudno-frekven~ni odziv konzolnega nosilca iz aluminija, ugotovljen s piezoelektri~nimi pretvorniki
Z. La{ová, R. Zem~ík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 95
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 49(1)1–177(2015)
MATER.
TEHNOL.
LETNIK
VOLUME 49
[TEV.
NO. 1
STR.
P. 1–177
LJUBLJANA
SLOVENIJA
JAN.–FEB.
2015
Micromechanical model of the substituents of a unidirectional fiber-reinforced composite and its response to the tensile cyclic
loading
Mikromehanski model nadomestkov kompozitov, oja~anih z enosmernimi vlakni, in njihov odgovor na cikli~no natezno obremenjevanje
T. Kroupa, H. Srbová, R. Zem~ík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99
Dry-cutting options with a chainsaw at the Hotavlje I natural-stone quarry
Mo`nosti suhega rezanja z veri`no `ago v kamnolomu naravnega kamna Hotavlje I.
J. Kortnik, B. Markoli . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 103
Effect of sliding speed on the frictional behavior and wear performance of borided and plasma-nitrided W9Mo3Cr4V
high-speed steel
Vpliv hitrosti drsenja na vedenje in obrabo boriranega in v plazmi nitriranega hitroreznega jekla W9Mo3Cr4V
I. Gunes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111
Tribological behaviour of A356/10SiC/3Gr hybrid composite in dry-sliding conditions
Tribolo{ko vedenje hibridnega kompozita A356/10SiC/3Gr pri suhem drsenju
B. Stojanovi}, M. Babi}, N. Miloradovi}, S. Mitrovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 117
New solid-polymer-electrolyte material for dye-sensitized solar cells
Novi elektrolitni material na osnovi trdnega polimera za son~ne celice, ob~utljive za svetlobo
V. K. Singh, B. Bhattacharya, S. Shukla, P. K. Singh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 123
Compacting the powder of Al-Cr-Mn alloy with SPS
Kompaktiranje prahu zlitine Al-Cr-Mn s SPS
T. F. Kubatík, Z. Pala, P. Novák. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 129
Effect of tempering on the microstructure and mechanical properties of resistance-spot-welded DP980 dual-phase steel
Vpliv popu{~anja na mikrostrukturo in mehanske lastnosti to~kasto varjenega dvofaznega jekla DP980
F. Nikoosohbat, S. Kheirandish, M. Goodarzi, M. Pouranvari . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 133
Optimization of the process parameters for surface roughness and tool life in face milling using the Taguchi analysis
Optimizacija procesnih parametrov glede na hrapavost povr{ine in trajnostno dobo orodja pri ~elnem rezkanju z uporabo Taguchijeve analize
M. Sarýkaya, H. Dilipak, A. Gezgin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 139
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Thin tin monosulfide films deposited with the HVE method for photovoltaic applications
Tanka plast HVE kositrovega monosulfida za uporabo v fotovoltaiki
N. Revathi, S. Bereznev, J. Lehner, R. Traksmaa, M. Safonova, E. Mellikov, O. Volobujeva . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 149
Mechanical properties of the austenitic stainless steel X15CrNiSi20-12 after recycling
Mehanske lastnosti avstenitnega nerjavnega jekla X15CrNiSi20-12 po recikliranju
A. Deli}, O. Kablar, A. Zuli}, D. Kova~evi}, N. Hod`i}, E. Bar~i}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153
Comparison of refractory coatings based on talc, cordierite, zircon and mullite fillers for lost-foam casting
Primerjava ognjevzdr`nih premazov na osnovi smukca, kordierita, cirkona in mulitnih polnil za ulivanje v forme z izparljivim modelom
Z. A}imovi}-Pavlovi}, A. Terzi}, Lj. Andri}, M. Pavlovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157
Selection of the most appropriate welding technology for hardfacing of bucket teeth
Izbira najbolj primerne tehnologije trdega navarjanja zoba zajemalke
V. Lazi}, A. Sedmak, R. R. Nikoli}, M. Mutavd`i}, S. Aleksandrovi}, B. Krsti}, D. Milosavljevi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 165
Characterization of TiO2 nanoparticles with high-resolution FEG scanning electron microscopy
Karakterizacija nanodelcev TiO2 z visokolo~ljivostno FEG vrsti~no elektronsko mikroskopijo
Z. Samard`ija, D. Lapornik, K. Gradi{ek, D. Verhov{ek, M. ^eh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 173
Erratum. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 177
M. GALINDO et al.: BORON-DOPED HYDROGENATED AMORPHOUS SEMICONDUCTOR MEMS
BORON-DOPED HYDROGENATED AMORPHOUS
SEMICONDUCTOR MEMS
Z BOROM DOPIRANI HIDROGENIRANI AMORFNI
POLPREVODNIK MEMS
Margarita Galindo1, Carlos Zúñiga2, Rodolfo Palomino-Merino1,
Francisco López3, Wilfrido Calleja2, Javier de la Hidalga2, Victor M. Castaño4
1Facultad de Ciencias Físico-Matemáticas, Benemérita Universidad Autónoma de Puebla, Av. San Claudio y 18 Sur, Col. San Manuel, Ciudad
Universitaria, 72570 Puebla, Mexico
2Instituto Nacional de Astrofísica, Óptica y Electrónica (INAOE), Luis E Erro # 1, Santa María Tonantzintla, Puebla, Mexico
3Facultad de Ciencias de la Electrónica, Benemérita Universidad Autónoma de Puebla, Av. San Claudio y 18 Sur, Col. San Manuel, Ciudad
Universitaria, 72570 Puebla, Mexico
4Centro de Fisica Aplicada y Tecnologia Avanzada, Campus Juriquilla, Queretaro, Mexico (on sabbatical leave at CIATEQ-Queretaro)
meneses@unam.mx
Prejem rokopisa – received: 2013-01-28; sprejem za objavo – accepted for publication: 2014-01-17
A micromachining process for the fabrication of micro-electro-mechanical systems (MEMSs), using both boron-doped silicon
and silicon-germanium amorphous films (a-SiB:H and a-Si0.50Ge0.50B:H) prepared with plasma-enhanced chemical vapor
deposition (PECVD), at 110 kHz, 300 °C and a low pressure (8 · 10–4 mbar) is presented. These MEMSs were fabricated using a
surface micromachining technology with wet and dry etching. The microstructures made of a-Si0.50Ge0.50B:H and a-SiB:H,
with a structural single layer with a thickness of 1 μm were fabricated. Diamond, ring, V and Vernier-beam structures were
included in the MEMS process. We found optical gaps of 1.55 eV and 0.97 eV for the a-SiB:H and a-Si0.50Ge0.50B:H films,
respectively. We varied the boron doping concentration for the a-Si0.50Ge0.50B:H samples, causing a decrease in the resistance
from 6.84 M to 0.2 M. A compressive stress of 13.44 MPa was obtained for the diamond-type microstructures of a-SiB:H.
An improved mechanical stability and lateral definition were obtained for the a-Si0.50Ge0.50B:H microstructures, with a
compressive stress of 10.66 MPa. These microstructures showed excellent characteristics for their integration in the production
of MEMSs.
Keywords: MEMS, Vernier beam, diamond-type structure, amorphous silicon, amorphous silicon germanium, PECVD
Predstavljen je postopek mikroobdelave pri izdelavi mikroelektromehanskega sistema (MEMS) z borom dopiranih silicijevih in
silicij-germanijevih amorfnih tankih plasti (a-SiB:H in a-Si0,50Ge0,50B:H), izdelanih s kemijskim postopkom nana{anja iz
parne faze v plazmi (PECVD) pri 110 kHz in 300 °C pri tlaku 8 · 10–4 mbar. MEMS so bili izdelani s tehnologijo mikroobdelave
povr{ine z mokrim in suhim jedkanjem. Izdelane so bile mikrostrukture iz a-Si0,50Ge0,50B:H in a-SiB:H, ki sestojijo iz ene
plasti debeline 1 μm. V proces MEMS so bile vklju~ene diamantne, obro~aste, V in Vernierjeve stebraste strukture. V plasti
a-SiB:H in a-Si0,50Ge0,50B:H smo na{li opti~ne vrzeli pri 1,55 eV in 0,97 eV. Spreminjanje koncentracije dopiranega bora v
vzorcu a-Si0,50Ge0,50B:H je pripeljalo do zmanj{anja upornosti od 6,84 M na 0,2 M. Tla~na napetost 13,44 MPa je bila
dose`ena v diamantni mikrostrukturi a-SiB:H. Izbolj{ana mehanska stabilnost je bila dobljena pri mikrostrukturi
a-Si0,50Ge0,50B:H s tla~no napetostjo 10,66 MPa. Te mikrostrukture ka`ejo odli~ne lastnosti za njihovo vklju~itev pri izdelavi
MEMS.
Klju~ne besede: MEMS, Vernierjev steber, diamantna struktura, amorfni silicij, amorfni silicij-germanij, PECVD
1 INTRODUCTION
Recent advances in micromachining technologies
related with the manufacturing of micro-electro-mecha-
nical systems (MEMSs) made it necessary to investigate
new materials to be used for the fabrication of modern
devices1–3. Some materials can be deposited at the
temperatures that are low enough (below 300 °C) to
remain unaffected during a regular process involving a
thermal load in the production of any complex micro-
electronics, for instance, when sensors and/or actuators
(accelerometers, anemometers, gyroscopes, thermal
devices, RF-transmitters, photodetectors, optical modu-
lators, micromirrors4–8) must be integrated using a com-
plementary metal-oxide semiconductor (CMOS) tech-
nology9,10. A variety of materials, such as amorphous Si
(a-Si), amorphous Ge (a-Ge) and the amorphous Si-Ge
alloy (a-SiGe), can be nowadays deposited onto different
substrates, using plasma-enhanced chemical vapor
deposition (PECVD)11–13. Currently, these materials are
attracting the attention because their electrical and
mechanical properties are suitable for manufacturing
microstructures.
The a-SiGe films were developed mainly for the
applications in optoelectronic devices. The optical band
gap of an amorphous material is determined with the
content ratio of Si/Ge in the film. In general, optical,
electrical and mechanical properties can be varied
according to the type and ratio of the introduced gases
during the deposition. In particular, the mechanical pro-
perties provide a new property of amorphous semicon-
ductors for fabricating microstructures.
On the other hand, most of the microstructures for
MEMSs are made with polysilicon (i.e., polycrystalline
silicon); however, this material is deposited and doped at
650 °C and 900 °C, respectively, and this affects the
Materiali in tehnologije / Materials and technology 49 (2015) 1, 3–8 3
UDK 678.84 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)3(2015)
thermal budget of the process. Thus, on the basis of our
knowledge of the PolyMEMSV process, we fabricated
microstructures using a-SiGe films. We used this mate-
rial because a-SiGe films, deposited at 300 °C, can show
properties similar to those presented by polysilicon
processed at much higher temperatures. Furthermore, the
a-SiGe properties such as Young’s modulus, bond con-
tent, thickness, resistivity and optical gap can be varied
with the deposition conditions (frequency, pressure,
temperature and gas-flow rate). It was found14 that Si-Ge
forms polymeric chains that do not alter at all the
procedure explained in this article.
1.1 Mechanical structures for monitoring mechanical
stress
One of the main problems in manufacturing surface
microstructures is the presence of residual stresses.
Residual stresses (i = E) are internal mechanical forces
that act on an isotropic material without the applications
of external forces or temperature gradients. Residual
stresses lead to mechanical deformations () in the ma-
terials, resulting in fractional changes in the dimensions
(linear, surface or volume)12,13.
In order to develop a microsystem technology, we
must include the structures for monitoring the efforts and
residual stress on the film so that we are able to observe
the presence of mechanical stresses due to the deforma-
tion. The presence of a mechanical stress causes an
unpredictable behavior in static or dynamic structures. A
bridge-type structure (clamped-clamped beam), under a
compressive stress, presents a deformation along its
structure when it is released (buckling). The critical
deformation cr can be estimated using Euler equation 1
for a beam with the critical length Lcr, as shown in
Figure 1:
 cr
cr
2= +
π 2 2 2
12
4 4
L
D hz( ) (1)
In the above equation h is the thickness of the beam,
Lcr is the critical length of the beam and Dz is the strain
amplitude.
1.2 Young’s modulus
Young’s modulus is a measure of the stiffness of an
elastic material. The residual stress (i) is directly pro-
portional to Young’s modulus; therefore, a higher
Young’s modulus corresponds to a higher stress in the
material. An ANSYS® finite-element simulator was
used to calculate the Young’s modulus for a-SiGe
films15,16 and this allowed a simulation of the dynamic
behavior of the microstructure beams. A modal analysis
was used to determine the first natural mode of the
resonance frequency of the beam with the following con-
ditions:
• The beam is considered as a harmonic oscillator.
• The beam is anchored at one end and free at the other
end.
• The resonance frequency is proportional to the
dimensions of the beam (L = 100 μm, h = 1 μm and
w = 10 μm).
• The properties of the materials are: thermal-expan-
sion coefficients, density, thermal conductivity17,18.
The resonance frequency is a function of the dimen-
sions and mechanical characteristics of the material:
F
E h
Lres
=
−
01604
1
2.
( )


(2)
Here h is the thickness of the thin film, L is the length
of the beam,  is the density of the material, E is Young’s
modulus,  is Poisson’s ratio and Fres is the resonance
frequency of the beam15.
1.3 Design and manufacturing of microstructures
Any PolyMEMSV chip includes passive structures
such as bridge, diamond, ring, V and Vernier-beam struc-
tures acting as the monitors of the residual mechanical
stress of a-SiB:H and a-Si0.50Ge0.50B:H films.
The microstructure manufacturing consists of depo-
sition and etching of the materials, while the micro-
structure release is done with the chemical-etching tech-
niques: wet and/or dry16–23. Surface micromachining is
used to remove the material used as a sacrificial film20–22.
In other cases, the bulk micromachining of silicon moves
the wafer back24. Dry etching is done with the reac-
tive-ion-etching system (RIE) or inductively coupled
plasma (ICP). In contrast, wet etching uses chemical
solutions, usually alkaline potassium hydroxide (KOH)
and sodium hydroxide (NaOH). In the cases of both wet
and dry etching, it is necessary to define the geometric
patterns on the wafer to manufacture the microstructures
with a photolithographic process.
The PolyMEMSV process is based on material depo-
sition and wet and dry attacks for surface microma-
chining. Wet etching with potassium hydroxide and
water (H2O:KOH) at a temperature of 32 °C is used. Dry
etching based on the reactive ion attack (RIE) is carried
out with MicroRIE-800 from Technics Inc.25; CHF3/O2
and SF6/O2 gasses are used24. Photolithography with a
positive photoresistance (ma-P1225) at 3000 r/min was
applied for 30 s.26 Our manufacturing processes are stan-
dardized to the level of a-SiGe of 1 μm on a phospho-
silicate-glass (PSG) layer of 3 μm as the sacrificial
material and an aluminum layer of 1 um.
M. GALINDO et al.: BORON-DOPED HYDROGENATED AMORPHOUS SEMICONDUCTOR MEMS
4 Materiali in tehnologije / Materials and technology 49 (2015) 1, 3–8
Figure 1: Bridge-type structure with a compressive stress
Slika 1: Oblika zgradbe mosti~ka s tla~no napetostjo
1.4 Experimental setup
The experimental setup consisted of a silicon oxide
(SiO2) film as the electrical insulator, a PSG material as
the sacrificial material (or a temporary mechanical
support), an amorphous-film deposition and aluminum
(Al) as the electrical conductor, and a crystalline silicon
wafer used as the substrate. The manufacturing process
is carried out in the following steps:
1. It begins with the growth of 200 nm SiO2 films at a
temperature of 1000 °C, for 30 min, in a H2O vapor
flow (Figure 2a).
2. A PSG sacrificial film 3 μm is deposited using an
atmospheric-pressure chemical vapor deposition
system (APCVD.) This film serves as a temporary
mechanical support for the structures suspended. The
conditions are: a temperature of 450 °C, nitrogen
(N2), silane (SiH4) and phosphine (PH3) flows, dura-
tion of 58 min (Figure 2b).
Lithography and wet etching are performed on the
PSG film in order to deposit a microstructure film,
applying a negative photoresistance (ma-1420) at 5000
r/min for 30 s, during an exposure time of 10 s.25 De-
veloped in 35 s and using a thermal annealing at 110 °C
for 15 min, wet etching is performed with potassium
hydroxide (KOH) to 10 % at a temperature of 32 °C for 3
min (Figure 2c).
The a-SiB:H and a-Si0.50Ge0.50B:H films 1 μm are
deposited with the PECVD system under the following
conditions: a temperature of 300 °C, a frequency of 110
kHz, a power of 500 W and a pressure of 8 · 10–4 mbar.
Silane (SiH4), germane (GeH4), hydrogen (H2) and dibo-
rane (B6H4) flows are used (Figure 2d).
The microstructures of the a-Si1B:H and
a-Si0.50Ge0.50B:H films are defined with lithography
and dry etching in RIE. Sulfur tetrafluoride (SF4) gas is
used to generate the RIE plasma, with an electrical
power of 200 W and the base pressure of 4 · 10–1 mbar
(Figure 2e).
Aluminum 1 μm used for electrical contacts is depo-
sited with a physical vapor deposition system (PVD).
Aluminum etching is performed using a chemical solu-
tion of H3PO4, CH3COOH and HNO3 for 7 min at 27 °C
(Figure 2f).
The last step includes a removal of the PSG sacrifi-
cial layer to release the suspended microstructures. The
technique used to release the microstructures consists of
eliminating the PSG with a hydrofluoric acid solution
(HF to 49 %) at room temperature for 4 min. The micro-
structure release is performed with 2-propanol at 62 °C,
water DI at 62 °C and acetone at 62 °C to remove the
residues in order to avoid a collapse (Figure 2g).
2 RESULTS AND DISCUSSION
The spectra of the optical absorption were measured
with a UV/Vis 300 Spectronic-Unicam spectrometer20.
The optical gap Eg of the films was extracted according
to the empirical formula derived by Tauc26–29. Figure 3
shows the band optical gap Eg versus the Ge content (Xg)
for the films studied here. The values of the optical gap
of 1.55 eV for the a-SiB:H film are also shown. As the
Ge content increases from 0 to 0.25 the optical gap
decreases rapidly from 1.55 eV to 1.06 eV for the
a-Si0.75Ge0.25B:H films. When the Ge content Xg > 0.5
(a-Si0.50Ge0.50B:H) the change in the optical gap
becomes slower, 0.97 eV and 0.78 eV for the a-GeB:H
films.
The electrical and mechanical properties, the stress
and roughness of a film are important in the fabrication
of microstructures. The data obtained from the a-SiB:H
thin films and the a-Si1-XGeXB:H alloy with different
molar concentrations of Ge show that the optical gap
M. GALINDO et al.: BORON-DOPED HYDROGENATED AMORPHOUS SEMICONDUCTOR MEMS
Materiali in tehnologije / Materials and technology 49 (2015) 1, 3–8 5
Figure 3: Band gap versus Ge content for a-Si1-XGeXB:H thin films
Slika 3: Energijska vrzel v tanki plasti a-Si1-XGeXB:H glede na
vsebnost Ge
Figure 2: Sequence of the steps for the production of the microstruc-
tures of a-Si1B:H and a-Si0.50Ge0.50B:H with a thickness of 1 μm
Slika 2: Zaporedje korakov pri izdelavi mikrostrukture a-Si1B:H in
a-Si0,50Ge0,50B:H debeline 1 μm
varies with the B content, but the preferential Ge incor-
poration in the solid phase is evident.
Figure 4 shows the current-voltage curves for the
sample of a-Si0.50Ge0.50B:H doped with different
amounts of boron; this graph shows that the slope of
each curve changes in the range of 0.12 μA to 1.7 μA.
These changes in the slope of the current-voltage curve
between different samples are due to the variation in the
electrical resistance of different samples. The electrical
resistance decreases with the amount of boron present in
the samples in the range of 6.84 M to 0.514 M. The
a-Si0.50Ge0.50B:H sample doped with 60 % of boron
has less electrical resistance. The result of an electrical
characterization was obtained with a Keithley Model
2400 SourceMeter16,30.
Figure 5 shows the resonance frequency obtained
through the finite-element method using ANSYS®16,
allowing us to calculate the Young’s modulus from
Equation 2 as a function of the germanium concentration
in the bridge-type microstructure.
The analysis showed that Young’s modulus can be
assumed as a linear combination of the amounts of sili-
con and germanium in thin films.
We designed a microstructure-manufacturing process
with the a-Si0.50Ge0.5B:H films deposited at the tem-
perature of 300 °C to use it for post-processing MEMS
with the CMOS technology. The SEM images were
recorded at an acceleration voltage of 15 kV and in a
high vacuum (HV) using a JEOL SEM model JSM-
5610LV (Hitachi, Japan). The samples were placed on a
specimen stub using a double-sided adhesive carbon
tape.
The microstructures manufactured at a low tempera-
ture of 300 °C, with a structural level of hydrogenated
amorphous silicon doped with boron (a-Si1B:H) exhi-
bited a higher roughness on the surface and a poor
definition at the edges of the structures; in addition to the
chemical attack on the surface of the wafer and the struc-
tures, Figure 6 shows SEM images of thermal micro-
actuator chevron (TVM) type microstructures.
Additionally, when using thin films of a-Si0.50
Ge0.50B:H a direct dependence on the roughness and
the definition of the edge is observed, while others do
not show lateral etching on the structures or on the wafer.
Figure 7 shows the profiles of thermal bimorph
micro-actuator (TBA) type microstructures manufactured
with the a-Si0.50Ge0.50B:H films including a structural
layer.
The microstructures of a-SiB:H and a-Si0.50Ge0.50
B:H show the presence of the residual stress (i) of
compression in suspended microstructures. Figure 8
M. GALINDO et al.: BORON-DOPED HYDROGENATED AMORPHOUS SEMICONDUCTOR MEMS
6 Materiali in tehnologije / Materials and technology 49 (2015) 1, 3–8
Figure 6: a) Geometric profile and b) close-to-TVM-type microstruc-
ture
Slika 6: a) Geometrijski profil in b) mikrostruktura blizu TVM
Figure 4: Current-voltage curves for the a-Si0.50Ge0.50B:H sample
doped with different amounts of boron
Slika 4: Odvisnost toka od napetosti pri vzorcu a-Si0,50Ge0,50B:H,
dopiranem z razli~nimi vsebnostmi bora
Figure 7: SEM images of the microstructure of a-Si0.50Ge0.50B:H:
a) TBA-type microstructure, b) close up of the TBA-type microstruc-
ture
Slika 7: SEM-posnetka mikrostrukture a-Si0,50Ge0,50B:H: a) TBA-
vrsta mikrostrukture, b) blizu TBA-vrsti mikrostrukture
Figure 5: Resonant frequency and Young’s modulus as functions of
the concentration
Slika 5: Resonan~na frekvenca in Youngov modul v odvisnosti od
koncentracije
shows images of diamond-type structures with different
lengths and a width of 10 μm across the central bar; the
structures present an evident deformation known as the
critical length of deformation (buckling), cr.
According to Equation 1, the deformation of a-SiB:H
diamond-type structures, with a length of 200 μm across
the central bar and a width of 10 μm, is 89.363 μm, and
the residual stress (i = E) is calculated to be 13.444
MPa. The a-Si0.50Ge0.50B:H diamond-type structure
does not show any lateral chemical attacks, but it indi-
cates a compressive residual stress of 10.661 MPa (Fig-
ure 8).
Figure 9 shows that the stress is higher when the
length of the central bar is smaller; this is because the
deformation (cr) is directly proportional to the length of
the bar and the stress – i is directly proportional to the
Young’s modulus. The Young’s modulus of a-Si:H is 160
GPa and for a-Si0.50Ge0.50B:H the Young’s modulus is
130 GPa. So, the a-Si0.50Ge0.50B:H microstructures
present less stress.
3 CONCLUSIONS
The work presents the results of manufacturing
microstructures with a-SiB:H and a-Si0.50Ge0.50B:H
films deposited with a PECVD system at a temperature
of 300 °C. The electrical and optical characterization of
a-Si1-XGeXB:H thin films was made. The value of the
optical gap of the a-SiB:H film is 1.55 eV and for the
a-Si0.50Ge0.50B:H film it is 0.97 eV. When germanium
content increases, the optical gap decreases rapidly from
1.55 eV to 0.78 eV for a-GeB:H films. The electrical
resistance of the a-Si0.50Ge0.50B:H film doped with
different concentrations of boron becomes reduced from
6.84 M to 0.56 M. The Young’s modulus of amor-
phous silicon-germanium alloys presents a linear
behavior of the silicon/germanium content.
The a-Si0.50Ge0.50B:H suspended microstructures
fabricated with a surface micromachining technique at a
temperature of 300 °C provide an excellent definition of
geometric patterns, having a high mechanical stability in
contrast with the microstructures of a-Si1B:H. The
advantage of working at low temperatures is a reduction
of the residual stress and the intrinsic gradient present at
high temperatures. However, diamond-type structures
exhibit a compressive mechanical stress of 13.44 MPa
for the a-SiB:H structural layer and 10.66 MPa for the
a-Si0.50Ge0.50B:H structural layer. We can, therefore,
conclude that the microstructures with the a-Si0.50
Ge0.50B:H structural layers deposited at 300 °C using
PECVD with the heat treatments < 350 °C can be
applied in post-CMOS processes and MEMSs as struc-
tural materials. In future we plan to integrate the manu-
facturing devices and microstructures within a die.
Acknowledgements
The work was supported by Consejo Naciuonal de
Ciencia y Tecnología through grant 160164 and by Red
Nacional de MEMS.
4 REFERENCES
1 H. Baltes, O. Brand, A. Hierlemann, D. Lange, C. Hagleitner,
CMOS MEMS – present and future, Proc. of the International
conference on Micro Electro Mechanical Systems, Las Vegas, 2002,
459–466
2 S. Beeby, G. Ensell, M. Kraft, N. White, MEMS mechanical sensors,
Artech House, Norwood, MA 2004
3 J. W. Judy, Microelectromechanical systems (MEMS): fabrication,
design and applications, Journal of Smart Materials and Structures,
10 (2001), 1115–1134
4 E. K. Petersen, Silicon as a mechanical material, Proceedings of the
IEEE, 70 (1982), 420–455
5 J. Stauth, Micromechanical electrostatic resonators in a standard
post-CMOS process, UC Berkeley-EECS, 2005, 1–4
M. GALINDO et al.: BORON-DOPED HYDROGENATED AMORPHOUS SEMICONDUCTOR MEMS
Materiali in tehnologije / Materials and technology 49 (2015) 1, 3–8 7
Figure 9: Compression residual stress of diamond-type microstruc-
tures versus the length
Slika 9: Zaostale tla~ne napetosti pri diamantni mikrostrukturi v od-
visnosti od dol`ine
Figure 8: Compression stress (–i) for diamond-type microstructures:
a) microstructure of a-SiB:H, b) microstructure of a-Si0.50Ge0.50B:H
Slika 8: Tla~na napetost (–i) za diamantno mikrostrukturo: a) mikro-
struktura a-SiB:H, b) mikrostruktura a-Si0,50Ge0,50B:H
6 F. Yuan, CMOS Current-Mode Circuits for Data Communications,
Springer, 2007
7 J. Laconte, D. Flandre, J. P. Raskin, Micromachined Thin-Film Sen-
sors for SOICMOS Co-Integration, Springer, 2006, 47–103
8 A. Witvrouw, F. Steenkiste, D. Maes, L. Haspeslagh, P. Gerwen, P.
Moor, S. Sedky, C. Hoof, A. Vries, A. Verbist, A. Caussemaeker, B.
Parmentier, K. Baert, Why CMOS integrated transducers? A review,
Microsystem Technologies, 6 (2000), 192–199
9 T. Hidekuni, S. Kazuaki, I. Makoto, Integrated inductors for RF
transmitters in CMOS/MEMS smart microsensor systems, Sensors, 7
(2007), 1387–1398
10 F. Mayer, A. Haberli, H. Jacobs, G. Ofner, H. Baltes, Single CHIP
CMOS anemometer, International Electron Devices Meeting,
Washington, 1997, 895–898
11 H. Guckel, D. Burns, C. Rutigliano, E. Lovell, B. Choi, Diagnostic
microstructures for the measurement of intrinsic strain in films, J.
Micromech, Microeng, 2 (1992), 86–95
12 H. Guckel, T. Randazzo, D. Burns, A simple technique for the deter-
mination of mechanical strain in thin films with applications to
polysilicon, J. Appl. Phys., 57 (1985) 5, 1671–1675
13 Research Institute. http://www.inaoep.mx/ (accessed 20 April 2010)
14 G. Craciun, H. Yang, W. Van Zeijl, L. Pakula, A. Blauw, E. Van der
Drift, J. French, Single step cryogenic SF6/O2 plasma etching
process for the development of inertial devices, Proceedings on
semiconductors sensors, Fabrication and characterization, 2002,
612–615
15 W. Brostow, V. M. Castaño, G. Martinez-Barrera, J. M. Saiter,
Pressure-volume-temperature properties of GexSe1-x inorganic
polymeric glasses, Physica B, 334 (2003), 436–443
16 F. López, A. Herrera, J. Estrada, C. Zúñiga, B. Cervantes, E. Soto, B.
S. Soto-Cruz, Alternative Post-Processing on a CMOS Chip to
Fabricate a Planar Microelectrode Array, Sensors, 11 (2011),
10940–10957
17 Software products for engineering simulation, http://www.ansys.
com/ (accessed in December 2009)
18 W. Hernández, I. Zaldivar, C. Zúñiga, A. Torres, C. Reyes, A. Itzmo-
yol, Optical and compositional properties of amorphous silicon-ger-
manium films by plasma processing for integrated photonics, Optical
Materials Express, 2 (2012), 358–370
19 Laboratory Equipments, Available online: http://www.echromtech.
com (accessed in September 2011)
20 B. Nelson, Y. Xu, D. Williamson, D. Han, R. Braunstein, M. Boshta,
B. Alavi, Narrow gap a-SiGe:H grown by hot-wire chemical vapor
deposition, Thin Solid Films, 430 (2003), 104–109
21 Electronic Instruments, Available online: http://www.keithley.com
(accessed in November 2011)
22 M. Christian, H. Qiuting, A Low-Noise CMOS instrumentation
amplifier for thermoelectric infrared detectors, IEEE Journal of
Solid-State Circuits, 32 (1997), 968–976
23 I. Toledo, R. Adler, Y. Knafo, O. Kalis, J. Kaplun, BCB etching
process using high density plasma, CS MANTECH Conference,
Chicago, Illinois, USA, 2008, 14–17
24 M. Tsai, C. Sun, Y. Liu, C. Wang, W. Fang, Design and application
of a metal wet-etching post-process for the improvement of
CMOS-MEMS capacitive sensors, Journal of Micromechanics and
Microengineering, 19 (2009), 1–7
25 D. Fernández, J. Ricart, J. Madrenas, Experiments on the release of
CMOS micromachined metal layers, Journal of Sensors, 2010
(2010), article ID 937301, 1–7
26 P. Merz, H. Quenzer, H. Bernt, B. Wagner, M. Zoberbier, A novel
micromachining technology for structuring borosilicate glass
substrates, The 12th International Conference on Solid State Sensors,
Actuators and Microsystems, Boston, MA, USA, 2003, 258–261
27 H. Berney, M. Hill, D. Cotter, E. Hynes, M. O’Neill, A. Lane,
Determination of the effect of processing steps on the CMOS
compatibility of a surface micromachined pressure sensor, Journal of
Micromechanics and Microengineering, 11 (2001), 402–408
28 A. Jihwan, V. Brian, Micromachined nerve electrode array integrated
with shape memory thin film position clamp, Proceedings of E240,
2007, 1–4
29 T. Gregory, I. Nadimi, E. Kurt, Bulk micromachining of silicon,
Proceedings of the IEEE, 86 (1998), 1536–1550
30 Microresist Technology, Available online: http://www.microresist.de/
home_en.htm (accessed in September 2011)
M. GALINDO et al.: BORON-DOPED HYDROGENATED AMORPHOUS SEMICONDUCTOR MEMS
8 Materiali in tehnologije / Materials and technology 49 (2015) 1, 3–8
T. THIYAGARAJAN et al.: WEAR BEHAVIOUR OF B4C REINFORCED HYBRID ALUMINUM-MATRIX COMPOSITES
WEAR BEHAVIOUR OF B4C REINFORCED HYBRID
ALUMINUM-MATRIX COMPOSITES
VEDENJE HIBRIDNEGA KOMPOZITA NA OSNOVI ALUMINIJA,
OJA^ANEGA Z B4C, PRI OBRABI
Thirumalai Thiyagarajan1, Ramanathan Subramanian2, Somasundaram
Dharmalingam3, Nachiappan Radika4, Annadurai Gowrisankar2
1Department of Mechanical Engineering, Erode Sengunthar Engineering College, Perundurai, Erode, India
2Department of Metallurgical Engineering, PSG College of Technology, Coimbatore, India
3Department of Mechanical Engineering, RVS Faculty of Engineering, Coimbatore, India
4School of Mechanical Engineering, Amrita Vishwa Vidyapeetham University, Coimbatore, India
thirumalaimet@gmail.com
Prejem rokopisa – received: 2013-08-12; sprejem za objavo – accepted for publication: 2014-03-24
Harnessing the benefits of aluminum hybrid composites requires strengthening of materials with better reinforcing materials.
Hence, this investigation was carried out to identify the optimum level of B4C reinforcement to enhance the wear resistance of
aluminum alloys. Four amounts of B4C reinforcement (3, 6, 9 and 12) % and a constant amount of 3 % graphite were mixed
with an aluminum alloy to produce aluminum hybrid composites. The hybrid composites thus produced were tested for wear
resistance at various loads and sliding speeds. The change in the wear resistance of the newly produced hybrid material was
studied with SEM micrographs to confirm the results. Al-B4C-Gr composites produced with the stir-casting method including
various amounts of B4C reinforcement show that the B4C reinforcement of up to 9 % was found beneficial as it increased the
wear resistance of the aluminum alloy. The addition of a constant amount of graphite provides for the lubrication effect, while
B4C increases the strength of the parent material (Al). SEM micrographs confirmed that aluminum hybrid composites showed
less debris and finer grooves with delamination at lower loads and sliding speeds. Hence, 9 % B4C reinforcement with 3 %
graphite is a better way of increasing the wear resistance of aluminum alloys.
Keywords: aluminum-matrix composites, boron carbide, graphite, wear, stir casting, sliding distance
Izkori{~anje prednosti hibridnih kompozitov na osnovi aluminija zahteva oja~anje materiala z bolj{imi materiali za utrjevanje.
Zato je bila izvr{ena ta preiskava, da bi se ugotovil optimalni dodatek B4C za pove~anje obrabne odpornosti aluminijevih zlitin.
[tiri vsebnosti B4C (3, 6, 9 in 12) % kot dodatka za utrjevanje in 3 % grafita so bile dodane aluminijevi zlitini, da bi dobili
aluminijeve hibridne kompozite. Tako izdelani hibridni kompoziti so bili preizku{eni v odpornosti proti obrabi pri razli~nih
obte`bah in hitrostih drsenja. Potrditev rezultatov in spremembe v odpornosti proti obrabi hibridnih materialov so bile ugotav-
ljane s SEM-posnetkov. Kompoziti Al-B4C-Gr, izdelani z metodo vme{avanja v talino, z razli~no vsebnostjo sredstva za utrditev
B4C so pokazali, da je dodatek do 9 % B4C ugoden in pove~a obrabno odpornost aluminijeve zlitine. Dodatek enake koli~ine
grafita zagotavlja u~inek mazanja, medtem ko B4C daje vi{jo trdnost osnovnemu materialu (Al). SEM-posnetki hibridnih
kompozitov aluminija so pokazali manj zlomnin in drobnej{e opilke z delaminacijo pri manj{ih obte`bah in hitrostih drsenja.
Torej je dodatek 9 % B4C za oja~anje in 3 % grafita bolj{a tehnologija za pove~anje obrabne odpornosti aluminijevih zlitin.
Klju~ne besede: kompoziti na osnovi aluminija, borov karbid, grafit, obraba, vme{avanje v talino, razdalja drsenja
1 INTRODUCTION
Aluminum-matrix composites (AMCs) have recently
gained momentum due to their structural applications in
aircraft, automotive, construction, packaging, electronics
and military industries. Next to iron and steel, aluminium
alloys are the most widely used metallic materials. They
are used particularly in the manufacturing of automotive
components such as cylinder liners, pistons, drive shafts,
brake rotors, cylinder heads, cylinder blocks, intake
manifolds, rear axles and differential housings1–3. Their
high strength-to-weight ratio, high thermal conductivity
and specific stiffness enhanced the interest of the re-
searchers to further improve the material characteristics
so that they are appropriate for all the applications. Alu-
minium represents around 8 % of the vehicle curb weight
in today’s cars, trucks and minivans. However, an
inadequate wear resistance and low seizure loads prevent
a direct use of aluminium alloys in automotive parts due
to the intensive friction, high thermal and mechanical
loading4.
Boron carbide (B4C) is found to be a promising cera-
mic material due to its high strength, low density (2.52
g/cm3), extremely high hardness and better chemical
stability. Extreme hardness of this material makes it a
better alternative to silicon carbide (SiC) and aluminum
oxide (Al2O3) for reinforcement with aluminum alloys. A
hybrid aluminum-matrix composite combines the high
strength and hardness of reinforcing materials with the
ductility and toughness of light metals5,6. A study con-
ducted to test the wear behaviour of Al-B4C and Al-SiC
composites fabricated with the stir-casting method
revealed that the wear rate and friction coefficient of
Al-B4C were lower than those of Al-SiC7. The alumi-
num-matrix composite reinforced with B4C particles and
SiC particles through the pressureless-infiltration method
indicated that the strength of the Al-B4C composite was
greater than that of the Al-SiC composite.8 Hence, a
Materiali in tehnologije / Materials and technology 49 (2015) 1, 9–13 9
UDK 669.018.9:539.538:669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)9(2015)
decrease in the reinforcement particle size to the nano-
meter range can improve the mechanical and tribological
properties of AMCs. The tribological behaviour of stir-
cast Al-Si/SiC composites (15 % and 20 % volume frac-
tion) against the automobile-brake-pad material using a
pin-on-disc tribotester showed an inverse proportionality
between the sliding speed and the wear rate9. Dry-sliding
wear performance of a hypereutectic A390 Al-Si alloy
reinforced with graphite particulates (4 % and 8 %)
showed that both the wear rate and COF of the compo-
sites decreased considerably with the addition of
graphite10.
The fabrication and characterization of bulk Al-B4C
nanocomposites containing different mass fractions of
B4C w = (5, 10 and 15) % showed an increase in the
wear resistance with the increasing B4C amount11. The
fabrication of mass fractions 40 % SiC/% Gr/Al compo-
sites with the additions of various amounts of graphite
using the squeeze-casting technology showed that the
addition of graphite decreased the friction coefficient of
the composites and increased the wear resistance by 170
to 340 times12. A study on the influence of the sliding
speed on dry-sliding wear behaviour and the extent of
subsurface deformation in an Al 2219/15 % SiCp com-
posite and an Al 2219/15 % SiCp-graphite hybrid com-
posite due to liquid metallurgy reported that the wear
rates of the composites remained almost unchanged with
the increasing sliding speed up to 4.6 m/s, after which an
increasing trend in the wear rate was observed13.
Extensive researches on individual reinforcements of
boron carbide and graphite improving the wear resi-
stance and the strength of hybrid aluminum composites
were carried out. However, very limited research was
conducted in order to explore the combined effect of
boron-carbide and graphite-particulate reinforcements on
hybrid aluminum composites. Hence, this study was
carried out i) to develop new hybrid aluminum compo-
sites with boron carbide and graphite and ii) to test the
improvement in the wear resistance of the material under
varying sliding speed and load.
2 MATERIALS AND METHODS
Hybrid aluminum composites with boron carbide
(B4C) in various amounts and constant graphite rein-
forcement were cast using the stir-casting method. The
hybrid Al composites were reinforced with (3, 6, 9 and
12) % of B4C and a constant amount of 3 % graphite.
The boron-carbide and graphite particles with the ave-
rage particle size of 25 μm to 75 μm were preheated up
to 300 °C. Degassing tablets were added during the
melting to remove gaseous molecules. Magnesium (w =
2 %) was added to the molten metal to improve the
wettability between the reinforcements and aluminium.
The mixture of the molten alloy was mechanically
stirred using a steel stirrer for 10 min to obtain a homo-
genous mixture. A speed of 400 r/min was maintained
and a pouring temperature of 710 °C was maintained.
After removing them from the metal dye, the specimens
were collected and subjected to the wear tests performed
with a pin-on-disc machine. The test specimens were in
the form of pins having a diameter of 10 mm and a
height of 40 mm. The ends of the specimens were
polished with abrasive paper of grade 600 followed by
grade 1000. The disc was made of EN-31 steel having a
hardness of 63 HRc under dry conditions. The wear loss
of the pin in microns was recorded during the wear test.
The pin surface wore out during the rubbing with the
counter disc, continuously moving down to have a con-
tact with the surface of the disc. The tests were con-
ducted as per ASTM G99-95a test standards with a
constant sliding distance of 2500 m and a normal load of
10 N to 30 N. The sliding speed of 1 m/s to 3 m/s was
tested under dry conditions and the wear resistance was
recorded.
3 RESULTS AND DISCUSSIONS
Dry-sliding wear tests were conducted for the hybrid
composites and an unreinforced aluminium alloy (the
parent material) using a pin-on-disc apparatus at
different normal loads (10 N to 30 N) and sliding speeds
(1 m/s and 2 m/s) with a constant sliding distance of
2500 m. The wear rate was calculated by weighing the
specimens before and after the test. The variation in the
wear rate with varied levels of the normal load and
reinforcement percentage at different sliding speeds is
shown in the graphs (Figures 1 to 4). It can be observed
that the wear loss increases with the increasing load and
the highest wear is noted at the load of 30 N at both
sliding speeds. Comparing various materials tested for
the wear loss, it was evident that the maximum wear loss
occurred on the parent material (Al) and the minimum
wear loss occurred on 3 % Gr + 9 % B4C + Al compo-
T. THIYAGARAJAN et al.: WEAR BEHAVIOUR OF B4C REINFORCED HYBRID ALUMINUM-MATRIX COMPOSITES
10 Materiali in tehnologije / Materials and technology 49 (2015) 1, 9–13
Figure 1: Wear loss of hybrid aluminum composites with load at 1 m/s
sliding speed
Slika 1: Izguba zaradi obrabe hibridnega kompozita aluminija z ob-
te`bo pri hitrosti drsenja 1 m/s
sites (C3) which might have happened due to its better
wear resistance.
The wear resistance increases with the increasing
B4C amount up to 9 % in the composites, but a mixture
containing more than 9 % B4C decreases the wear resi-
stance. The beneficial effect noticed up to 9 % B4C rein-
forcement could be due to its higher hardness, causing a
better wear resistance of the composites11. Like the
parent material, the other compositions, 3 % Gr + 3 %
B4C + Al (C1), 3 % Gr + 6 % B4C + Al (C2) and 3 % Gr
+ 12 % B4C + Al (C4), also showed higher wear losses
which can be ascribed to larger quantities of the parti-
cles. Many investigators reported that the low density
(2.52 g/cm3), high hardness (HV = 30 GPa) and thermal
stability of boron carbide might be possible reasons for
the increased wear resistance14–17. A decreased wear
resistance at a higher level of B4C reinforcement (w =
12 %) could be due to a reduction in the interfacial
strength. This reduction in the interfacial strength is
attributed to the formation of a boron-rich Al-B phase at
the grain boundaries18,19. The sliding speed is an import-
ant factor that needs to be tested to understand the wear
loss of any new composite. Hence, two sliding speeds
were tested under various loads and B4C reinforcement
percentages and the results revealed that the wear loss
decreased with the increasing sliding speed irrespective
of the reinforcement percentage and the load. The wear
loss of the parent material was the highest at the higher
load (30 N) and it decreased with the sliding speed
which must have been due to the reduced particle reten-
tion time for abrasion11,20. The formation of the surface
T. THIYAGARAJAN et al.: WEAR BEHAVIOUR OF B4C REINFORCED HYBRID ALUMINUM-MATRIX COMPOSITES
Materiali in tehnologije / Materials and technology 49 (2015) 1, 9–13 11
Figure 5: SEM micrographs of the worn surfaces of hybrid Al-B4C
composites and aluminum (sliding speed 1 m/s, sliding distance 2500
m)
Slika 5: SEM-posnetki obrabljene povr{ine hibridnega kompozita
Al-B4C in aluminija (hitrost drsenja 1 m/s, razdalja drsenja 2500 m)
Figure 4: Wear loss of hybrid aluminum composites with reinforce-
ment fraction at 2 m/s sliding speed
Slika 4: Izguba zaradi obrabe hibridnega kompozita aluminija z dele-
`em oja~itvene faze pri hitrosti drsenja 2 m/s
Figure 3: Wear loss of hybrid aluminum composites with load at 2
m/s sliding speed
Slika 3: Izguba zaradi obrabe hibridnega kompozita aluminija z
dele`em oja~itvene faze pri hitrosti drsenja 2 m/s
Figure 2: Wear loss of hybrid aluminum composites with reinforce-
ment fraction at 1 m/s sliding speed
Slika 2: Izguba zaradi obrabe hibridnega kompozita aluminija z dele-
`em oja~itvene faze pri hitrosti drsenja 1 m/s
coatings of oxidized debris at the higher sliding speed
might have reduced both the wear rate and the coefficient
of friction.
Wear-track patterns of hybrid composites and the
parent material at three different loads and two sliding
speeds were studied by taking SEM micrographs. Fig-
ures 5 and 6 show the typical wear-track patterns deve-
loped on the surfaces of the composites and the parent
material. On the parent material, the increasing load
increases the formation of coarse grooves, debris parti-
cles and surface delamination. At the lower load (10 N),
more debris particles and abrasive wear of the surfaces
were noticed. With the increased load (20 N) wider
grooves and more debris were formed. Delamination,
wider grooves with more debris and mixed adhesive
wear were noticed when the parent material was
subjected to the 30 N load21,22.
Upon sliding, the higher speed caused abrasive wear
which pulled some of the material from the surface and
formed loose abrasive debris. As the sliding process pro-
ceeded with the increasing speed, the material attached
to alumina was delaminated and defoliated to form more
small debris and deep widened grooves. As the load
increased, greater amounts of the material from the
surfaces were defoliated, forming deep grooves with
wider widths due to ploughing in the aluminum alloy23.
In the case of major wear mechanisms, the abrasion wear
was predominant at the lower load (10 N) while at the
higher load (30 N) the adhesive wear was higher. In
contrast, the reinforcement with B4C and graphite
showed a beneficial effect on the wear resistance of the
composites.
The reinforced hybrid composites exhibited fine
grooves and less debris at the 10 N load. Like in the case
of the parent material, the increasing load increased the
delamination and the adhesive wear of the composites to
a lesser extent. The B4C reinforcement decreased the
groove width and debris formation due to an improve-
ment in the load-carrying capacity of the composites,
enhancing the abrasion resistance of the composites and,
thus, decreasing the wear loss24,25. At the lower load and
sliding speed (L = 10 N and S = 1 m/s) the worn pin
surface predominantly revealed fine and shallow grooves
in the sliding direction. Such features are characteristic
of the abrasive wear, where hard asperities of the counter
face plough into the hybrid composite pin, causing wear
by removing small fragments of the material.
4 CONCLUSIONS
Al-B4C-Gr composites were produced with the stir-
casting method using various levels of B4C reinforce-
ment and a constant amount of 3 % graphite addition. An
evaluation of the morphological properties and the wear
showed that the B4C reinforcement of up to 9 % was
beneficial in increasing the wear resistance of aluminum
hybrid composites. In the presence of graphite, B4C
made the parent material (Al) even stronger, resisting
both the abrasive and the adhesive wear. Subjecting the
composites and the unreinforced aluminum parent mate-
rial to various levels of the load showed that the in-
creased load increased the wear of the composites. At the
lower load and sliding speed, less debris and fine
grooves with delamination were formed which was con-
firmed with SEM micrographs. The improved wear
resistance of the composites is due to the hardness of
B4C and the lubricating effect of graphite. Hence, 9 %
B4C reinforcement with 3 % graphite can be a better way
of increasing the wear resistance of aluminum alloys.
5 REFERENCES
1 W. S. Miller, L. Zhuang, J. Bottema, A. J. Wittebrood, P. D. Smet, A.
Haszler, A. Vieregge, Recent development in aluminium alloys for
the automotive industry, Materials Science & Engineering A, 280
(2000), 37–49
2 I. Sinclair, P. J. Gregson, Structural performance of discontinuous
metal matrix composites, Materials Science and Technology, 13
(1997), 709–726
3 I. A. Ibrahim, F. A. Mohamed, E. J. Lavernia, Particulate Reinforced
Metal Matrix Composites – A Review, Journal of Materials Science,
24 (1991), 1137–1156
4 V. Kevorkijan, Engineering wear-resistant surfaces in automotive
aluminium, JOM, 55 (2003), 32–34
5 B. R. Henriksen, T. E. Johnsen, Influence of microstructure of fiber/
matrix interface on mechanical properties of Al/SiC composites,
Materials Science and Technology, 6 (1990), 857–863
T. THIYAGARAJAN et al.: WEAR BEHAVIOUR OF B4C REINFORCED HYBRID ALUMINUM-MATRIX COMPOSITES
12 Materiali in tehnologije / Materials and technology 49 (2015) 1, 9–13
Figure 6: SEM micrographs of the worn surfaces of hybrid Al-B4C
composites and aluminum (sliding speed 2 m/s, sliding distance 2500
m)
Slika 6: SEM-posnetki obrabljene povr{ine hibridnega kompozita
Al-B4C in aluminija (hitrost drsenja 2 m/s, razdalja drsenja 2500 m)
6 Y. H. Kim, C. S. Lee, K. S. Han, Fabrication and Mechanical Pro-
perties of Aluminum Matrix Composite Materials, Journal of Com-
posite Materials, 26 (1992), 1062–1086
7 K. M. Shorowordi, A. S. M. A. Haseeb, J. P. Celis, Tribo-surface
characteristics of Al–B4C and Al–SiC composites worn under diffe-
rent contact pressures, Wear, 261 (2006), 634–641
8 K. B. Lee, H. S. Sim, S. Y. Cho, H. Kwon, Tensile properties of 5052
Al matrix composites reinforced with B4C, Metall. Mater. Trans. A,
32 (2001), 2142–2147
9 R. K. Uyyuru, M. K. Surappa, S. Brusethaug, Tribological behavior
of Al-Si-SiCp composites/automobile brake pad system under dry
sliding conditions, Tribology International, 40 (2007), 365–373
10 T. S. Mahmoud, Tribological behaviour of A390/GrP metal-matrix
composites fabricated using a combination of rheocasting and
squeeze casting techniques, Journal of Mechanical Engineering
Science, 222 (2008), 257–265
11 E. Mohammad Sharifi, F. Karimzadeh, M. H. Enayati, Fabrication
and evaluation of mechanical and tribological properties of boron
carbide reinforced aluminum matrix nanocomposites, Materials and
Design, 32 (2011), 3263–3271
12 L. Jinfeng, J. Longtao, W. Gaohui, T. Shoufu, C. Guoqin, Effect of
Graphite Particle Reinforcement on Dry Sliding Wear of SiC/Gr/Al
Composites, Rare Metal Materials and Engineering, 38 (2009),
1894–1898
13 S. Basavarajappa, G. Chandramohan, A. Mahadevan, M. Thanga-
velu, R. Subramanian, P. Gopalakrishnan, Influence of sliding speed
on the dry sliding wear behaviour and the subsurface deformation on
hybrid metal matrix composite, Wear, 262 (2007), 1007–1012
14 S. Basavarajappa, G. Chandramohan, K. Mukund, M. Ashwin, M.
Prabu, Dry Sliding Wear Behavior of Al 2219/SiCp-Gr Hybrid Metal
Matrix Composites, Journal of Materials Engineering and Perfo-
rmance, 15 (2006) 6, 668–674
15 M. Gui, S. B. Kang, J. M. Lee, Dry sliding wear behavior of spray
deposited AlCuMn alloy and AlCuMn/SiCp composite, Journal of
Materials Science, 35 (2000), 4749–4762
16 R. M. Mohanty, K. Balasubramanian, S. K. Seshadri, Boron car-
bide-reinforced alumnium 1100 matrix composites: fabrication and
properties, Materials Science and Engineering A, 498 (2008), 42–52
17 J. Oñoro, M. D. Salvador, L. E. G. Cambronero, High-temperature
mechanical properties of aluminium alloys reinforced with boron
carbide particles, Materials Science and Engineering A, 499 (2009),
421–426
18 I. Kerti, F. Toptan, Micro structural variations in cast B4C-reinforced
aluminium matrix composites (AMCs), Materials Letters, 62 (2008),
1215–1218
19 U. B. Gopala Krishna, K. V. Sreenivas Rao, B. Vasudeva, Effect of
boron carbide reinforcement on aluminium matrix composites, Inter-
national Journal of Metallurgical & Materials Science and Engineer-
ing, 3 (2013), 41–48
20 Y. Sahin, S. Murphy, The effect of sliding speed and microstructure
on the dry wear properties of metal-matrix composites, Wear, 214
(1998), 98–106
21 D. Poirier, R. A. L. Drew, M. L. Trudeau, R. Gauvin, Fabrication and
properties of mechanically milled alumina/aluminum nanocompo-
sites, Materials Science and Engineering A, 527 (2010), 7605–7614
22 N. Hosseini, F. Karimzadeh, M. H. Abbasi, M. H. Enayati, Tribo-
logical properties of Al6061-Al2O3 nanocomposite prepared by
milling and hot pressing, Mater. Des., 31 (2010), 4777–4785
23 D. Cree, M. Pugh, Dry wear and friction properties of an A356/SiC
foam interpenetrating phase composite, Wear, 272 (2011), 88–96
24 M. Ramachandra, K. Radhakrishna, Effect of reinforcement of flyash
on sliding wear, slurry erosive wear and corrosive behavior of
aluminium matrix composite, Wear, 262 (2007), 1450–1462
25 M. Jafari, M. H. Enayati, M. H. Abbasi, F. Karimzadeh, Compressive
and wear behaviors of bulk nanostructured Al2024 alloy, Mater.
Des., 31 (2010), 663–669
T. THIYAGARAJAN et al.: WEAR BEHAVIOUR OF B4C REINFORCED HYBRID ALUMINUM-MATRIX COMPOSITES
Materiali in tehnologije / Materials and technology 49 (2015) 1, 9–13 13

J. MALCHARCZIKOVÁ et al.: INFLUENCE OF THE HIP PROCESS ON THE PROPERTIES ...
INFLUENCE OF THE HIP PROCESS ON THE PROPERTIES OF
AS-CAST Ni-BASED ALLOYS
VPLIV VRO^EGA IZOSTATSKEGA STISKANJA NA LASTNOSTI
Ni-ZLITIN Z LITO STRUKTURO
Jitka Malcharcziková1, Martin Pohludka1, Vít Michenka2, Tomá{ ^egan1,
Jan Juøica1, Miroslav Kursa1
1V[B-Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of Non-Ferrous Metals, Refining
Processes and Materials Recycling, Regional Materials Science and Technology Centre, 17. listopadu 15/2172, 708 33 Ostrava - Poruba, Czech
Republic
2VÚH@ a.s., Dobrá 240, 739 51 Dobrá, Czech Republic
jitka.malcharczikova@vsb.cz
Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2014-01-30
The main goal of this work was to evaluate the application of the samples prepared by centrifugal casting as the test samples for
a tensile test. Selected types of modified superalloys were prepared as experimental samples. The samples were molten by
vacuum-induction melting and then cast centrifugally into a shaped graphite mould. The final castings had the shape
corresponding approximately to the test specimens. Some of the samples were subjected to hot isostatic pressing (HIP). After
HIP the castings contained substantially fewer casting defects. Selected mechanical properties were compared for the samples in
the as-cast state and after HIP. For the majority of the investigated alloys the HIP process led to an increase in the strength and
ductility.
Keywords: Ni-based alloys, centrifugal casting, hot isostatic pressing, mechanical properties
Glavni cilj tega dela je bila ocena uporabnosti vzorcev za natezni preizkus, izdelanih s centrifugalnim litjem. Preizkusni vzorci
so bili izdelani iz izbrane vrste modificiranih superzlitin. Vzorci so bili staljeni v vakuumski indukcijski pe~i in centrifugalno
uliti v izoblikovane grafitne forme. Ulitki so imeli obliko, ki ustreza vzorcem za natezne preizkuse. Del vzorcev je bil vro~e
izostatsko stisnjen (HIP). Ulitki so imeli po obdelavi HIP ob~utno manj livarskih napak. Izbrane mehanske lastnosti so bile
primerjane pri vzorcih v litem stanju in po obdelavi HIP. Pri ve~ini preiskovanih zlitin je obdelava HIP povzro~ila pove~anje
trdnosti in duktilnosti.
Klju~ne besede: Ni-zlitine, centrifugalno litje, vro~e izostatsko stiskanje, mehanske lastnosti
1 INTRODUCTION
Ni-based alloys, including the alloys based on inter-
metallic compounds (IMC) and superalloys are still in
the forefront of interest. These alloys that can be used
even at high temperatures are continuously investigated
thanks to their excellent mechanical and corrosion
properties1,2. The technology of hot isostatic pressing
(HIP) belongs, within the powder metallurgy, to the
methods of hot sintering, during which the powder
achieves the required properties as a result of the effect
of predominantly physical processes. The HIP method is
successfully used also for additional compacting of
castings, ensuring a homogenisation of the structure and
a reduction of the pores and shrinkage cavities. This
method is used also for the alloys based on nickel and
Ni3Al. Numerous data were published on sintering
powder-pressed pieces made of the materials based on
Ni3–6. Kim7 describes a procedure for additional com-
pacting of the castings made of alloys IC396. Many
problems occur when processing the samples with the
usual production methods. The methodology for process-
ing to the specified dimensions is highly demanding and
expensive. A large amount of experimental material is
damaged8 and the testing of the mechanical properties of
this type of material is, therefore, very difficult.
The most frequent use of hot isostatic pressing is a
compaction of powder materials4,9–12. There are several
papers dealing with the compaction of nickel-alloy
castings with the HIP method13–15. Nickel-alloy castings
consist of large and irregular grains of the ’ phase. This
phase continues to coarse during the HIP process at a
high temperature and pressure4,9,12,15. In some cases,
carbides and the 	 phase precipitate at the grain
boundaries after a HIP process11. All of the above has a
negative influence on the mechanical properties of nickel
alloys. Therefore, after a HIP process, nickel alloys
should be heat treated by solution annealing and
aging11,12.
2 EXPERIMENT
Selected types of modified superalloys IC50, IC396,
IC221M and IC438 were prepared as the experimental
samples. They were molten by vacuum-induction melt-
ing and then cast centrifugally into a shaped graphite
mould (Figure 1). The final castings had the shape
corresponding approximately to the test specimens
Materiali in tehnologije / Materials and technology 49 (2015) 1, 15–18 15
UDK 669.245:621.74.042:621.7.016.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)15(2015)
(Figure 2). The compositions of the alloys were derived
from the already verified and industrially used types of
alloys for the given applications, and their compositions
are presented in Table 1 (A: the state after HIP, B: the
as-cast state).
Table 1: Chemical composition of the used samples in mass fractions,
w/%
Tabela 1: Kemijska sestava uporabljenih vzorcev v masnih dele`ih,
w/%
Alloy Sample
Amounts of elements
Al Cr Mo Zr B Ni
IC50 3A, 3B 11.30 - - 0.60 0.01 88.08
IC396 4A, 4B 7.98 7.72 3.02 0.85 0.01 80.42
IC221M 5A, 5B 8.00 7.70 1.43 1.70 0.01 81.10
IC438 6A, 6B 8.10 5.23 7.02 0.13 0.01 79.52
A flaw-detection method with X-ray radiation was
used for determining the quality of the castings. The
pattern of distribution and the type of casting defects
were similar in all the centrifugally cast samples. Figure
3 schematically shows their location and size. Defects
occurred on the tapered part of the test piece (the tensile
bar body), or in the direction towards the head, on the
part of the casting without a riser.
2.1 Hot isostatic pressing (HIP)
It was established after a flaw-detection analysis, that
the castings contained numerous casting defects. That is
why some of the samples were subjected to hot isostatic
pressing (HIP). The conditions for HIP were derived
from the literature sources3–7. Hot isostatic pressing was
performed on a press of the EPSI company. In the first
stage of the additional compacting we chose a pressure
of 100 MPa, the time of 4 h and a temperature of 1100
°C. Using flaw detection, we established that the castings
subjected to HIP under the pressure of 100 MPa still
contained a considerable amount of casting defects. For
this reason we increased the pressure to 150 MPa for the
HIP process, while the time of 4 h and the temperature of
1100 °C remained unchanged. However, even after
applying a higher pressure the castings still contained
casting defects, but to a much smaller extent than in the
state just after centrifugal casting.
Due to sufficient compaction, hot isostatic pressing
can create a homogenous structure with a reduced num-
ber of pores and casting defects (shrinkage cavities and
contractions) even in castings. This was confirmed in our
case as well. After HIP, the castings contained substan-
tially fewer casting defects.
2.2 Testing of the mechanical properties
Some of the samples in the as-cast state and the
samples after HIP were ground to precise dimensions of
the test piece and the tensile testing was performed.
Short tensile bars with a length of 55 mm and a diameter
of the central part of the bar of 5 mm were prepared by
lathe turning. The strain rate was approximately 9.1 ·
10–4 s–1 for all the samples. Table 2 gives the measured
values of the tensile stress at yield, the ultimate engi-
neering stress and the proportional elongation after
braking. The table also contains the average values of the
porosity and micro-hardness.
Table 2: Obtained mechanical characteristics
Tabela 2: Dobljene mehanske lastnosti
Sample
No.
Rp0.2/
MPa
Rm/
MPa
A/
%
Porosity
% HV 0.05
3A 239 949 32 0.0309 302
3B 313 824 19 0.0393 249
4A 539 668 6 0.0123 353
4B 720 754 6 0.0743 336
5A 550 995 28 0.0486 354
5B 653 776 7 0.0496 324
6A 463 807 10 0.0589 359
6B 771 817 5 0.0614 351
J. MALCHARCZIKOVÁ et al.: INFLUENCE OF THE HIP PROCESS ON THE PROPERTIES ...
16 Materiali in tehnologije / Materials and technology 49 (2015) 1, 15–18
Figure 2: Sample after the casting
Slika 2: Vzorec po ulivanju
Figure 3: Scheme of casting defects
Slika 3: Shematski prikaz livarskih napak
Figure 1: Graphite-mould drawing
Slika 1: Risba grafitne forme
3 DISCUSSION
We performed a comparison of the selected mecha-
nical properties in the as-cast state and after HIP. In the
majority of the investigated alloys the HIP process led to
an increase in the determined strength and elongation. It
is evident from the comparison of the values in Table 2
that a significant enhancement of the mechanical pro-
perties after the HIP process occurred in samples 3 and
5, while no significant improvement of these values took
place in samples 4 and 6. The Rm value of sample 4 does
not correspond to the expected values. The casting
probably contained a larger volume of casting defects
before the HIP process. Figure 4 shows the stress-strain
curves for samples 3 and 5 in the as-cast state and after
HIP. It may be concluded from the results that the
process of additional compaction of the castings with the
HIP method under the given conditions has a positive
influence.
Figures 5 to 8 show the structures of samples 3 and 5
in the as-cast state and after HIP. The HIP process
caused a coarsening of the grains as observed on Figures
5 and 7. This change in the structure was reflected by the
changes in the yield strength. The value of the yield
strength dropped for all the types of alloys (Table 2,
Figure 4). The gained results were already confirmed in
the conclusions of the papers of other authors4–15.
According to these results every nickel alloy reacts
differently to the HIP process. To avoid a decrease in the
mechanical properties of the castings, their grains should
be as small as possible and the microstructural changes
J. MALCHARCZIKOVÁ et al.: INFLUENCE OF THE HIP PROCESS ON THE PROPERTIES ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 15–18 17
Figure 4: Stress-strain curves for: a) samples No. 3 and b) samples
No. 5
Slika 4: Krivulje napetost – raztezek za: a) vzorca {t. 3 in b) vzorca {t. 5
Figure 7: Microstructure of sample 5A
Slika 7: Mikrostruktura vzorca 5A
Figure 6: Microstructure of sample 3B
Slika 6: Mikrostruktura vzorca 3B
Figure 5: Microstructure of sample 3A
Slika 5: Mikrostruktura vzorca 3A
of these alloys should be studied during the HIP process.
The porosity and micro-hardness are approximately the
same for both the as-cast state and the HIP state. How-
ever, for all the experimental alloys the value of porosity
decreased after the HIP process.
In practise it is possible to prepare the samples for
evaluating the selected mechanical properties using the
method of centrifugal casting of precision castings. The
use of precision castings as the test specimens brings
considerable savings of the materials and machining
costs. The alloys of this type are used for high-tempe-
rature applications and a modification of the alloy
composition is still performed16.
4 CONCLUSIONS
Modified Ni superalloys were prepared as the
experimental samples. The samples were molten by
vacuum-induction melting and then cast centrifugally
into a shaped graphite mould. The final castings had the
shape corresponding to the test specimens. After a
flaw-detection analysis, it was established that the
castings contained numerous casting defects. As a result,
some of the samples were subjected to hot isostatic
pressing (HIP). Hot isostatic pressing reduced the
number of the pores and casting defects in the castings.
We performed a comparison of selected mechanical
properties in the as-cast state and after HIP. For the
majority of the investigated alloys, the HIP process led to
an increase in the strength and ductility.
Acknowledgement
The presented results were obtained within research
project TA 01011128 "Research and development of
centrifugal casting technology of the Ni-based inter-me-
tallic compounds" and project CZ.1.05/2.1.00/01.0040
"Regional materials science and technology centre".
5 REFERENCES
1 A. Milosavljevic et al., The influence of the heat-treatment regime on
a fracture surface of nickel-based supperalloys, Mater. Tehnol., 46
(2012) 4, 411–417
2 M. Oru~ et al., Alloys with modified characteristics, Mater. Tehnol.,
45 (2011) 5, 483–487
3 C. García et al., Improvement of the ductility of a bimodal micro-
structured PM Ni3Al by thermal treatment, Materials Science and
Engineering A, 303 (2001), 11–18
4 J. G. Berriocanal et al., Microstructure and mechanical properties of
Ni3Al base alloy reinforced with Cr particles produced by powder
metallurgy, Intermetallics, 14 (2006), 456–463
5 V. A. C. Haanappel et al., The cyclic oxidation behavior of PM
Ni3Al in air at elevated temperatures: the effect of Cr- and Y-implan-
tation, Intermetallics, 6 (1998), 347–356
6 P. Pérez et al., Influence of the powder particle size on tensile pro-
perties of Ni3Al processed by rapid solidification and hot isostatic
pressing, Materials Science and Engineering A, 199 (1995), 211–218
7 H. K. Kim et al., High temperature deformation and fracture mecha-
nism in a dendritic Ni3Al alloy, Acta Metallurgica et Materialia, 42
(1994), 679–687
8 J. Malcharcziková et al., Non-alloyed Ni3Al based alloys – prepara-
tion and evaluation of mechanical properties, Metalurgija, 52 (2013),
309–312
9 C. L. Qui et al., Influence of hot isostatic pressing temperature on
microstructure and tensile properties of a nickel-based superalloy
powder, Materials Science and Engineering A, 564 (2013), 176–185
10 L. S. Ng et al., Cold-hot isostatic pressing of Mar M200 superalloy
powder, Journal of Materials Processing Technology, 67 (1997),
143–149
11 L. Chang et al., Effect of heat treatment on microstructure and me-
chanical properties of the hot-isostatic-pressed Inconel 718 powder
compact, Journal of Alloys and Compounds, 590 (2014), 227–232
12 C. Qiu et al., Influence of heat treatment on microstructure and ten-
sile behaviour of a hot isostatically pressed nickel-based superalloy,
Journal of Alloys and Compounds, 578 (2013), 454–464
13 M. T. Kim et al., Effect of high pressure on the solid-liquid phase
change of a nickel base superalloy during hot isostatic pressing, Jour-
nal of Alloys and Compounds, 477 (2009), 224–232
14 M. T. Kim et al., Effect of ’ precipitation during hot isostatic press-
ing on the mechanical property of a nickel-based superalloy,
Materials Science and Engineering A, 480 (2008), 218–225
15 H. Y. Bor et al., Influence of isostatic pressing on the fracture tran-
sitions in the fine grain MAR-M247 superalloys, Materials Chemi-
stry and Physics, 84 (2004), 284–290
16 J. Malcharcziková et al., Evaluation of structure characteristics of
centrifugally cast alloy based on Ni3Al, 21st International Confe-
rence on Metallurgy and Materials, Metal 2012, Brno, 2012,
1280–1287
J. MALCHARCZIKOVÁ et al.: INFLUENCE OF THE HIP PROCESS ON THE PROPERTIES ...
18 Materiali in tehnologije / Materials and technology 49 (2015) 1, 15–18
Figure 8: Microstructure of sample 5B
Slika 8: Mikrostruktura vzorca 5B
I. KUCUK, C. SARIOGLU: CORROSION OF CrN-COATED STAINLESS STEEL IN A NaCl SOLUTION ...
CORROSION OF CrN-COATED STAINLESS STEEL IN A NaCl
SOLUTION (w = 3 %)
KOROZIJA NERJAVNEGA JEKLA S CrN-PREVLEKO V
RAZTOPINI NaCl (w = 3 %)
Israfil Kucuk, Cevat Sarioglu
Marmara University, Dept. of Metallurgical and Materials Engineering, Göztepe kampusu, 34722 Kadiköy-Istanbul, Turkey
cevat.sarioglu@marmara.edu.tr
Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2014-02-12
A CrN coating deposited by arc PVD was characterized by XRD and SEM. The in-situ measurement of the corrosion of the
CrN-coated substrate was made by corrosion potential (Cor. Pot.), the polarization resistance (PR) method and electrochemical
impedance spectroscopy (EIS) in a NaCl solution w = 3 % as a function of the immersion time (about 24 h). A semiconductor
scale that formed on the CrN was identified by Mott-Shottky analysis as a p-type semiconductor with flat band potentials, 0.49
V (SCE). The CrN coating (0.5 μm thick) consisted of a mixture of cubic Cr and hexagonal Cr2N phases exhibiting equiaxed
grains and a dense coating with a small quantity of pinholes, voids and porosities. The “transition in corrosion resistance” for
the CrN coatings at an early stage was found based on Cor. Pot., PR and EIS data. The CrN did not exhibit any pitting for about
24 h, while the corrosion resistance (Rp and Rtotal) decreased rapidly with time after 5 h of incubation. The transition from high
resistance (3 M cm2) to low resistance (0.24 M cm2) was explained as a result of the penetration of the electrolyte through
the Cr2O3 oxide layer to the Cr2O3/CrN interface. The resistance of the CrN against pitting corrosion was explained based on the
blocking character of the equiaxed, dense, CrN coating against the penetration of the electrolyte.
Keywords: stainless steel, coating, CrN, EIS, polarization resistance, pitting corrosion
CrN-prevleka, nanesena z oblo~nim PVD-postopkom, je bila pregledana z XRD in SEM. In-situ merjenje korozije podlage s
CrN-prevleko je bilo izvr{eno z metodo korozijskega potenciala (Cor. Pot.), metodo polarizacijske upornosti (PR) in z elektro-
kemijsko impedan~no spektroskopijo (EIS) v raztopini NaCl (w = 3 %) v odvisnosti od ~asa namakanja (okrog 24 h).
Mott-Shottkyjeva analiza je odkrila nastanek p-vrste polprevodni{ke plasti s plo{~inskim potencialom 0,49 V (SCE).
CrN-prevleka (debela 0,5 μm) je sestavljena iz me{anice kubi~nega Cr in heksagonalne faze Cr2N ter ima enakoosna zrna, je
gosta, z majhno koli~ino luknjic, praznin in poroznostjo. Na podlagi podatkov Cor. Pot., PR in EIS je bil `e v zgodnji fazi
ugotovljen "prehod v obstojnost proti koroziji". CrN ni pokazal korozijskih jamic po 24 h, medtem ko se je korozijska upornost
(Rp in Rtotal) hitro zmanj{ala po 5 h inkubacijskega ~asa. Prehod iz velike upornosti (3 M cm2) v majhno (0,24 M cm2) se
razlaga kot rezultat penetracije elektrolita skozi sloj oksida Cr2O3 na mejo Cr2O3/CrN. Odpornost CrN proti jami~asti koroziji je
razlo`ena z zadr`evalnim zna~ajem goste, enakoosne prevleke CrN za penetracijo elektrolita.
Klju~ne besede: nerjavno jeklo, prevleka, CrN, EIS, polarizacijska upornost, jami~asta korozija
1 INTRODUCTION
Hard ceramic coatings such as CrN have mainly been
used for tribological applications (e.g., the injection
molds for plastic and metals and forming dies).1–4 CrN
coatings possess excellent wear and oxidation resistance
and low coefficients of friction. In addition, CrN has
been used for decorative applications (particularly for re-
placing decorative electroplating chrome coatings).5,6
Besides the excellent mechanical properties and
oxidation resistance, CrN coated on metallic substrates
(steel and stainless steels) possesses significant corrosion
resistance in saline solutions compared to other hard
coatings such as TiN and TiAlN.7–15 Even though the
corrosion resistance of a CrN coating is superior to other
coatings, the corrosion resistance of CrN-coated sub-
strates degrades with time.9,14,15 Generally, this degrada-
tion ends up with pitting corrosion by galvanic corrosion
between the substrate and the coatings.14,15 In the litera-
ture there has not been sufficient research into degra-
dation and pitting-corrosion mechanisms and the
relationships with microstructure.
Martensitic stainless steels (used in this work, EN
1.4034) generally used for blades in kitchen appliances,
were coated for both decorative (chrome appearance)
and wear-resistance requirements. The corrosion of the
CrN coating deposited by arc PVD on a stainless-steel
substrate was studied in detail with corrosion potential
(Cor. Pot), polarization resistance (PR) and electroche-
mical impedance spectroscopy (EIS) techniques. The
mechanism of corrosion was evaluated with respect to
the microstructure of the CrN coating and EIS mo-
delling.
2 EXPERIMENTAL PROCEDURES
The substrate material obtained from ThyssenKrupp
was EN 1.4034 (X46Cr13) stainless steel. The compo-
sition analysed by optical emission spectroscopy (Q4
Tasman Bruker) was 0.40 C, 0. 44 Si, 0.53 Mn, 14.05 Cr
and Fe (balance) (all in mass fractions, w/%). The 4034
stainless steel belonged to the martensitic group of
stainless steels and was commercially available annealed
and cold-rolled thin-plate form (1.5 mm thickness). The
Materiali in tehnologije / Materials and technology 49 (2015) 1, 19–26 19
UDK 669.14.018.8:621.793:620.193 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)19(2015)
specimens were cut into widths 50 mm and lengths 100
mm. All the specimens were electropolished for 10 min
in a solution containing 30 % H2SO4, 60 % H3PO4 and
10 % Neutral H2O at 50 °C with a voltage of 10 V, and
then ultrasonically cleaned at 60 °C, dumped in distilled
water and dried with nitrogen in order to obtain a clean
and smooth corrosion-resistant surface. The electro-
polished 4034 stainless-steel substrate is subsequently
referred to as "substrate" in this paper.
The coating of the substrates was performed in an
industrial-sized arc PVD coating chamber (AFS ltd.
Cop., Turkey) equipped with six evaporators (cathodes)
and rotating substrate holders. During the deposition,
four round cathodes (opposing each other) were in use.
An Ar glow discharge at a –600 V bias for 10 min was
formed in the chamber in order to clean the surface of
the substrates. The heating of the substrates was per-
formed with Cr cathodes at about a –800 V bias for
about 15 min until the starting temperature reached 350
°C. Then the electropolished specimens were coated with
a Cr interlayer for 1 min to improve the adhesion of the
coating and later with the CrN layer for 20 min at 1.1 ×
10–3 mbar nitrogen pressure and a total pressure of 10–2
mbar with a bias voltage of –200 V. The final deposition
temperature was 250 °C. The argon gas was purged into
the chamber before opening the chamber door.
Electrochemical corrosion units used to perform the
EIS, polarization resistance and Mott-Shottky scan were
a Gamry PC14/750 Poteniostat/Galvanostat/ZRA System
equipped with DC105 corrosion testing and an EIS300
A/C (10 mHz to 300 kHz). All the tests were performed
in a NaCl 0.5 M (w = 3 %) aerated water solution at
25 °C using a three-electrode system (working (sample),
auxiliary (graphite) and reference (standard calomel
electrode (SCE)) using a Gamry paint cell unit. This
Gamry Paint cell unit, presented in Figure 1 and used for
flat plate samples, consisted of a hollow glass cylinder
(capacity 50 mL) mounted on a plate sample using a
stainless-steel clamp with an O-ring. The exposed
surface area of all the samples was kept at 15 cm2, the
maximum value without using a mask in order to incor-
porate a statistically significant number of defects in the
CrN coatings. The computer-controlled automatic
measurements (corrosion potential, polarization resi-
stance and EIS) were carried out in the following order:
10 min cor. pot; 2.7 min PR; 5 min waiting period; and 8
min EIS meas. for each cycle 68 min during a total of
about 24 h of exposure after the corrosion potential was
stable. The amplitude of the AC voltage was selected to
be 10 mV with respect to the corrosion potential in the
EIS. During the polarization resistance scan the corro-
sion potential was changed by 20 mV with a 0.25 mV/s
scan rate from the cathodic to the anodic polarization
directions. Mott-Shottky analyses at a frequency of 1 Hz
between +430 mV and –570 mV were performed on a
CrN-coated substrate after the corrosion potential was
stable.
A microstructural analysis of the uncoated and coated
samples was performed before and after the corrosion
using a scanning electron microscope (SEM, Jeol, JSM-
5910LV) and energy-dispersive spectroscopy (EDS). The
cross-section of the coatings was observed after fracture
in liquid nitrogen. X-ray diffraction (Rigaku, D-MAX
2200, Cu K
 radiation) was used to identify the structure
of the coatings deposited on the substrate.
3 RESULTS
3.1 Microstructural characterization of CrN coated
substrate
The 4034 EP stainless-steel substrate possessed a
ferritic structure (
-Fe) matrix (Figure 2) with a homo-
genous distribution of the carbide phases (Cr and C-rich
phase identified by EDS). After coating with CrN at
1.1 × 10–3 mbar of N2 partial pressure, the surface mor-
phology of the CrN coatings reflected the morphology of
I. KUCUK, C. SARIOGLU: CORROSION OF CrN-COATED STAINLESS STEEL IN A NaCl SOLUTION ...
20 Materiali in tehnologije / Materials and technology 49 (2015) 1, 19–26
Figure 2: 2 scan obtained from EP 4034 steel substrate and CrN-
coated substrates using Bragg-Brentano symmetric X-ray diffraction
Slika 2: Posnetek 2 z Bragg-Brentanovo simetri~no rentgensko di-
frakcijo podlage iz jekla EP 4034 in iz podlage s prevleko CrN
Figure 1: Gamry paint cell used in this work for electrochemical
measurements
Slika 1: Elektrokemijske meritve so bile izvr{ene z Gamryjevo celico
za preizku{anje barvnih premazov
the EP substrate surface where etched carbide phases
were covered by the CrN coating (Figure 3). The same
observations were made for the cross-section micro-
graphs of the CrN-coated substrates where the coatings
formed on the carbide phases represented a hemispheric
depression (Figure 4). Based on a detailed surface and
cross-section investigation, it was found that the CrN
coatings (Figures 3 and 4) exhibited a dense coating
morphology (a small quantity of pinholes and porosities)
with equiaxed grains (less than 100 nm grain size) and
smooth coverage over the carbide precipitates and
matrix. The thickness of the CrN coatings measured
from the cross-section (Figure 4) was about 0.5 μm.
Due to the arc PVD process, droplets formed on the
surface of the CrN-coated substrate. There were two
different types of droplets found on the surface of the
CrN coatings, the droplets embedded in the scale, the
coating and the un-embedded macro-particles (spherical
bright particles) (Figures 3 and 4). They were believed
to deposit on the surface as a result of vapour-phase
precipitation after the coating was finished (when a bias
was interrupted). Both types of droplets analysed by
EDS contained mainly Cr and N (Cr-rich particles). The
CrN coating contained cubic Cr and hexagonal Cr2N
phases. The Cr peak exhibited line broadening and a
shift in the line position of the (200) peak due to the
dissolution of nitrogen in the Cr and a small crystallite
size (Figure 2).
3.2 In-situ measurement of the corrosion of the
CrN-coated substrate during 24 h of exposure with
corrosion potentials, PR- and EIS-measurements
The corrosion potentials were plotted as a function of
exposure time in Figure 5. The corrosion potential of the
substrate (–468 mV (SCE) at 20 min) dropped rapidly
and reached a steady-state value of –620 mV (SCE)
within 180 min (3 h) (Figure 5). The corrosion potential
of the CrN remained at around –150 mV (SCE) up to
267 min before decreasing to a level close to the sub-
strate’s potential value (Figure 5). At all times the corro-
sion potential of the substrate was lower compared to the
I. KUCUK, C. SARIOGLU: CORROSION OF CrN-COATED STAINLESS STEEL IN A NaCl SOLUTION ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 19–26 21
Figure 3: SEM (SEI) micrographs taken from the CrN coating surface
at: a) low magnification and b) high magnification. The CrN coating
deposited on the carbide phases was distinguished with a depression
over the surface at a) and b) and was marked at b). Embedded droplets
were marked at a). Bright particles on coating surface at a) and b)
were un-embedded droplets (spherical particles).
Slika 3: SEM (SEI)-posnetka povr{ine prevleke CrN pri: a) majhni
pove~avi in b) veliki pove~avi. CrN-prevleke na karbidnih fazah so
lo~ene s poglobitvijo na povr{ini a) in b) in ozna~ene na b). Vgne-
zdene kapljice so ozna~ene na a). Svetli delci na povr{ini prevleke
a) in b) so nevgnezdene kapljice (sferi~ni delci).
Figure 4: SEM (SEI) micrographs of fractured surface (in liquid
nitrogen) of CrN coating from different areas a) and b). Equiaxed
grains through the fractured coating were seen at high magnification
b). Embedded particles were marked at b). Hemispheric depression in
fractured coating was shown at a).
Slika 4: SEM (SEI)-posnetka povr{ine preloma (v teko~em du{iku)
CrN-prevleke z razli~nih podro~ij a) in b). Pri veliki pove~avi so vidna
enakoosna zrna na prelomu prevleke. Vgnezdeni delci so ozna~eni na
b). Polkro`na poglobitev na prelomu prevleke je prikazana na a).
CrN coating during about 24 h of exposure. The surface
appearance of the uncoated substrate and the CrN-coated
substrate did not vary during the 24 h of exposure.
The polarization-resistance values (Rp) were deter-
mined using the polarization resistance method and
plotted (Figure 5). The polarization resistance (Rp) of the
CrN coatings (Figure 5), deduced from the plots in Fig-
ure 6, remained at a high level of 3 M cm2 until 267
min and then rapidly decreased from 3 M cm2 to 0.24
M cm2 within the next 272 min. This decrease could be
detected in the polarization plots (expansion of the
x-axis) (Figure 6). After 1491 min the polarization resis-
tance decreased to 66 k cm2. The plateau in the
polarization resistance up to 267 min matched with the
plateau in the corrosion potentials in the same exposure
range (Figure 5). The polarization resistance (Rp) of the
substrate increased up to 27 k cm2 with time (Figure 5).
After about 24 h the Rp of the CrN coating was two times
greater than the Rp of the substrate (Figure 5).
The Bode and Nyquist plots of the CrN coatings from
the EIS measurements are shown in Figure 7. At the
beginning of the immersion in the salt solution, the Bode
and Nyquist plots exhibited a strong capacitive beha-
viour. In the Bode plot (Figure 7b), the plateau in the
phase shift (at about –80°) extended to the whole low-
frequency region and the almost perpendicular line in the
Nyquist plot indicated a strong imaginary part of the
impedance (a strong capacitive character) (Figure 7a).
Between 275 min and 411 min, the Bode and Nyquist
plots changed noticeably, the plateau in the low-fre-
quency region shrank and the Nyquist plot’s shape
changed from a line to an arch (Figure 7a). The
evolution of the impedance spectrum coincided with the
change in the corrosion potential (Figure 5) and the
polarization resistance during the first 411 min of
exposure in the salt solution (Figures 5 and 6). The one
time constant clearly appeared in the Bode plot (Figure
7b). The imaginary impedance values decreased,
resulting in the shrinkage of the arch in the Nyquist plot
(Figure 7a). During 24 h of exposure to the salt solution
there were no pits and no colour changes on the surface
I. KUCUK, C. SARIOGLU: CORROSION OF CrN-COATED STAINLESS STEEL IN A NaCl SOLUTION ...
22 Materiali in tehnologije / Materials and technology 49 (2015) 1, 19–26
Figure 7: EIS data of CrN coating for selected immersion time: a)
Nyquist plots and b) the Bode plots
Slika 7: EIS-podatki za CrN-prevleko pri izbranih ~asih namakanja: a)
Nyquistov diagram in b) Bodejev diagram
Figure 6: Polarization resistance scan of CrN coating in NaCl (w = 3
%) solution as a function of immersion time
Slika 6: Posnetek polarizacijske upornosti CrN-prevleke v raztopini
NaCl (w = 3 %) v odvisnosti od ~asa namakanja
Figure 5: Corrosion potentials and polarization resistance (Rp) of the
substrate and CrN-coated substrates determined from polarization
resistance scan in NaCl (w = 3 %) solution as a function of immersion
time
Slika 5: Korozijski potenciali in polarizacijska upornost (Rp) podlage
in CrN-prevleke, dolo~eni s posnetka polarizacijske upornosti v raz-
topini NaCl (w = 3 %) v odvisnosti od ~asa namakanja
of the CrN coatings, while the impedance values con-
tinued to decrease (Figure 7).
3.3 The Mott-Shottky measurements
The Mott-Shottky measurements were made for the
CrN coating and plotted in Figure 8. Before the Mott-
Shottky measurement was made the corrosion potential
was stable and the EIS data indicated a strong capacitive
response at an early stage of the exposure (after 60 min).
The plot for CrN (Figure 8) showed a linear part with a
negative slope between 110 mV and 435 mV (SCE). The
negative slope implied that there was a p-type semicon-
ductor oxide layer on the CrN surface.
4 DISCUSSION
4.1 Structure of CrN coating
The CrN coating exhibited a mixture of two phases of
Cr and Cr2N in the N2 partial pressure of 1.1 × 10–3 mbar
(Figure 2). Significant peak broadening of the (200)
plane and a shift in the peak position were observed for
Cr. Similar results were found by Barata.16 With
increasing nitrogen content, Cr, Cr2N, CrN and mixtures
of these phases were obtained, based on reports in16–21.
The Cr phase was soft, while the Cr2N phase was
hard.16,18,22 The CrN coatings in this work exhibited
significant ductility, as can be seen from the fractured
cross-section (Figure 4). Droplets, considered as a prefe-
rential site for pitting in8, were found on the CrN coating
(Figures 3 and 4). The CrN coating exhibited a very
dense and very smooth surface over the substrate,
including etched carbide phases and droplets, resulting in
less porosity and fewer pinholes. Thus, the CrN coating
could effectively prevent the electrolyte from reaching
the CrN coating/substrate interface directly.
The development of an equiaxed grain morphology
(Figure 4) requires significant mobility of the adatoms
during the growth of the coating according to the struc-
ture zone model (SZM).8 The equiaxed grain morpho-
logy was observed for similar coating conditions in6,14,15.
In addition, from the surface and cross-section micro-
graphs (Figures 3 and 4) it was clear that the CrN
coating was dense and provided smooth coverage over
the etched carbide phases, indicating significant surface
mobility of the adatoms. In the structure Zone Model,
the TS/TM ratio, where TS is the substrate deposition
temperature (taken as an average of 300 °C) and TM is
the melting temperature of the coating, determines the
morphology of coatings (equiaxed vs. columnar)
together with the bias voltage and the total pressure.3 As
the ratio increases the coating morphology becomes
more equiaxed due to an increase in the surface
mobility.3 When the melting point of the CrN coatings is
taken as TM (Cr) = 1907 °C 23 and TM (Cr2N) = 1500
°C 24, the ratio of TS/TM for CrN is in the range
0.26–0.32. For this ratio range we expected to observe
significant surface mobility for the CrN coatings
according to SZM3, supporting the observation of
equiaxed grains, and the denser and smoother coverage
of the CrN coating.
4.2 Mott-Shottky analysis of CrN coatings
In the literature, the Mott-Shottky analysis has been
used to characterize the semiconductor layer formed on
the surface of the materials and coatings.9,25–30 Based on
Figure 8, it was concluded that the semiconductor layer
formed on the CrN-coated substrate was p-type (pre-
sumably Cr2O3). In agreement with this result, Lavigne31
and Mendibide9 found p-type Cr2O3 on the CrN coating
deposited by PVD on glass substrates.
The Mott-Shottky equation on page 127 in25 was used
to determine the flat band potential and the density of the
charge (the density of acceptors for the p-type semi-
conductor) in the space-charge region. After taking the
dielectric constant of the Cr2O3 as 30 (cited in26 as ref.
17), the flat band potentials and the density of charges
were determined from the linear portion of the plot in
Figure 8 using the equation in25. The density of accep-
tors in the space-charge region in the p-type Cr2O3 was
2.14 × 1025 cm–3. Lavigne31 obtained a density of the
acceptors of 1.3 × 1020 cm–3 for the p-type oxide scale
(mainly Cr2O3) formed on the CrN coating in a solution
of Na2SO4 0.07 M and a small addition of H2SO4 with
pH = 4. The Cr2O3 scales with the high charge densities
(1020–25 cm–3) in the space-charge layer were considered
as highly doped semiconductor layers.25 The flat bond
potentials indicated the position of the valance-band
edge for the p-type Cr2O3 since the Fermi levels were
generally close to the valance-band edge for the p-type.25
The calculated flat band potentials from the intercept of
the plots (Figure 8) was +0.49 V (SCE) for the p-type
Cr2O3. The flat band potential reported for the oxide
scale for a similar coating and nitrogen pressure by
Lavigne31 was in agreement with this result after con-
sidering the pH effect.
I. KUCUK, C. SARIOGLU: CORROSION OF CrN-COATED STAINLESS STEEL IN A NaCl SOLUTION ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 19–26 23
Figure 8: The Mott-Shottky measurements made for CrN coatings at
frequency of 1 Hz in NaCl (w = 3 %) solution
Slika 8: Mott-Shottkyjeva meritev CrN-prevleke pri frekvenci 1 Hz v
raztopini NaCl (w = 3 %)
4.3 Corrosion of CrN-coated substrate and EIS mo-
delling
Even though there was no colour change or pitting on
the CrN-coated substrate during immersion, the transi-
tion in PR and EIS data (this transition was reproduced)
from more resistance to less resistance at the early stage
between 275 min and 411 min was found simultaneously
using the EIS and PR measurements (Figures 5 and 7).
The EIS data was modelled as a one-time constant
(Figure 7a), the same as the EIS model for the substrate
since there was no pitting corrosion and the EIS data
exhibited one time constant (Figure 7). The total polari-
zation resistance (Rtotal = Rpassive + Rpore) included the
resistance of the passive layer (Rpassive) and the pore
resistance (Rpore) to where the electrolyte gained access.
The fitted EIS data are presented in Figure 7 and Table
1. The 2 values were in the range of 2–10 × 10–3, while
the n value was about 0.88 and remained constant with
time. The polarization resistance (Rp) and the total
resistance (Rtotal) determined from the EIS data were
almost the same (Figure 9). The Rtotal values dropped
rapidly from about 2.62 M cm2 levels to 0.33 M cm2
levels between 275 min and 411 min (in about 2 h)
(Figure 9). In the same time interval, the admittance, Y0
(Figure 10 and Table 1), decreased from 34.6 μcm–2 sn 
to 29.7 μcm–2 sn  by 14 %, indicating an increase in the
capacitive behaviour of the passive layer (Cr2O3). The
total resistance of the capacitive layer Rtotal and Rp
continued to decrease gradually after 411 min until the
end of the exposure (1499 min) to the level of 60 k cm2
without any observation of pitting on the CrN surface
(Figures 5 and 9). The corrosion potential also followed
the same trend with the EIS and PR results (Figures 5
and 9).
Table 1: The EIS fit parameters of the CrN coating from the electrical
circuit model (EC model); Y0, n, Rsol. and Rtotal were determined by
the best non-linear least-square fit with a goodness of the fit value (2)
Tabela 1: EIS-parametri CrN-prevleke iz modela elektri~nega toko-
kroga (EC-model); Y0, n, Rsol. in Rtotal so bili dolo~eni z najbolj{im
ujemanjem nelinearnih najmanj{ih kvadratov z vrednostjo ujemanja
(2)
Immer-
sion time
(min)
Rsol/
( cm2)
Rtotal/
(k cm2)
Y0/
(μcm–2 sn ) n
Goodness
of Fit
(2 × 10+3)
71 10.6 3457.5 42.8 0.86 2.143
275 10.6 2632.5 34.6 0.87 4.065
343 10.6 1054.5 33.1 0.87 4.014
411 10.6 340.5 29.7 0.88 3.299
819 10.7 195 24.6 0.88 2.72
1091 10.7 109.5 22.7 0.88 10.18
1499 10.6 63 23 0.88 2.237
It was believed that a rapid drop in the corrosion
potential and the polarization resistance (Rp and Rtotal) of
the passive oxide (Cr2O3) between 275 min and 411 min
(in the transition period) (Figure 9) was due to penetra-
tion of the electrolyte into the CrN coating/Cr2O3 inter-
face. Once the electrolyte reached this interface, the
galvanic coupling between the passive layer (Cr2O3) and
the bare mixture of the Cr and Cr2N surfaces (CrN sur-
face) proceeded and polarization on both areas (the
cathodic polarization on the Cr2O3 scale, the anodic
polarization on the bare CrN coatings) took place. The
new corrosion potential established (more cathodic, more
negative) as observed (Figures 5 and 9).
After 411 min, the polarization resistance (Rp and
Rtotal) and the corrosion potential continued to drop
gradually (Figures 5 and 9). This could be explained as
follows. The anodic areas of the bare exposed surface in
the CrN scale increased with the penetration of the
electrolyte into the Cr2O3/CrN interface and towards the
substrate/CrN coating interface through equiaxed grain
boundaries, pinholes and embedded droplets. The
increase in the anodic areas increased the anodic currents
I. KUCUK, C. SARIOGLU: CORROSION OF CrN-COATED STAINLESS STEEL IN A NaCl SOLUTION ...
24 Materiali in tehnologije / Materials and technology 49 (2015) 1, 19–26
Figure 9: Corrosion potential (Cor. Pot.), polarization resistance (Rp
and Rtotal calculated from EIS data) for the CrN coating in NaCl (w =
3 %) solution as a function of immersion time
Slika 9: Korozijski potencial (Cor. Pot.), polarizacijska upornost (Rp
in Rtotal, izra~unani iz podatkov EIS) CrN-prevleke v raztopini NaCl
(w = 3 %) v odvisnosti od ~asa namakanja
Figure 10: Admittance (Y0) for the CrN coating in NaCl (w = 3 %)
solution as a function of immersion time
Slika 10: Admitanca (Y0) za CrN prevleko v raztopini NaCl (w = 3 %)
v odvisnosti od ~asa namakanja
and cathodic currents, resulting in a more cathodic pola-
rization of the Cr2O3 surface as bare surfaces of CrN
coating continued to be exposed to the electrolyte with
the immersion time. As a result of this, the polarization
resistance and corrosion potentials decreased, as was
observed (Figures 5 and 9).
In aerated salt solution (w(NaCl) = 3 %) with pH = 8
and at the corrosion potential range (Figure 5), the
Pourbaix diagram for the Cr indicated that the Cr could
exhibit passivation and the anodic polarization of surface
could result in passivation.32 Mixture phases (Cr and
Cr2N) of the CrN coating were thought to behave in a
similar way to the pure Cr phase when the surface was
anodically polarized in the saline solution. In support of
this, Lavigne31 observed a passive layer on the Cr and Cr
+ Cr2N coating in a pH = 4 solution. Passivation could
continue with time until the substrate containing Fe is
reached. The electrolyte was believed to not yet reach the
substrate/CrN interface since no pit or Fe oxide residue
(brownish colour residue) on the surface after 24 h was
observed.
The Y0 (measure of capacity) value decreased with
time logarithmically up to 1091 min (Figure 10). The
capacity (C) decreases either with a decrease in the
dielectric constant of the Cr2O3 passive film and the area
of the passive film or an increase in the thickness of the
Cr2O3 oxide layer. Both the area of the passive layer
could increase (as explained above) and the passive film
of Cr2O3 on Cr or Cr2N coating could grow in this solu-
tion of pH = 8 during immersion in a 3 NaCl (w = 3 %)
solution.
The corrosion resistance of the CrN was greater than
the substrate during about 24 h of exposure (Figure 5),
indicating more resistance to the passive scale on the
CrN-coated substrate in the NaCl (w = 3 %) solution.
Also, the corrosion resistance of the CrN coating was
more resistance than that of the TiN coating under simi-
lar deposition conditions.33 Similar results were reported
in the literature for different PVD processes and coating
conditions.8–15 There were two main differences in the
microstructural morphologies between the TiN and CrN
coatings for the same deposition condition.14,15,33 First,
the CrN exhibited equiaxed grains (Figure 4), while the
TiN coating was columnar. Second, the CrN coating was
denser (contained a smaller quantity of pinholes and
fewer porosities) compared to the TiN coating. The CrN
coatings with equiaxed grains were believed to suppress
the time for the electrolyte to reach the substrate
interface by increasing the diffusion length (more grain
boundaries) (Figure 4).14,15 In addition to this, with
denser and better coverage of the CrN coatings over the
substrate, it resisted more against the penetration of the
electrolyte, delaying the pitting corrosion (Figures 3 and
4). On the other hand, the TiN coating with columnar
grain boundaries extending from the surface to the sub-
strate interface and containing more porosity (pinholes,
voids, cracks and less coverage) allowed the electrolyte
to reach the interface easily. As a result of this, the
pitting formation and growth as captured by the PR and
EIS measurements took place on the TiN coating as
quickly as after 1 h. In conclusion, the blocking charac-
ter of the coating against the penetration of the elec-
trolyte driven by galvanic corrosion was found to be the
predominant factor in the corrosion of CrN and TiN.
5 CONCLUSIONS
1. The corrosion of CrN is directly related to the coating
defects and coating structure. The CrN coating de-
posited on the substrate consisted of a mixture of
cubic Cr and hexagonal Cr2N and a passive p-type
oxide (presumably Cr2O3) film, determined by Mott-
Shottky analysis. The CrN coating consisted of
equiaxed grains and a smooth coverage over the sub-
strate surface, resulting in a dense coating morpho-
logy with a small quantity of pin holes and porosities.
2. The CrN did not exhibit any pitting for about 24 h,
while the corrosion resistance (Rp and Rtotal) de-
creased rapidly with time after 5 h of incubation
time. Based on the PR and EIS results, the transition
from high resistance (3 M cm2) to low resistance
(0.24 M cm2) was explained as a result of the
penetration of the electrolyte through the Cr2O3 oxide
layer to the Cr2O3/CrN interface.
3. The corrosion resistance (Rp and Rtotal) of CrN was
greater than that of the TiN and the substrate during
about 24 h of immersion in NaCl (w = 3 %) solution.
The cause of the greater corrosion resistance of the
CrN was explained on the basis of the greater block-
ing character of the equiaxed, dense, CrN coating
against the penetration of the electrolyte through the
coatings compared to the columnar, porous, TiN
coating.
Acknowledgement
Marmara University is greatly acknowledged for
financial support through the Contract (No: FEN-KPS-
080808-0178).
6 REFERENCES
1 B. Warchholinski, A. Gilewicz, Surface Engineering, 27 (2011) 7,
491–497
2 E. Lugscheider, K. Bobzin, Th. Horning, M. Macs, Thin Solid Films,
420–421 (2002), 318–323
3 C. Ducros, V. Benevent, F. Sanchette, Surface and Coatings Tech-
nology, 163–164 (2003), 681–688
4 L. Hultman, J. E. Sundgren, Structure/property relationships for hard
coatings, In: R. F. Bunshah (ed.), Handbook of Hard Coatings,
William Andrew Publishing, New York 2001, 108–180
5 B. Navinsek, P. Panjan, I. Milosev, Surf. Cot. Tech., 116–119 (1999),
476–487
6 A. Hurkmans, D. B. Lewis, W. D. Munz, Surface Engineering, 19
(2003) 3, 205–210
7 H. A. Jehn, Surface and Coatings Technology, 125 (2000), 212–217
I. KUCUK, C. SARIOGLU: CORROSION OF CrN-COATED STAINLESS STEEL IN A NaCl SOLUTION ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 19–26 25
8 M. Fenker, M. Balzer, H. Kappl, Thin Solid Films, 515 (2006) 1,
27–32
9 C. Mendibide, P. Steyer, J. P. Millet, Surface and Coatings
Technology, 200 (2005), 109–112
10 M. Urgen, A. F. Cakir, Surface and Coatings Technology, 96 (1997),
236–244
11 M. A. M. Ibrahim, S. F. Korablov, M. Yoshimura, Corrosion Science,
44 (2002), 815–828
12 V. K. William Grips, H. C. Barshilia, V. Ezhil Selvi, Kalavati, K. S.
Rajam, Thin Solid Films, 514 (2006), 204–211
13 L. Cunha, M. Andritschky, L. Rebouta, R. Silva, Thin Solid Films,
317 (1998), 351–355
14 C. Liu, Q. Bi, A. Leyland, A. Matthews, Corrosion Science, 45
(2003), 1243–1256
15 C. Liu, Q. Bi, A. Leyland, A. Matthews, Corrosion Science, 45
(2003), 1257–1273
16 A. Barata, L. Cunha, C. Moura, Thin Solid Films, 398–399 (2001),
501–506
17 E. Forniés, R. Escobar Galindo, O. Sánchez, J. M. Albella, Surface
and Coatings Technology, 200 (2006), 6047–6053
18 S. K. Pradhan, C. Nouveau, T. A. Vasin, M. A. Djouadi, Surface and
Coatings Technology, 200 (2005), 141–145
19 Z. B. Zhaoa, Z. U. Rekb, S. M. Yalisovec, J. C. Bilell, Surface and
Coatings Technology, 185 (2004), 329–339
20 B. C. Ruden, E. Restrepo-Parra, A. U. Paladines, F. Sequeda,
Applied Surface Science, 270 (2013), 150–156
21 P. Hones, R. Sanjines, F. Levy, Surface and Coatings Technology,
91–95 (1997), 398–402
22 P. Engel, G. Schwarz, G. K. Wolf, Surface and Coatings Technology,
98 (1998), 1002–1007
23 R. T. DeHoff, Thermodynamics in materials science, McGraw-Hill,
Inc., 1993
24 W. Kim, C. Y. Suh, K. M. Roh, J. W. Lim, S. L. Du, I. J. Shan, Mate-
rials Transactions, 53 (2012) 8, 1543–1546
25 S. Roy Morrison, Electrochemistry at Semiconductor and Oxidized
Metal Electrodes, Plenum Press, New York 1980
26 S. Rudenja, C. Leygraf, J. Pan, P. Kulu, E. Talimets, V. Mikli, Sur-
face and Coatings Technology, 114 (1999), 129–136
27 J. Wielant, V. Goossens, R. Hausbrand, H. Terryn, Electrochimca
Acta, 52 (2007), 7617–7625
28 K. S. Raja, D. A. Jones, Corrosion Science, 48 (2006), 1623–1638
29 N. E. Hakiki, M. F. Montemor, M. G. S. Ferreira, M. da Cunha Belo,
Corrosion Science, 42 (2000), 687–702
30 G. Goodlet, S. Faty, S. Cardoso, P. P. Freitas, A. M. P. Simoes, M. G.
S. Ferreira, M. Da Cunha Belo, Corrosion Science, 46 (2004),
1479–1499
31 O. Lavigne, C. Alemany-Dumont, B. Normand, M. H. Berger, C.
Duhamel, P. Dekichere, Corrosion Science, 53 (2011), 2087–2096
32 D. A. Jones, Principles and Prevention of Corrosion, Prentice hall,
Inc., 1992
33 I. Kucuk, Corrosion of hard ceramic (TiN, CrN) coatings deposited
on stainless steels, Master Thesis, Sakarya University, Turkey, 2009
I. KUCUK, C. SARIOGLU: CORROSION OF CrN-COATED STAINLESS STEEL IN A NaCl SOLUTION ...
26 Materiali in tehnologije / Materials and technology 49 (2015) 1, 19–26
J. JUØICA et al.: PREPARATION AND PROPERTIES OF MASTER ALLOYS Nb-Al AND Ta-Al FOR MELTING ...
PREPARATION AND PROPERTIES OF MASTER ALLOYS Nb-Al
AND Ta-Al FOR MELTING AND CASTING OF -TiAl
INTERMETALLICS
PRIPRAVA IN LASTNOSTI PREDZLITIN Nb-Al IN Ta-Al ZA
TALJENJE IN ULIVANJE INTERMETALNIH ZLITIN -TiAl
Jan Juøica, Tomá{ ^egan, Kateøina Skotnicová, Daniel Petlák, Bedøich Smetana,
Vlastimil Matìjka
V[B – Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of RMSTC, 17. listopadu 15,
70833 Ostrava-Poruba, Czech Republic
jan.jurica@vsb.cz
Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2014-03-10
The advantages of gamma TiAl-based alloys including their specific modulus, specific high-temperature strength and oxidation
resistance make them attractive candidates as high-temperature structural materials in the automotive, aerospace and power
industries. Currently most attention is paid to the alloys of the third and fourth generations. However, this type of alloys contains
relatively high amounts of refractory metals such as Nb and Ta. The high melting points of these metals (2477 and 3017) °C are
problematic for the preparation of these products with the conventional casting, because it is necessary to use higher
temperatures and thus, generally, longer total melting times. This may result in increased oxygen amounts in the products and in
decreased mechanical properties. The use of Nb-Al and Ta-Al master alloys for the preparation of the resulting Ti-Al-Nb and
Ti-Al-Ta alloys is highly suitable because of the reduction in the temperature during melting.
This article describes the preparation of selected master alloys Nb-60Al and Ta-80Al (x/%) with the melting points of about
1600–1650 °C using plasma melting. The optimum conditions for the preparation of these master alloys (current density, feed
speed, distribution and size of charge) were characterised in order to maximise the purity and homogeneity. The prepared alloys
were studied with light microscopy (LM), backscattered scanning electron microscopy (BSE), energy-dispersive spectrometry
(EDS), and the melting temperature was evaluated with a differential thermal analysis (DTA).
Keywords: intermetallics, plasma melting, microstructure, differential thermal analysis (DTA)
Prednost zlitin na osnovi gama Ti-Al je v specifi~nem modulu, specifi~ni visokotemperaturni trdnosti in odpornosti proti
oksidaciji, kar jih dela zanimive za izdelavo visokotemperaturnih komponent v avtomobilski industriji, letalstvu in energetiki.
Sedaj se najve~ pozornosti namenja zlitinam tretje ali ~etrte generacije. Ta vrsta zlitin vsebuje relativno veliko koli~ino ognje-
varnih kovin, kot sta Nb in Ta. Visoko tali{~e teh materialov (2477 in 3017) °C ote`uje izdelavo teh zlitin s klasi~nim ulivanjem,
ker je treba uporabiti vi{je temperature in s tem dalj{e ~ase taljenja. To lahko povzro~i pove~anje vsebnosti kisika v proizvodih
in poslab{anje mehanskih lastnosti. Uporaba predzlitin Nb-Al in Ta-Al za izdelavo zlitin Ti-Al-Nb in Ti-Al-Ta je zelo primerna
zaradi zni`anja temperature pri taljenju.
^lanek opisuje pripravo izbranih predzlitin Nb-60Al in Ta-80Al (x/%) s tali{~em okrog 1600–1650 °C s taljenjem v plazmi. Za
maksimiranje ~istosti in homogenosti so bili dolo~eni optimalni pogoji za pripravo teh predzlitin (gostota, hitrost dodajanja,
razporeditev in velikost zatehte). Pripravljene zlitine so bile preiskane s svetlobno mikroskopijo (LM), vrsti~no elektronsko
mikroskopijo s povratno sipanimi elektroni (BSE), energijsko disperzijsko spektroskopijo (EDS), temperatura taljenja pa je bila
dolo~ena z diferen~no termi~no analizo (DTA).
Klju~ne besede: intermetalne zlitine, taljenje s plazmo, mikrostruktura, diferen~na termi~na analiza (DTA)
1 INTRODUCTION
The binary diagrams shown in Figures 1a and 1b
demonstrate the objective of this experiment, e.i., how to
transform refractory Ta or Nb into the master alloys,
suitable for the preparation of the -TiAl intermetallics
of the third or fourth generations that usually contain
5–10 % of these metals1–4. Suitable master alloys should
have the melting temperature of approximately 1600 °C,
high purity and low amount of interstitial elements,
particularly oxygen. The chosen chemical compositions
of the master alloys were Nb-60Al and Ta-80Al
(amount-of-substance fractions, x/%). As it can be seen
in the binary diagrams, the master alloys with these
compositions are characterised by both an acceptable
melting temperature and a chemical composition that can
be considered as suitable for the preparation of the
-TiAl alloys with respect to an easy preparation of the
charges.
2 EXPERIMENTAL WORK
For the preparation of the Al-Nb and Al-Ta alloys we
used the possibility of re-melting in a plasma furnace
with a horizontal mould. The charge consisted of a lump
material including pure metals in the form of plates,
wires and rods. The charge was uniformly distributed in
a water-cooled copper mould placed in the furnace. The
melting was performed using a quadruple passage
through the zone (twice on each side) for the Nb-60Al
alloy and a sextuple passage through the zone (three
times on each side) for the Ta-80Al alloy. An inert
Materiali in tehnologije / Materials and technology 49 (2015) 1, 27–30 27
UDK 669.017.12:621.74.046:66.046.51 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)27(2015)
atmosphere was ensured during the melting using an
argon flow with a flow rate of 27 L/min through the
furnace chamber. The rate of shift of the mould was
0.5 cm/min and the current intensity was up to 800 A.
The samples for metallographic observations were
prepared with the standard metallographic method
including the initial grinding on the sandpapers with the
granularities from 60 to 2000, and polishing with an
Al2O3 suspension with the particles of the sizes from 1
μm to 0.3 μm. For the observation of the microstructure
and the analysis of the chemical composition we used a
scanning electron microscope QUANTA FEG 450
equipped with an EDAX APOLLO X probe. The sam-
ples for the X-ray diffraction analysis were prepared by
cutting them in an electro-spark cutter. X-ray diffraction
patterns were obtained with a diffractometer Bruker D8
Advance equipped with a detector VÅNTEC 1. The
measurements were performed in the reflection mode.
The phase composition was evaluated using the PDF-2
2004 database from the International Centre for Diffrac-
tion Data. The oxygen amount was measured with the
thermo-evolution method using an ELTRA ONH–2000
instrument. For the measurement we always used at least
three pieces from each sample. The DTA analysis was
performed with the experimental equipment Setaram
SetSys 1750. The samples with the dimensions of 3 mm
× 3 mm × 3 mm, prepared by electro-spark cutting and
by wet grinding them on the sandpaper with a granula-
rity of 600 were analysed in corundum crucibles with an
internal coating of Y2O3 at a heating rate of 10 °C/min
and in an atmosphere of argon with the purity of 6 N.
3 RESULTS AND DISCUSSION
Table 1 shows the chemical compositions, deter-
mined with the EDS method, of individual sections of
the prepared ingots of the master alloys. The sections
were always taken from the front and rear parts of an
ingot. It was established that a slight loss of mass
occurred during the melting, which was probably a result
of the evaporation of the Al part, due to the high tem-
perature of the melt (the melting point of Ta is 3017 °C,
the evaporation temperature of Al is 2519 °C). A similar
phenomenon was also observed during the melting of
alloy Ti46Al8Ta5 and during the melting of -TiAl inter-
metallics in ceramic crucibles6. This loss of mass was
approximately 1 g for alloys NbAl-1 and NbAl-2, and
2 g for alloy TaAl-1. The loss of mass was probably a
result of the evaporation of the Al part due to a very high
temperature of the melt. A higher loss of mass in the
case of master alloy TaAl-1 also corresponds with it,
since higher outputs were used. The following charges
were, therefore, modified. This modification consisted of
an increase in the Al amount by 1 g for the charge of
master alloy NbAl-3, and by 2 g for the charges of the
TaAl-2 and TaAl-3 master alloys. It follows from Table
1 that in the cases of the master alloys whose Al amounts
in the charges were not increased, the Al amounts
decreased. For the master alloys, whose charges were
thus increased, the decrease of Al was lower. This
confirms the fact that a loss of mass after melting was
caused by the evaporation of the Al part. Nevertheless,
for all the master alloys, with the exception of the TaAl-3
master alloy, the determined amount of Al was lower
than for the nominal composition. The determined values
J. JUØICA et al.: PREPARATION AND PROPERTIES OF MASTER ALLOYS Nb-Al AND Ta-Al FOR MELTING ...
28 Materiali in tehnologije / Materials and technology 49 (2015) 1, 27–30
Figure 1: Binary-phase diagrams for: a) Nb-Al and b) Ta-Al with
indicated chemical compositions3,4
Slika 1: Binarna fazna diagrama: a) Nb-Al in b) Ta-Al z ozna~eno
kemijsko sestavo3,4
Table 1: Al amounts in different parts of the ingots detected with EDS
Tabela 1: Vsebnost Al v razli~nih delih ingota, dolo~ena z EDS
Alloy
The nominal
composition
(x/%)
EDS-measured composition (x/%)
Front part Rear part
NbAl-1 Nb-60Al 56.82 ± 0.58 58.41 ± 0.24
NbAl-2 Nb-60Al 57.93 ± 0.46 59.46 ± 0.12
NbAl-3 Nb-60Al 58.98 ± 0.51 59.73 ± 0.39
TaAl-1 Ta-80Al 76.79 ± 3.00 78.43 ± 3.86
TaAl-2 Ta-80Al 78.87 ± 2.98 79.23 ± 3.02
TaAl-3 Ta-80Al 79.01 ± 4.01 81.72 ± 3.28
of chemical compositions also manifest substantial diffe-
rences between individual parts of an ingot. This can be
explained with a slow shift of the mould and a segrega-
tion of Al. However, a non-homogenous distribution of
the elements in the ingots of the master alloys does not
represent the principal problem, since we always used
whole ingots for the preparation of the TiAlNb and
TiAlTa alloys. The determined amounts of oxygen in the
prepared master alloys were very low, (29 ± 2) μg/g for
the Nb-60Al master alloy and (123 ± 21) μg/g for the
Ta-80Al master alloy.
The microstructures of the Nb-60Al master alloys are
shown in Figures 2 and 3. In principle, these microstruc-
tures correspond to the binary diagram (Figure 1a) con-
sisting of Nb2Al dendrites, presented as light areas, and a
regular eutectic consisting of Nb2Al and NbAl3 phases.
The NbAl3 phase is presented, in the micrographs below,
as darker areas. The average amounts of the elements in
individual phases or areas, calculated from the measured
chemical composition according to EDS, are presented
in Table 2.
The occurrence of the Nb2Al and NbAl3 phases was
also confirmed with the X-ray diffraction pattern shown
in Figure 4.
Figures 5 and 6 show BSE images of the micro-
structure of the TaAl-2 master alloy. The microstructures
of all the prepared master alloys were dendritic, not
showing any significant changes between individual
ingots and parts of the ingots. The established chemical
composition of dendrites and interdendritic areas is
presented in Table 3. According to the binary diagram of
the Ta-Al binary system (Figure 1b), the chemical com-
position of dendrites corresponds to the TaAl3 phase, and
the chemical composition of interdendritic areas corres-
ponds to the solid solution of Ta and Al. The occurrence
of these two phases was also confirmed with the X-ray
diffraction analysis (Figure 7).
J. JUØICA et al.: PREPARATION AND PROPERTIES OF MASTER ALLOYS Nb-Al AND Ta-Al FOR MELTING ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 27–30 29
Figure 3: Microstructure of master alloy NbAl-1
Slika 3: Mikrostruktura predzlitine NbAl-1
Table 2: Al amounts measured in individual structural fields shown in
Figure 3
Tabela 2: Vsebnost Al, izmerjena na posameznih podro~jih strukture,
prikazanih na sliki 3
Al amount by EDS (x/%)
Ingot Nb2Al (1) NbAl3 (2) Eutectic (3)
NbAl-1 46.80 ± 0.33 71.68 ± 1.34 59.66 ± 0.50
NbAl-2 46.48 ± 0.15 76.05 ± 0.41 60.62 ± 0.22
NbAl-3 46.39 ± 0.11 75.58 ± 0.39 60.88 ± 0.34
Figure 2: Microstructure of master alloy NbAl-2
Slika 2: Mikrostruktura predzlitine NbAl-2
Figure 5: Microstructure of master alloy TaAl-2
Slika 5: Mikrostruktura predzlitine TaAl-2
Figure 4: Diffraction pattern of NbAl-2 master alloys: 1-NbAl3, 2-
Nb2Al
Slika 4: Uklonska slika predzlitin NbAl-2: 1-NbAl3, 2- Nb2Al
The results of the differential thermal analysis of the
master alloys are shown in Figure 8. The determined
temperature of the liquidus of the NbAl alloy was 1624
°C and the temperature of the liquidus of the TaAl alloy
was 1541 °C. The temperatures presented in Figure 8
are not corrected and must be corrected to the melting
temperature of pure Pd (5N, a calibration standard). We
obtained the melting temperature of Pd under identical
conditions. After the correction, the determined liquidus
temperature of NbAl was 1615 °C and that of TaAl was
1532 °C.
4 CONCLUSION
Master alloys of NbAl and TaAl with the nominal
compositions of Nb-60Al and Ta-80Al (x/%) were
prepared by melting in a plasma furnace. These master
alloys contain very low amounts of oxygen, namely, (29
± 2) μg/g for the Nb-60Al master alloy and (123 ± 21)
μg/g for the Ta-80Al master alloy. The melting tempe-
ratures of these master alloys were 1615 °C for NbAl
and 1532 °C for TaAl. The DTA analysis, therefore, con-
firmed low melting temperatures of the prepared master
alloys and, thus, their suitability for the preparation of
-TiAl alloys alloyed with Nb or Ta.
Acknowledgement
This paper was created within projects No.
CZ.1.05/2.1.00/01.0040 "Regional Materials Science and
Technology Centre" within the frame of the operational
programme "Research and Development for Innova-
tions" financed by the Structural Funds and the state
budget of the Czech Republic and SP2013/64 "Specific
research in metallurgy, materials and process engi-
neering".
5 REFERENCES
1 V. Imayev, Advanced Engineering Materials, 10 (2008), 1095–1100
2 D. Hu, Intermetallics, 15 (2007), 327–332
3 T. Takeuchi [cit. 2013-08-27], http://www.nims.go.jp/smcMetal/
takeuchi.html
4 ASM International phase diagram database [cit. 2013-08-27]
http://www1.asminternational.org/asmenterprise/apd/
5 K. Frkáòová, Conference proceedings METAL, Brno, 2011
6 F. Gomes, Materials Science Forum, 730–732 (2012), 697–702
J. JUØICA et al.: PREPARATION AND PROPERTIES OF MASTER ALLOYS Nb-Al AND Ta-Al FOR MELTING ...
30 Materiali in tehnologije / Materials and technology 49 (2015) 1, 27–30
Figure 8: DTA curves of NbAl and TaAl master alloys
Slika 8: DTA-krivulji predzlitin NbAl in TaAl
Figure 7: Diffraction pattern of TaAl-2 master alloys: 1-TaAl3, 2-Al
Slika 7: Uklonska slika predzlitin TaAl-2: 1-TaAl3, 2-Al
Figure 6: Microstructure of master alloy TaAl-3
Slika 6: Mikrostruktura predzlitine TaAl-3
Table 3: Al amounts measured in individual structural fields shown in
Figure 6
Tabela 3: Vsebnost Al, izmerjena na posameznih podro~jih strukture,
prikazanih na sliki 6
Al amount by EDS (x/%)
Ingot Dendrites Interdendriticspace
TaAl-1 75.08 ± 0.19 99.73 ± 0.13
TaAl-2 75.18 ± 0.27 99.56 ± 0.35
TaAl-3 75.15 ± 0.21 99.71 ± 0.28
P. MARTÍNEK et al.: DECREASING THE CARBONITRIDE SIZE AND AMOUNT IN AUSTENITIC STEEL ...
DECREASING THE CARBONITRIDE SIZE AND AMOUNT IN
AUSTENITIC STEEL WITH HEAT TREATMENT AND
THERMOMECHANICAL PROCESSING
ZMANJ[ANJE VELIKOSTI KARBONITRIDOV V AVSTENITNEM
JEKLU S TOPLOTNO OBDELAVO IN TERMOMEHANSKO
PREDELAVO
Petr Martínek, Pavel Podaný, Jan Nacházel
COMTES FHT a.s., Dobøany, Czech Republic
pmartinek@comtesfht.cz
Prejem rokopisa – received: 2013-10-01; sprejem za objavo – accepted for publication: 2014-03-07
The article deals with the heat treatment and forging of AISI 321 austenitic stainless steel with a combination of the main
alloying elements, Cr-Ni, in the mass ratio of about 18 : 10 and its variants used in the energy industry. The experiment was
focused on the influence of heat treatment and thermomechanical processing on the microstructure and, especially, on the distri-
bution of titanium carbo/nitrides and their agglomerations. Three experimental heats with various amounts of carbon, titanium
and boron were prepared and subjected to different heat-treatment regimes. Also, different solution-annealing treatments were
applied after the forging. The distribution of titanium carbo/nitrides was analyzed in different areas of the tested ingots. The
microstructure of the samples was analyzed by means of light and scanning electron microscopy. A numerical simulation in the
DEFORM HT software was used for simulating the cooling in different environments after the forging.
Keywords: austenitic steels, heat treatment, titanium carbonitrides
^lanek obravnava toplotno obdelavo in kovanje avstenitnega nerjavnega jekla AISI 321 s kombinacijo glavnih legirnih ele-
mentov Cr-Ni v masnem razmerju 18 : 10 in njegovih razli~ic, ki se uporabljajo v energetiki. Preizkus je bil usmerjen na vpliv
toplotne obdelave in termomehanske predelave na mikrostrukturo in predvsem na razporeditev titanovih karbonitridov in
njihovo grupiranje. Pripravljene so bile tri eksperimentalne taline z razli~no vsebnostjo ogljika, titana in bora ter izpostavljene
razli~nim na~inom toplotne obdelave. Po kovanju so bila izvr{ena razli~na raztopna `arjenja. Na razli~nih podro~jih eksperi-
mentalmih ingotov je bila analizirana razporeditev titanovih karbonitridov. Mikrostruktura vzorcev je bila analizirana s svet-
lobno in vrsti~no elektronsko mikroskopijo. Za simulacijo ohlajanja v razli~nih okoljih po kovanju je bila uporabljena nume-
ri~na simulacija s programsko opremo DEFORM HT.
Klju~ne besede: avstenitna jekla, toplotna obdelava, titanovi karbonitridi
1 INTRODUCTION
Stainless steels with an austenitic microstructure and
an approximate composition of mass fractions w = 18 %
chromium, 10 % nickel and additions of molybdenum,
titanium or niobium are today widely used in the com-
ponents designed for high-temperature applications like
nuclear power stations, boilers or superheaters.1 AISI
321 is a typical austenitic stainless steel with this com-
bination of the main alloying elements of Cr-Ni.
The main goal of this study was to reduce the amount
of large carbo/nitride agglomerations. Large agglomera-
tions usually cause echoes during ultrasonic inspections
and, thus, almost finished products have to be rejected.
In normal practice, a sufficient amount of titanium is
added to steel to combine with all the carbon. Titanium
prevents the formation of Cr23C6 carbides, which locally
deplete the matrix of chromium2; however, a reduction in
the carbon amount below w = 0.03 % does more to
improve the sensitization resistance than an increase in
the Ti amount.3
Titanium and niobium carbides are much less soluble
in austenite than chromium carbide, so they form at
much higher temperatures as relatively stable particles.
These should remain relatively inert during commercial
heat treatments involving solution temperatures no
higher than 1050 °C, thus minimizing the possible
nucleation of Cr23C6. However, TiC and NbC have some
solubility in austenite at 1050 °C and can subsequently
precipitate at lower temperatures. During high-tempe-
rature processes, these carbides dissolve to a greater
extent in austenite and can then reprecipitate at lower
temperatures. Therefore, NbC and TiC do not always
form inert dispersions and are often likely to be redistri-
buted due to heat treatment. They do, however, have a
great advantage of not depleting the matrix of chromium,
particularly in the sensitive areas such as grain bounda-
ries.4
The precipitation of carbides and nitrides also occurs
during the rapid quenching from high solution tempera-
tures;4 nevertheless, if only the final heat treatment is
applied on thermomechanically processed products
(long-time annealing at the temperatures above 1150
°C), grain coarsening can be observed. M23C6 particles
precipitated predominantly at the grain boundaries,
causing an earlier onset of an abnormal grain growth as
compared to the samples, in which precipitation occurred
Materiali in tehnologije / Materials and technology 49 (2015) 1, 31–36 31
UDK 669.15-194.56:539.4.016:621.78 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)31(2015)
in the interior of the grains (at the lowest temperature,
1090 °C). This effect is associated with the stability and
kinetics of the dissolution of M23C6 particles.5
The solution heat treatment in the range of
1100–1250 °C also has a beneficial effect on the creep
resistance due to the subsequent precipitation of MX pre-
cipitates.6
2 EXPERIMENT
Three different heats of austenitic stainless steel were
prepared. The chemical composition meets the require-
ment of the Russian standard 08Ch18N10T and Ameri-
can standards AISI 321 and AISI 321H (in the case of a
heat enrichment with boron). The chemical composition
of the experimental material is summarized in Table 1.
Table 1: Chemical composition of experimental heats
Tabela 1: Kemijska sestava eksperimentalnih talin
Element Heat number
(w/%) 46471 46736 46777
C 0.0400 0.0400 0.0600
Mn 1.5100 1.6100 1.7000
Si 0.5300 0.5100 0.5700
P 0.0250 0.0220 0.0210
S 0.0070 0.0020 0.0020
Cr 17.600 17.650 17.600
Ni 10.300 10.100 10.100
Ti 0.2200 0.2800 0.3900
B 0.0001 0.0037 0.0045
N 0.0215 0.0265 0.0116
The main variation of the chemical composition
consists of different amounts of titanium, carbon and
boron. The experiment was mainly focused on the
microstructure examination and image analysis. This was
done on the ingots. The ingot size and shape were 8K 1.1
with a mass of 1000 kg (the mean diameter was about
500 mm, the length without the head was 1120 mm). The
microstructure was inspected in three localities (the ingot
axis – 1; the half of radius – 2; the edge – 3), of the
ingots after different heat treatments, on the samples
forged from the ingots and cooled in two different envi-
ronments and on these samples after the high-tempera-
ture annealing. The microstructure examination and
image analyses were focused on the amount and distri-
bution of titanium carbonitrides, on the volume fraction
of delta ferrite and the grain size after forging and heat
treatment. The particular parameters of the experiment
described above are shown in Table 2. All the samples
for the microstructure documentation were cut in a cubic
shape with the dimensions of about 12 mm × 12 mm ×
12 mm.
Table 2: Parameters of the experiment (for the samples from ingots
and the samples forged from ingots)
Tabela 2: Parametri preizkusa (na vzorcih iz ingotov in na kovanih
vzorcih iz ingotov)
Experimental
procedure
Experiments on
ingots
Experiments on forged
samples
Heat treatment of
the samples from
ingots (solution
annealing)
Cooling after
forging
(forging at
1200 °C)
Annealing
temperature
(°C)
Parameters
Temp.
(°C)
Time
(h) Vermiculite
or
water
quenching
1020 or
1100
900 1; 5; 10
1100 1; 5; 10
1300 1; 5; 10
Analyses Microstructure examination and image analysis
The ingots were cut with a conventional cutting saw.
The samples for the microstructure examination were
subjected to a conventional metallographic procedure
(grinding and subsequent polishing with an alumina
suspension). The Beraha 2 etchant was used for reveal-
ing delta ferrite. Austenite grains were revealed by
means of the V2A etchant.
The identification of carbonitrides was done by
means of an HKL Nordlyss EBSD camera and EDS
detector INCAx-sight on a JEOL 7400F scanning elec-
tron microscope (Figure 1).
The image analysis was made with a Nikon MA 200
light microscope with the NIS Elements software for
image analyses.
All the samples were evaluated at a 500× magnifica-
tion and inspected in 30 image fields of each sample
(one image field represents an area of 0.022728 mm2).
The entire experimental heat treatment was done with
an annealing furnace Heraus. The forging of the samples
from the experimental ingots was done with a hydraulic
press Zeulenroda (PYE 40).
P. MARTÍNEK et al.: DECREASING THE CARBONITRIDE SIZE AND AMOUNT IN AUSTENITIC STEEL ...
32 Materiali in tehnologije / Materials and technology 49 (2015) 1, 31–36
Figure 1: EBSD and EDX analysis of TiCN
Slika 1: EBSD- in EDX-analiza TiCN
3 RESULTS AND DISCUSSION
3.1 Microstructure of the ingots
A quantitative image analysis of the carbonitride
count and the carbonitride size found that there is no
essential difference between particular heats. However, it
was found that there are substantial differences between
the localities within an ingot. The highest amount of car-
bonitrides was concentrated on the edge (near the sur-
face) of an ingot. On the other hand, the size (the average
area) of these particles near the surface is the lowest (see
Figure 2 where locality 1 is the ingot axis, locality 2 is at
the half of its radius and locality 3 is on the edge of the
ingot). The term "object count" in the figure represents
the amount of particles per analysed image field, while
"object area" represents the average area of one particle.
Agglomerations of carbonitrides are classified as one
large carbonitride during the image analyses. Histograms
show the distribution of carbonitrides into 10 size factors
on the x-axis (from 0 μm2 to 10 μm2 with a 1 μm2 step).
Class "other" indicates carbonitride agglomerations. It is
clearly visible that the amount (the y-axis) of large car-
bonitride agglomerations decreases from the ingot axis
to its surface (Figure 3).
Delta ferrite formed a complex network around the
as-cast grains in the middle of the ingot (Figures 4 and
P. MARTÍNEK et al.: DECREASING THE CARBONITRIDE SIZE AND AMOUNT IN AUSTENITIC STEEL ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 31–36 33
Figure 2: Amount of carbonitride particles and the average area of
carbonitride particles in ingots: a) amount of particles per field, b) ave-
rage area of particles
Slika 2: Dele` delcev karbonitridov in povpre~na velikost karbonitrid-
nih delcev v ingotih: a) koli~ina delcev v enem polju, b) povpre~na
velikost podro~ja delcev
Figure 5: Size and distribution of delta ferrite in ingot 46471 – the
edge of the ingot, Beraha 2 etchant
Slika 5: Velikost in razporeditev delta ferita v ingotu 46471 – rob
ingota, jedkalo Beraha 2
Figure 4: Size and distribution of delta ferrite in ingot 46471 – the
middle part of the ingot, Beraha 2 etchant
Slika 4: Velikost in razporeditev delta ferita v ingotu 46471 – sredina
ingota, jedkalo Beraha 2
Figure 3: Histograms of carbonitride classification for ingot 46471: a)
ingot axis, b) half of radius, c) edge of ingot
Slika 3: Histogrami razporeditve velikosti karbonitridov v ingotu
46471: a) os ingota, b) polovica polmera ingota, c) rob ingota
5). Near the edge of the ingot, delta ferrite forms more
separate islands and its volume fraction is lower. Both
delta-ferrite and carbonitride distribution within the ingot
are related to the ingot solidification. Carbonitrides form
agglomerations due to the reduced speed of solidification
and delta ferrite is formed due to the segregation of the
alpha-phase-forming elements.
3.2 Heat treatment of ingots
The samples from the ingots were annealed at (900,
1100 and 1300) °C for (1, 5 or 10) h. The results of the
image analyses after the solution annealing are shown
Table 3.
The annealing at 900 °C for a short time (1 h and 5 h)
does not substantially affect the size and count of car-
bonitrides. The effect of annealing on this temperature is
only visible after the annealing for 10 h when the amount
of carbonitrides decreases and its average area slightly
increases.
An increase in the solution temperature to 1100 °C
led to a similar effect on the particle size even after 1 h.
But the average amount of carbonitrides per inspected
field did not change substantially. A longer annealing at
this temperature did not bring any other effect.
The best result was obtained with the annealing at
1300 °C for 1 h and 5 h (Figure 6). The average area of
one particle decreased and the average count of particles
increased. However, the annealing prolonged to 10 h
gave almost the same values of the particle count and the
average area was the same as in the initial state before
the annealing. Nevertheless, the annealing at this tempe-
rature for 10 h led to a more homogeneous distribution
of carbonitrides.
The best result was reached after the annealing at
1300 °C for 1 h and 5 h when the average area of parti-
cles decreased and the average amount of particles per
inspected image field increased.
3.3 Forging of the samples from the ingots and diffe-
rent cooling processes
The samples made from the edges of the ingots were
heated up to 1200 °C and solution annealed for 2 h and
after that they were forged, with a degree of reduction of
2.5. Two different cooling processes were applied:
quenching in water and simulation of cooling a whole
forged bar with a 250 mm diameter in the air. An
insertion of a forged sample into insulating vermiculite
simulated this process. The condition of cooling was
verified by means of the DEFORM HT software, which
proved a similarity of the cooling conditions for a small
sample in vermiculite. The influence of forging and the
way of cooling on the carbonitride size and distribution
were then studied. The forging itself has a substantial
P. MARTÍNEK et al.: DECREASING THE CARBONITRIDE SIZE AND AMOUNT IN AUSTENITIC STEEL ...
34 Materiali in tehnologije / Materials and technology 49 (2015) 1, 31–36
Figure 6: Size and distribution of carbonitrides in ingot 46471: a) ini-
tial state, as polished, b) after the heat treatment (1300 °C, 5 hours), as
polished
Slika 6: Velikost in razporeditev karbonitridov v ingotu 46471: a) za-
~etno stanje, polirano, b) po toplotni obdelavi (1300 °C, 5 h), polirano
Table 3: Results of the image analyses after the solution annealing (object count represents the amount of particles per analysed image field,
object area represents the average area of one particle)
Tabela 3: Rezultati analize slik po raztopnem `arjenju ("{tevilo objektov" pomeni {tevilo delcev na enem analiziranem polju, velikost podro~ja
pomeni povpre~no velikost enega delca)
TiCN
46471 46736 46777
Obj. count Obj. area Obj. count Obj. area Obj. count Obj. area
Temperature
Time (–) (ìm2) (–) (ìm2) (–) (ìm2)
0 13.00 3.49 14.1 3.67 14.70 3.31
900 °C
1 12.10 3.48 12.2 3.21 14.17 3.05
5 12.87 3.37 13.67 3.04 14.47 3.33
10 10.57 5.40 6.8 5.90 14.53 5.33
1100 °C
1 16.57 4.15 8.93 5.77 12.80 4.88
5 16.33 4.60 10.53 4.68 18.33 5.67
10 12.37 4.06 7.67 6.64 14.77 5.32
1300 °C
1 36.13 1.72 30.13 2.10 25.43 2.05
5 37.03 1.98 28.23 2.31 26.27 1.84
10 20.53 2.50 12.77 3.44 15.20 2.60
influence on the size and distribution of carbonitrides.
An increase in the amount of carbonitrides was observed
for all three heats. Heat 46777 showed the strongest in-
crease in carbonitrides due to forging. The observed
particles were unfortunately too small for an EDS anal-
ysis. But with respect to their positioning within the
grain interior, the carbides are not detrimental. Only a
low increase in the particle count was observed after the
cooling in vermiculite in comparison to the cooling in
water. The forging of all the samples led to a decrease in
the size of carbonitrides.
3.4 Annealing of forged samples
The forged samples were afterwards annealed at two
different temperatures, 1020 °C and 1100 °C, to achieve
a further dissolution of carbonitrides. It was found that in
the case of heats 46736 and 46777 annealing the forgings
quenched in water at 1020 °C led to a further increase in
the particle count (Figure 7) and a decrease in the ave-
rage area of particles (Table 4).
Increasing the annealing temperature for these heats
was beneficial for a further particle refinement but ne-
gligible for the final grain size. The grain size increased
from 8 (after the forging) to 5 (after the annealing at
1100 °C) according to ASTM E 112 (Figure 8). This
result was the same for all the heats, there were only
minor differences between them.
3.5 Mechanical testing
There was not enough material for standard mecha-
nical testing, thus, mini-tensile test samples were made
P. MARTÍNEK et al.: DECREASING THE CARBONITRIDE SIZE AND AMOUNT IN AUSTENITIC STEEL ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 31–36 35
Table 4: Average values of the particle count and size for the initial state, after the forging (1 – quenching in water, 2 – cooling in vermiculite)
and annealing
Tabela 4: Povpre~ne vrednosti {tetja delcev in velikost v za~etnem stanju, po kovanju (1 – ohlajanje v vodi, 2 – ohlajanje v vermikulitu) in po
`arjenju
TiCN
46471 46736 46777
Obj. count Obj. area Obj. count Obj. area Obj. count Obj. area
state – ìm2 – ìm2 – ìm2
initial state 13.0 3.49 14.1 3.67 14.70 3.31
forging 1 27.97 2.17 18.13 2.86 166.53 0.92
1020 °C 23.9 2.63 24.87 2.83 192.57 0.90
1100 °C 23.73 2.95 45.13 2.10 203.83 0.89
forging 2 25.7 3.05 46.20 2.70 171.73 0.95
1020 °C 23.03 2.81 32.73 1.98 180.37 0.97
1100 °C 19.43 3.05 57.23 1.56 215.57 0.90
Figure 8: Microstructure of a forged sample after forging, cooling in
vermiculite and annealing: a) heat 46736 annealed at 1020 °C, V2A
etchant, b) heat 46736 annealed at 1100 °C, V2A etchant
Slika 8: Mikrostruktura kovanega vzorca po kovanju, ohlajanju v
vermikulitu in `arjenju: a) talina 46736, `arjena pri 1020 °C, jedkalo
V2A, b) talina 46736, `arjena pri 1100 °C, jedkalo V2A
Figure 7: Microstructure of a forged sample after forging, cooling in
vermiculite and annealing: a) heat 46736 annealed at 1020 °C, po-
lished, b) heat 46777 annealed at 1020 °C, polished
Slika 7: Mikrostruktura kovanega vzorca po kovanju, ohlajanju v ver-
mikulitu in `arjenju: a) talina 46736, `arjena pri 1020 °C, polirano, b)
talina 46777, `arjena pri 1020 °C, polirano
from the forged specimens. A sample for the mini-tensile
testing is flat, with very small dimensions (the functio-
nal-body length is in units of millimetres). The tests were
done on a mini-tensile test machine MTS with a video-
extensometer MESSPHYSIK. The results of testing the
samples annealed at 1020 °C are in Table 5. The yield
strength increases due to the increasing count of carbo-
nitride particles (precipitation strengthening). Thus, heat
46777 shows the highest YS and TS of all the heats. The
mechanical test of the samples annealed at 1100 °C was
not performed because of insufficient grain size.
Table 5: Results of mechanical testing after forging 1 (subsequent
quenching in water) and annealing at 1020 °C
Tabela 5: Rezultati mehanskih preizkusov po kovanju 1 (sledilo je
ohlajanje v vodi) in `arjenju pri 1020 °C
Specimen Rp0.2/MPa Rm/MPa A/%
46471 295.5 609.3 70.5
46736 323.4 595.4 56.7
46777 343.2 628.7 63.0
4 CONCLUSIONS
Complex experiments were performed on austenitic
stainless steel. Three experimental heats of AISI 321(H)
with various amounts of carbon, titanium and boron were
subjected to a particular annealing heat treatment, forg-
ing with various cooling procedures and a further anneal-
ing. Three different temperatures of solution annealing
were applied on the ingots – (900, 1100 and 1300) °C,
with different holding times – (1, 5 and 10) h. The forged
samples were cooled in two different environments –
rapid quenching in water and slow cooling in vermicu-
lite. The last heat treatment – solution annealing – was
applied on the forged samples – at 1020 °C and 1100 °C
for 1 h.
The highest amount of carbonitrides is concentrated
on the edge (near the surface) of an ingot where rapid
solidification takes place. Carbonitrides become
coarsened, forming agglomerations towards the ingot
axis, together with the decreasing solidification velocity.
Delta ferrite forms a network structure near the ingot
axis; it divides into separate islands near the ingot sur-
face.
A substantial decrease in the carbonitride amount
was reached after the heat treatment of the ingots at the
temperature 1300 °C for 1 h or 5 h when the average area
of particles decreased and the average amount of
particles per inspected image field increased. Increasing
both the annealing temperature and the holding time led
to a substantial decrease in the delta-ferrite volume
fraction. No substantial differences between the amounts
of the observed carbonitrides, in dependence of particu-
lar carbon and titanium amounts in the ingots were ob-
served.
Further experiments were concentrated on the sam-
ples taken from the edges of the ingots. Thermomecha-
nical processing – the forging itself – substantially
decreased the size of particles in all the heats. On the
other hand, the amount of carbonitrides increased. The
minimum increase was observed on heat 46471 with the
lowest carbon and titanium amounts. The highest
increase in the carbonitride amount was observed on heat
4677 with the highest titanium, carbon and boron
amounts. Considering the difference between the cooling
procedures in two environments, a substantial increase in
the particle count was observed for heat 46736 after the
cooling in vermiculite in comparison to the cooling in
water. However, the differences for the other heats were
not crucial.
The highest increase in carbonitrides after the solu-
tion annealing at 1020 °C was observed in heat 46777
(the heat with the highest carbon and boron amounts).
Using a higher annealing temperature (1100 °C) was not
beneficial because of the coarsening of the grain size.
The grain size increased from 8 (after the forging) to 5
(after the annealing at 1100 °C) according to ASTM E
112. This correlates with the other studies5. The heats
with higher titanium and boron amounts (46736 and
46777) exhibited good mechanical properties. The
precipitation of very fine carbonitrides has a positive
effect on the yield strength, as previously stated by
Sourmail6.
Acknowledgments
The results presented in this paper were supported
by: project FR-TI1/222 of the Ministry of Industry and
Trade of the Czech Republic and project West-Bohemian
Centre of Materials and Metallurgy CZ.1.05/2.1.00/
03.0077, co-funded by the European Regional Develop-
ment Fund.
5 REFERENCES
1 G. Kaishu et al., Effect of aging at 700 °C on precipitation and
toughness of AISI 321 and AISI 347 austenitic stainless steel welds,
Nuclear Engineering and Design, 235 (2005), 2485–2494
2 K. Kaneko et al., Formation of M23C6-type precipitates and chro-
mium-depleted zones in austenite stainless steel, Scripta Materialia,
65 (2011) 6, 509–512
3 A. Pardo et al., Influence of Ti, C and N concentration on the
intergranular corrosion behaviour of AISI 316Ti and 321 stainless
steels, Acta Materialia, 55 (2007), 2239–2251
4 W. K. Honeycombe, H. K. D. H. Bhadeshia, Steels: Microstructure
and Properties, 3rd Ed., Elsevier Ltd., 2006, 267
5 J. C. Dutra et al., Interaction Between Second-Phase Particle Disso-
lution and Abnormal Grain Growth in an Austenitic Stainless Steel,
Materials Research, 5 (2002) 3, 379–384
6 T. Sourmail, Precipitation in creep resistant austenitic stainless steels,
Materials Science and Technology, 17 (2001) 1, 1–14
P. MARTÍNEK et al.: DECREASING THE CARBONITRIDE SIZE AND AMOUNT IN AUSTENITIC STEEL ...
36 Materiali in tehnologije / Materials and technology 49 (2015) 1, 31–36
P. SUCHMANN et al.: DEEP CRYOGENIC TREATMENT OF H11 HOT-WORKING TOOL STEEL
DEEP CRYOGENIC TREATMENT OF H11 HOT-WORKING TOOL
STEEL
GLOBOKA KRIOGENSKA OBDELAVA ORODNEGA JEKLA H11
ZA DELO V VRO^EM
Pavel Suchmann1, Dagmar Jandova2, Jana Niznanska1
1COMTES FHT a. s., Prùmyslová 995, 334 41 Dobrany, Czech Republic
2VZU PLZEN s. r. o., Tylova 46, 323 00 Plzen, Czech Republic
pavel.suchmann@comtesfht.cz
Prejem rokopisa – received: 2013-10-02; sprejem za objavo – accepted for publication: 2014-02-07
Unlike a conventional cold treatment, which is commonly used for the elimination of retained austenite, a deep cryogenic
treatment (DCT) primarily improves the wear resistance of tools. This effect is supposed to result from the preferential
precipitation of fine 
-carbides, whose formation mechanism has been the subject of several recent investigations, performed
mainly on high-speed steels. This article describes the influence of a DCT on the microstructure and properties of
X37CrMoV5-1 (H11) hot-working tool steel. The wear resistance of specimens treated using DCT was analysed using a
pin-on-disc wear tester and compared to that of specimens treated using standard quenching and tempering. In addition, the
specimens’ microstructures were analysed by TEM. The results show a significant improvement in the wear resistance as a
consequence of the DCT, especially at the high sliding velocities that are typical of many industrial applications for hot-working
steels (e.g., closed die forging). Apart from this, some effects of DCT on the microstructure were found, which contributed to a
better understanding of this process.
Keywords: tool steel, wear resistance, microstructure
Razli~no od konvencionalnega podhlajevanja, ki se navadno uporablja za odpravo zaostalega avstenita, se kriogena obdelava
(DCT) uporablja predvsem za pove~anje obrabne odpornosti jekla. Domneva se, da to izvira iz preferen~nega izlo~anja drobnih

-karbidov. Mehanizem njihovega nastanka je predmet {tevilnih raziskav, narejenih predvsem pri hitroreznih jeklih. ^lanek
opisuje vpliv DCT na mikrostrukturo in lastnosti X37CrMoV5-1 (H11) orodnega jekla za delo v vro~em. Obrabna odpornost
vzorcev, obdelanih z DCT, je bila dolo~ena na napravi za dolo~anje obrabe "pin-on-disc" in primerjana z obrabno odpornostjo
vzorcev, obdelanih z navadnim kaljenjem in popu{~anjem. Mikrostruktura vzorcev je bila analizirana s TEM. Rezultati ka`ejo
ob~utno izbolj{anje odpornosti proti obrabi pri DCT-obdelavi, posebno pri velikih hitrostih drsenja, ki so zna~ilne pri mnogih
vrstah industrijske uporabe jekel za delo v vro~em (npr. kovanje v utopih). Poleg tega so bili ugotovljeni vplivi DCT na
mikrostrukturo, kar prispeva k bolj{emu razumevanju tega procesa.
Klju~ne besede: orodno jeklo, odpornost proti obrabi, mikrostruktura
1 INTRODUCTION
The deep cryogenic treatment of steels has been the
subject of numerous research and experimental studies
over the past twenty years. Its practical significance has
been growing with new findings, which suggest that the
process may substantially extend the life of various types
of tools. From the technical viewpoint, a deep cryogenic
treatment is a single processing operation that immedi-
ately follows the conventional quenching (and precedes
tempering). One should, however, distinguish between a
deep cryogenic treatment and a conventional cold treat-
ment. The latter has been commonly used in industry
since the 1950s. The purpose of a cold treatment (same
as in multiple tempering) is to eliminate the retained
austenite from the microstructure of the hardened steel.
In high-alloyed tool steels the retained austenite remains
stable down to the approximate temperature range
–80 °C to –120 °C. In a cold treatment, a temperature
below this level is reached only once in the process. The
treatment improves the material’s mechanical properties
(i.e., the strength) and enhances the dimensional stability
for tools processed in this manner.
In contrast, in a deep cryogenic treatment the elimi-
nation of retained austenite is no more than a side effect.
A number of reports suggest that in steels the super-
cooling below –150 °C and a subsequent long holding
time (of the order of hours or tens of hours) at the low
temperature leads to a substantial improvement in their
resistance wear in service. This applies even in compa-
rison with materials subjected to a conventional cold
treatment (i.e., materials with a zero retained austenite
content).
The improvement in tool life due to a deep cryogenic
treatment has recently been reported in various types of
forming and cutting tools.1,2 In forging dies from the
X37CrMoV5-1 (H11) steel, an improvement in service
life by 40 % was proven repeatedly,2 when compared to
conventional quenching and tempering without a deep
cryogenic treatment. An extensive series of experiments
on various types of punching tools, milling cutters and
drill bits from several types of tool steels demonstrated a
definite contribution of the deep cryogenic treatment to a
longer tool life.3 Another widely published series of
experiments was conducted on turning tools made from
HS10-4-3-10 high-speed steel.4 In this case, too, the
Materiali in tehnologije / Materials and technology 49 (2015) 1, 37–42 37
UDK 621.78.08:539.538 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)37(2015)
improvement in the cutting edge’s life thanks to a deep
cryogenic treatment was demonstrated.
The microstructural aspects of these effects were
explored in the 1990s by Collins and other researchers.5–8
In the case of the X155CrVMo12-1 steel (the nearest
equivalent grade: D2), a large content of fine secondary
carbides was found in the deep cryogenically treated
material, in addition to its improved wear resistance over
that of the conventionally quenched and tempered
material. The changes in microstructure taking place in
the course of the deep cryogenic treatment are
summarised and described by Collins as a cold treatment
modification of martensite. The author believes that the
modification consists of the formation of a large number
of lattice defects that serve as nuclei for the precipitation
of these fine carbides during tempering. It should be
noted that all these conclusions have been drawn on the
basis of nothing more than a quantitative image analysis
of optical micrographs. From today’s perspective, this is
an insufficiently detailed analysis of the microstructure.
Another in-depth investigation of the behaviour of
X155CrVMo12-1 steel (the nearest equivalent grade:
D2) after a deep cryogenic treatment was reported by
Das et al.9–11 Tests of the wear resistance revealed its
improvement by 12–39 % after cold treatment and by
34–88 % after a deep cryogenic treatment, when com-
pared to conventional quenching and tempering. Das
reported that a deep cryogenic treatment leads to a per-
manent change in the carbide precipitation kinetics.
According to his report, the material upon deep cryo-
genic treatment contains 22 % more carbides per unit
volume than the conventionally quenched and tempered
material.
The most recent findings on deep cryogenic treat-
ment were reported in 2011 by Oppenkowski,1 who
carried out an in-depth analysis of several types of
high-speed steels that were processed using various tech-
nologies, including cold and deep cryogenic treatments.
He proved that steels, after various hardening proce-
dures, contain different types of martensite. According to
his findings, deep cryogenically treated materials are
characterized by a higher content of martensite with less
tetragonal distortion and finer twin structures. These
microstructural changes are probably the cause of the
precipitation of a large amount of fine carbides during
the subsequent tempering.
The present paper describes the investigation proce-
dure and the results of exploring the impact of a deep
cryogenic treatment on the wear resistance and micro-
structure of the X37CrMoV5-1 (H11) hot-work steel.
The wear resistance was measured at 400 °C using the
pin-on-disc method with a rotary tribometer. The micro-
structure of the steel was examined using light and
transmission electron microscopes.
2 EXPERIMENTAL
2.1 Heat Treating
The experimental specimens of the X37CrMoV5-1
(H11) tool steel had a diameter of 55 mm and a height of
10 mm. These specimens were treated using the sche-
dules listed in Table 1. Two specimens were prepared
with each schedule. One was used for wear-resistance
testing and the other for microstructural observations.
The hardness of all the heat-treated specimens was
52 HRC.
Table 1: Experimental heat-treatment schedules
Tabela 1: Pregled eksperimentalnih toplotnih obdelav
Specimen
Designation Heat Treating Schedule
1 (H+2T) Heating to 1030 °C, quenching in oil,
double tempering (610 °C)
2 (H+6C+2T) Heating to 1030 °C, quenching in oil, deep
freezing at –160 °C for 6 h, double
tempering (610 °C)
3 (H+12C+2T) Heating to 1030 °C, quenching in oil, deep
freezing at –160 °C for 12 h, double
tempering (610 °C)
4 (H+20C+2T) Heating to 1030 °C, quenching in oil, deep
freezing at –160 °C for 20 h, double
tempering (610 °C)
2.2 Wear Resistance
The wear-resistance tests were performed using the
pin-on-disc method. The principle of this test is forcing a
ceramic ball into the surface of a rotating flat specimen.
The holder with the ball indenter is pressed using a
defined force (exerted by a weight) against the specimen.
The holder is attached to a flexible arm, through which
the dependence of the friction coefficient on the length
of the ball’s path is recorded by means of strain gauges.
The testing equipment is schematically shown in Figure
1.
After the test, the profile of the resulting wear track is
measured using a contact profilometer. The wear of the
P. SUCHMANN et al.: DEEP CRYOGENIC TREATMENT OF H11 HOT-WORKING TOOL STEEL
38 Materiali in tehnologije / Materials and technology 49 (2015) 1, 37–42
Figure 1: Pin-on-disc equipment: 1-test sample, 2-indenter holder
with ball indenter, 3-weight, 4-flexible arm
Slika 1: Naprava "pin-on-disc": 1-vzorec, 2-nosilec kroglastega utis-
kovalnika, 3-ute`, 4-gibljiva ro~ica
test specimen is calculated from the measured data as
follows:
W/(μm3/N m) =
=
Wear track volume ( m )
Load (N) Path travelled by
3μ
⋅ ball indenter (m)
The specimens of X37CrMoV5-1 (H11) were tested
with the following parameters:
• Indenter holder with a ball indenter (Si3N4) with the
diameter of D = 6 mm
• Wear-track diameter: r = 3.00 mm
• Temperature: 400 °C
• Load: 10 N
• 5000 cycles (specimen revolutions)
The choice of a temperature of 400 °C is based on the
conditions for the most frequent industrial applications
of the steel in question: forging dies, injection moulds
and other tools whose surfaces may experience a high
temperature in service. For example, during closed-die
forging, the temperatures of the die surface may reach
approximately 250–500 °C.
The wear rates of the specimens tested are given in
Figure 2. The results of the measurement suggest that a
deep cryogenic treatment dramatically improves the
material’s wear resistance. However, with the holding
time at the deep cryogenic temperature becoming longer,
the resistance declines again. This occurrence was men-
tioned in several reports (e.g.,1) but no satisfactory
theoretical explanation has been found.
2.3 Tensile tests
Standard tensile tests of specimens treated according
to Table 1 were performed at room temperature and at
400 °C. For each testing temperature, three specimens
were tested. As the following results (Figures 3 and 4)
show, no significant influence of the deep cryogenic
treatment on the tensile strength or the elongation was
found.
2.4 Impact tests
Standard specimens (10 mm × 10 mm × 55 mm, with
a U-formed notch, the radius of the notch is 1 mm, the
depth of the notch is 5 mm) treated according to the
above-mentioned regimes undertook the Charpy impact
tests at 20 °C and 400 °C. Similar to the results of the
tensile tests, no significant influence of the deep cryo-
genic treatment on the notch toughness of the investi-
gated material was observed (Figure 5).
2.5 Light Microscopy Observation
The specimens were prepared using a standard metal-
lographic procedure: grinding and subsequent polishing.
P. SUCHMANN et al.: DEEP CRYOGENIC TREATMENT OF H11 HOT-WORKING TOOL STEEL
Materiali in tehnologije / Materials and technology 49 (2015) 1, 37–42 39
Figure 4: Results of tensile tests
Slika 4: Rezultati nateznih preizkusov
Figure 2: Wear rates of pin-on-disc test specimens
Slika 2: Hitrost obrabe na preizku{ancih "pin-on-disc"
Figure 5: Results of impact tests
Slika 5: Rezultati udarnih preizkusov
Figure 3: Results of tensile tests
Slika 3: Rezultati nateznih preizkusov
Their microstructures were revealed by etching with nital
3 % and photographed using a NIKON EPIPHOT 200
light microscope. As the micrographs below show, the
microstructures in all the specimens examined are very
similar. They consist in all cases of martensite laths,
within which some fine carbides can be found. Neither
non-metallic inclusions nor other microstructure defi-
ciencies were revealed. Micrographs of the individual
specimens are shown in Figures 6 to 9.
2.6 Transmission Electron Microscopy Observation
In order to document the impact of a deep cryogenic
treatment on the martensite morphology more accurately,
the tempering at 610 °C was omitted in the case of speci-
mens for transmission electron microscopy observation.
The reason was that some recent reports (e.g.,1) claim
that at higher tempering temperatures, the precipitation
of fine carbides in tool steels is accompanied by a gra-
dual annihilation of some types of lattice imperfections
formed during the deep cryogenic treatment. For this
reason, only low-temperature tempering at 180 °C was
applied. The heat-treatment schedules for the specimens
investigated by means of transmission electron micro-
scopy are detailed in Table 2.
Table 2: Heat treatment of specimens for transmission electron micro-
scopy observation
Tabela 2: Toplotna obdelava vzorcev za presevno elektronsko mikro-
skopijo
Specimen Code Heat Treating Schedule
1T (H+2T) Heating to 1030 °C, quenching in oil,tempering (180 °C)
2T (H+6C+2T)
Heating to 1030 °C, quenching in oil,
deep freezing at –160 °C for 6 h,
tempering (180 °C)
The microstructures of specimens 1T and 2T consist
of tempered martensite. No significant differences bet-
ween these specimens were observed in the light micro-
scope (Figure 10 – etched with Vilella-Bain reagent).
The bands occur in both samples as a result of the segre-
gation of alloying elements – this is a typical pheno-
menon in the forgings of hot-working tool steels. A
Widmannstätten structure was evident in the more inten-
sively etched regions (bands), whereas a rather feature-
less structure with a high density of precipitates was
found in adjacent regions.
P. SUCHMANN et al.: DEEP CRYOGENIC TREATMENT OF H11 HOT-WORKING TOOL STEEL
40 Materiali in tehnologije / Materials and technology 49 (2015) 1, 37–42
Figure 8: Microstructure of specimen 3 (H+12C+2T); etched in nital
3 %
Slika 8: Mikrostruktura vzorca 3 (H+12C+2T); jedkano v nitalu 3 %
Figure 9: Microstructure of specimen 4 (H+20C+2T); etched in nital
3 %
Slika 9: Mikrostruktura vzorca 4 (H+20C+2T); jedkano v nitalu 3 %
Figure 7: Microstructure of specimen 2 (H+6C+2T); etched in nital 3
%
Slika 7: Mikrostruktura vzorca 2 (H+6C+2T); jedkano v nitalu 3 %
Figure 6: Microstructure of specimen 1 (H+2T); etched in nital 3 %
Slika 6: Mikrostruktura vzorca 1 (H+2T); jedkano v nitalu 3 %
Thin foils for the transmission electron microscopy
observations were prepared from specimens 1T and 2T
using jet electrolytic polishing in a 6 % solution of per-
chloric acid in methanol at a temperature of approxima-
tely –50 °C. Thin areas, transparent to electrons, accele-
rated by a voltage of 120 kV were found predominantly
in the above-mentioned bands of the featureless struc-
ture. Two different substructures were observed. Tem-
pered martensite of plate-like, rather than lath-like,
character with very thin twins (a width of several tens of
nanometres) was observed in foils from both specimens
(Figure 11). In the 2T specimen, very fine precipitates at
boundaries of the ferritic laths and twins were observed.
A crystallographic investigation has not been performed
yet, as the electron-diffraction patterns obtained are still
not finished and ready for a reliable phase identification.
The second type of substructure revealed a specific
striped contrast and diffraction patterns with streaks typi-
cal for structures after spinodal decomposition (Figure
12). In these areas, some globular and relatively coarse
particles were observed – probably those that were also
observed using the light microscope.
It seems that both phase transformations occur in the
steel during a deep cryogenic treatment: the displacive
martensitic transformation and the spinodal decomposi-
tion. It can be supposed that the spinodal decomposition
occurred predominately in areas enriched with carbon
and other alloying elements. Furthermore, the above-des-
cribed substructures seem to promote the nucleation of
very fine carbides during the annealing of the investi-
gated steel.
3 CONCLUSIONS
Tests carried out in this study have shown that a deep
cryogenic treatment of the X37CrMoV5-1 (H11) steel
has no influence on its tensile strength and notch tough-
ness, but it dramatically improves its resistance to wear
at high temperatures (as determined by the pin-on-disc
test). However, the length of the holding time at the tem-
perature of the deep cryogenic treatment is crucial. The
optimum time appears to be 6 hours. A microstructural
observation in a light microscope did not reveal any
substantial differences between the specimens hardened
in the conventional manner and the specimens after a
deep cryogenic treatment. On the other hand, the analy-
sis of specimens upon deep cryogenic treatment by
means of transmission electron microscopy found two
types of substructure, which are likely to facilitate the
precipitation of fine carbides during the final tempering
of the steel. At this point, the impact of longer holding
times at the deep cryogenic temperature on this substruc-
ture has not been analysed. The cause of the decline in
the wear resistance with extended holding times cannot
thus be determined.
P. SUCHMANN et al.: DEEP CRYOGENIC TREATMENT OF H11 HOT-WORKING TOOL STEEL
Materiali in tehnologije / Materials and technology 49 (2015) 1, 37–42 41
Figure 12: TEM image of specimen 2T
Slika 12: TEM-posnetek vzorca 2T
Figure 11: TEM image of specimen 1T
Slika 11: TEM-posnetek vzorca 1T
Figure 10: Microstructure of 2T specimen; etched in Vilella-Bain
Slika 10: Mikrostruktura vzorca 2T; jedkano z jedkalom Vilella-Bain
4 REFERENCES
1 A. Oppenkowski, Cryobehandlung von Werkzeugstahl, PhD thesis,
Ruhr-Universität Bochum, 2011
2 P. [uchmann, A. Ciski, Enhancing of lifetime of forging dies by
means of deep cryogenic treatment, 19th International Forging Con-
gress, Chicago, USA, 2008
3 F. P. Stratton, Optimizing nano-carbide precipitation in tool steels,
Materials Science and Engineering, A449 (2007), 809–812
4 G. R. Tated, R. S. Kajale, K. Iyer, Improvement in tool life of cutting
tools by application of deep cryogenic treatment, 7th International
Tooling Conference, Torino, Italy, 2006
5 N. D. Collins, G. O’Rourke, Response of Tool Steels to Deep Cryo-
genic Treatment, Effect of Alloying Elements, ASM International,
Materials Park, OH 1999
6 N. D. Collins, Deep Cryogenic Treatment of Tool Steels: a Review,
Heat Treatment of Metals, (1996), 40–42
7 N. D. Collins, J. Dormer, Deep Cryogenic Treatment of a D2 Cold-
Work Tool Steel, Heat Treatment of Metals, (1997), 71–74
8 K. Moore, N. D. Collins, Cryogenic Treatment of Three Heat-
Treated Tool Steels, Key Engineering Materials, 86-87 (1993), 47–54
9 D. Das, K. A. Dutta, V. Toppo, K. K. Ray, Effect of deep cryogenic
treatment on the carbide precipitation and tribological behavior of
D2 steel, Materials and Manufactirung Processes, 22 (2007) 4,
474–480
10 D. Das, K. A. Dutta, K. K. Ray, Correlation of microstructure with
wear behavior of deep cryogenically treated AISI D2 steel, Wear,
267 (2009) 9–10, 1371–1380
11 D. Das, K. A. Dutta, V. Toppo, K. K. Ray, Inconsistent wear behavior
of cryotreated tool steels: the role of mode and mechanism, Materials
Science and Technology, 25 (2009) 10, 1249–1257
P. SUCHMANN et al.: DEEP CRYOGENIC TREATMENT OF H11 HOT-WORKING TOOL STEEL
42 Materiali in tehnologije / Materials and technology 49 (2015) 1, 37–42
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
NUMERICAL PREDICTION OF THE COMPOUND LAYER
GROWTH DURING THE GAS NITRIDING OF Fe-M BINARY
ALLOYS
NUMERI^NO NAPOVEDOVANJE RASTI SPOJINSKE PLASTI MED
PLINSKIM NITRIRANJEM BINARNIH ZLITIN Fe-M
Ramdhane Kouba, Mourad Keddam
Laboratoire de Technologie des Matériaux, Département SDM, Faculté de Génie Mécanique et Génie des Procédés, USTHB BP 32 El-Alia
16111 Alger, Algeria
r_kouba11@yahoo.fr
Prejem rokopisa – received: 2013-10-03; sprejem za objavo – accepted for publication: 2014-01-20
A numerical model was developed to simulate the gas nitriding of Fe-M binary alloys. The suggested model takes into account
the nitrogen diffusion in the compound layer and in the diffusion zone as well as the displacement of the (’/
) interface. The
precipitation of fine MN nitrides in the diffusion zone was also considered during the modelling. A numerical resolution of the
problem was done with the front-tracking method via the finite-difference technique. The model was capable of predicting the
growth of the ’-compound layer and also the nitrogen-depth profile.
In order to validate the model, the compound-layer thicknesses obtained with the simulation were compared with those obtained
experimentally for the nitrided Fe-Cr binary alloys. Good concordance between the numerical results and the experimental data
was noticed.
Keywords: gas nitriding, modelling, compound layer, front-tracking method
Razvit je bil numeri~ni model za simulacijo plinskega nitriranja binarnih zlitin Fe-M. Predlagani model upo{teva difuzijo du{ika
v spojinski plasti in tudi v difuzijski plasti, kot tudi premik stika (’/
). Model upo{teva tudi izlo~anje drobnih MN-nitridov v
difuzijski coni. Numeri~na re{itev problema je bila izvr{ena z uporabo metode sledenja fronte z uporabo tehnike kon~nih
diferenc. Model je sposoben predvidevanja rasti ’ spojinskega sloja in tudi profila du{ika v globino.
Za oceno modela je bila pri simulaciji dobljena debelina spojinske plasti primerjana z eksperimentalno dobljeno debelino pri
nitriranih binarnih zlitinah Fe-Cr. Ugotovljeno je dobro ujemanje med numeri~nimi rezultati in eksperimentalnimi podatki.
Klju~ne besede: nitriranje v plinu, modeliranje, spojinski sloj, metoda sledenja fronte
1 INTRODUCTION
Gas nitriding is a thermochemical process in which
nitrogen atoms diffuse into the material surface after the
dissociation of ammonia gas. The nitriding temperature
varies between 500 °C and 600 °C with the time duration
ranging from a few hours to a few days.
After this treatment, the nitride layers are formed and
we can distinguish a compound layer and a diffusion
zone. The compound layer, also called the white layer,
generally contains two sublayers: the ’ phase, mainly
composed of iron nitride (Fe4N) and the  phase as
Fe2N1–x. The thickness of the compound layer can reach
a value of 50 μm. Beneath the compound layer, the
diffusion zone can extend up to a depth of 1200 μm. For
binary alloys, the diffusion zone is a ferrite which
contains atomic nitrogen dissolved interstitially, with
dispersed fine metallic nitrides of the MN type where M
is a nitride-forming element. This diffusion zone is
responsible for increasing the resistance to fatigue
observed after nitriding.
The main objective of modelling nitriding is to
quantitatively describe different phenomena occurring
during the gas-nitriding treatment. Several models1–7
were reported in the literature, describing the nitriding
process for iron alloys and steels. These models allowed
the prediction of the nitriding kinetics and microstructure
of the nitrided zone.
Most of these models focus on the diffusion zone, in
which nitrogen diffusion takes place simultaneously with
the precipitation of the nitrides of the MN type (where M
is an alloying element). In the case of steels, the presence
of carbides and carbonitrides was also considered in the
modelling5 on the basis of the appropriate thermodyna-
mic data.
However, these models did not take into account the
presence of the compound layer, despite the fact that the
formation of this layer has an important influence on the
properties of a nitrided material. Indeed, the white layer
can lead to a significant improvement in tribological and
anti-corrosion properties8. In the case of the nitriding of
steels, it has been shown that the microstructure of the
compound layer has an influence on the hardening depth
in the diffusion zone9. With respect to these considera-
tions, it is interesting to include this layer in the
modelling.
The prediction of the nitriding process including
several phases (, ’ and 
) was already achieved for the
case of nitrided pure iron. This modelling was performed
Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53 43
UDK 519.61/.64:621.785.53:669.017.12 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)43(2015)
analytically and reported on in10–13, while the numerically
performed modelling was described in14–18. However, a
corresponding model of nitriding a multi-component
system, such as Fe-M binary alloys and steels, is not yet
available.
The main objective of the present work is to present
such a model, relating to the gas nitriding of a Fe-M
binary alloy. However, the numerical approach
developed in the present work only deals with a single
compound layer consisting of ’ nitride.
2 MATHEMATICAL FORMULATION OF THE
MODEL
During the gas nitriding of pure iron and binary
alloys or even steels, the formation of the compound
layer depends on the operating parameters, particularly,
the nitriding potential rN (see references15,16 for further
details). In fact, three configurations are possible:
• the absence of the compound layer (and therefore the
treatment is completed in the ferritic phase)
• the compound zone with a single-phase ’
• the compound zone with a dual-phase (/’).
In the present work, the aim was to simulate the
nitriding process with the presence of only ’ in the
compound zone. Therefore, the chosen value of rN must
correspond to this configuration according to the Lehrer
diagram16.
During the nitriding treatment, if the thermodynamic
conditions are satisfied, ’ precipitates appear in the
vicinity of the external surface if the solubility limit of
nitrogen in ferrite is reached. Thereafter, the ’ phase
tends to grow at the expense of the ferrite. This phase
transformation (
  ’) involves a displacement of the
(’/
) interface.
A diffusion problem, with one or more moving
boundaries, is commonly called the Stefan problem. Its
solving is usually performed with the use of one of the
front-tracking methods.
In19–22 several models based on this approach are
presented in order to simulate phase transformations in
metals and alloys. The model presented in20 was imple-
mented in the DICTRA software and used to simulate
the nitriding process in a Fe-N system13,14.
In the present work, a new front-tracking method is
presented considering both the long-range diffusion in
the diffusion zone and the interstitial diffusion of
nitrogen within the iron nitrides. The nitrogen diffusion
takes place simultaneously with the advancement of the
(’/
) nterface.
It should be noted that the diffusion of heavy
elements (Fe, M) was neglected as the nitrogen diffusion
controls the growth of the compound layer. This hypo-
thesis has already been adopted by several authors1–7.
The main objective of the present model is to simu-
late the growth of the ’ layer at the expense of ferrite.
However, in the diffusion zone, the nitrogen diffusion
occurs simultaneously with the precipitation of MN
nitrides which also affect the growth kinetics of the ’
layer. For this reason, this precipitation phenomenon was
also included in the modelling.
After the ’ phase is precipitated, the problem can be
represented schematically as shown in Figure 1, where
the two phases, ’ and 
, are henceforth adjacent. At any
time t, the position of the interface (’/
) is represented
by the distance  from the origin (z = 0) which
corresponds to the (’/gas) interface.
The nitrogen diffusion in each phase is governed by
the Fick’s second law. The system of Equations 1 and 2
can be established in the one-dimensional space as:
∂
∂
∂
∂
∂
∂
N
t z
D N
N
z

 
’
’ ’
’
( )=
⎛
⎝
⎜
⎞
⎠
⎟ in the ’ phase (1)
∂
∂
∂
∂
∂
∂
N
t z
D N
N
z



 


=
⎛
⎝
⎜ ⎞
⎠
⎟( ) in the 
 phase (2)
where N’ and N
 represent, respectively, the nitrogen
concentrations in the ’ and 
 phases. These concen-
trations are expressed as the numbers of moles per unit
volume and they are dependent on the depth (z) and time
(t). The nitrogen amount can also be expressed as a
nitrogen mole fraction (X’, X
) using the relations:
N X /V  ’ ’ ’= and N X /V
 
 
= , where V’ and V
 are the
molar volumes of the two phases that are considered
independent of the compositions. The concentration
may also be converted into mass fractions.
D’(N’) and D
(N
) are defined as the nitrogen
diffusion coefficients dependent on the N amount.
During the incremental time dt the interface position
advances with the value of d. In order to conserve the
number of nitrogen moles and considering the nitrogen
flux arriving at the interface and the nitrogen flux
leaving the interface, the following balance equation can
be established (Equation 3):
( )’ ’
’
’
’
X X
v
V
D
N
z
D
N
z
z
 
 








− = −
⎡
⎣⎢
⎤
⎦⎥
− −
⎡
⎣⎢
⎤
⎦⎥
=
∂
∂
∂
∂
z = 
(3)
Where v is the interface velocity expressed as v =
d/dt.
The terms X
’ and X’
 are the nitrogen molar frac-
tions at the interface of the two phases, 
 and ’, respec-
tively (refer to Figure 1). It is assumed that the thermo-
dynamic equilibrium is continuously established
between the two phases, 
 and ’ at the considered inter-
face. So, X
’ and X’
 can be read from the corresponding
phase diagram.
The diffusion problem with the moving boundaries is
then governed by the set of three partial differential
equations (Equations 1, 2 and 3). The solving of the
problem of interest requires the knowledge of the follow-
ing boundary conditions:
In the external surface (z = 0), the nitrogen concen-
tration in ’ is considered constant (Ns) corresponding to
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
44 Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53
the equilibrium between ’ and the gas mixture. There-
fore, at z = 0, N’(0,t) = Ns, at z = , N’(,t) = N’
 and
N
(,t) = N
’.
For greater depths (i.e., z > 1200 ìm), this zone
corresponds to the non-nitrided core where the nitrogen
amount in the matrix is neglected. Therefore, at z = +,
N
(,t) = 0.
3 NUMERICAL SOLVING OF THE PROBLEM
The finite-difference method is used to solve the
problem of the nitrogen diffusion with a moving boun-
dary. For this purpose, the zones occupied by both
phases, ’ and 
, are divided into the cells with thick-
nesses z’ and z
, respectively, as shown in Figure 1.
Two adjacent cells are separated by a node. It is assumed
that at time t, the number of nodes in the ’ phase is n’
and in 
 it is n
 (the numbers of cells in both phases will
then be (n’ – 1) and (n
 – 1), respectively).
The thicknesses of all the cells are kept constant
except for the last cell in ’ and the first cell in the 

phase, which are considered as time-dependent during
the simulation process (Figure 1). Indeed, as the inter-
face advances, there is an expansion of the left cell and
shrinkage of the one on the right. By introducing these
two variable cells, it becomes possible to follow the
progress of the ’/
 interface during the numerical
calculation as explained below.
When using the finite-difference method, the time is
also discretized with an equal time step (t). The solu-
tion is assumed to be known at time t, then it is calcul-
ated at the next time step (t + t). At the instant t, the
nitrogen concentration in each node is known for both ’
and 
 and so are the sizes of the two variable cells, d’(t)
and d
 (t).
In order to determine these values at time (t + t), the
process can be divided in two steps: first, a diffusion step
where the nitrogen diffusion occurs in ’ and 
 by con-
sidering the displacement of the (’/
) interface.
In the diffusion step, the resolution consists of sepa-
rate solutions of diffusion Equations 1 and 2 in their
corresponding areas. This allows a determination of the
nitrogen concentration profile for the whole nitrided
zone.
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53 45
Figure 1: Schematic diagram illustrating the problem of nitriding with the presence of ’as the compound layer. In order to carry out a numerical
resolution, the zone occupied by the two phases, ’ and 
, is divided into small cells. Two adjacent cells are separated by a node also called a grid
point. It is supposed that the number of nodes is equal to n’ and n
 in the ’ phase and 
 phase, respectively. In order to consider the
displacement of the interface, the thicknesses of the last cell in ’ and the first cell in 
 are considered variable during the process. When the cell
on the right-hand side dilates, the left one shrinks, while the sum (d’ + d
) remains constant being equal to d.
It should be noted that all the nodes of the grid are immobile, except for the last node in ’ and the first one in 
, which are mobile. At the
interface, the N concentration is kept constant in both phases with the values of N’
 and N
’. The corresponding N mole fractions (i.e., X’
 and
X
’) can be read from the phase diagram.
Slika 1: Shematski prikaz problema pri nitriranju s prisotnostjo spojinske plasti ’. Za numeri~no resolucijo je bila plast, v kateri sta dve fazi, ’
in 
, razdeljena v majhne celice. Dve sosednji celici sta lo~eni z vozlom oz. to~ko v mre`i. Predvideva se, da je {tevilo vozlov enako n’ in n
 v
’-fazi in 
-fazi. Da bi upo{tevali premik sti~ne ploskve, se kot spremenljivki v procesu vzameta debelina zadnje celice v ’ in prve celice v 
.
Medtem ko se celica na desni strani {iri, se na levi o`i. Medtem pa vsota (d’ + d
) ostaja konstantna in je enaka d.
Treba je omeniti, da so vsi vozli v mre`i nemobilni, razen zadnjega vozla v ’ in prvega v 
, ki sta mobilna. Na stiku je koncentracija N
konstantna v obeh fazah z vrednostima N’
 in N
’ Ustrezni molski dele` N (to je X’
 in X
’) se prebere iz faznega diagrama.
In order to solve the diffusion equation in ’ using the
finite-difference technique, Equation 1 is rewritten in the
following form:
∂
∂
∂
∂
∂
∂
∂
∂
N
t
D N
N
z
D N
z
N
z

 
   ’
’ ’
’ ’ ’ ’
( )
( )
= +
⎛
⎝
⎜
⎞
⎠
⎟
⎛
⎝
⎜
2
2
⎞
⎠
⎟ (4)
When the implicit scheme is used, a relation between
the unknown nitrogen concentrations in three successive
nodes (i – 1, i and i + 1) can be found. This is done sim-
ply by replacing the partial derivatives, which appear in
Equation 4, with their corresponding finite diffe-
rences23,24. The following relation can be established:
A N i t ti  ’ ( , )− +1 + B N i t ti  ’ ( , )+ + C N i t ti  ’ ( , )+ +1 =
= N i t’ ( , ) 2  i  n’ – 2 (5)
The values for Ai, Bi and Ci are obtained with the
following expressions:
[ ]
A
D i D i D i t
zi
=
+ − + −  

’ ’ ’
’
( ) ( ) ( )
( )
1 1 4
4 2
Δ
Δ
B
D i t
zi
= +1
2
2


’
’
( )
( )
Δ
Δ
[ ]
C
D i D i D i t
zi
= −
+ + + +  

’ ’ ’
’
( ) ( ) ( )
( )
1 1 4
4 2
Δ
Δ
Obviously, the diffusion coefficient D’ changes along
the depth, so the term D’(i) corresponds to its value at
the node i. The variation of the diffusion coefficient in ’
with the nitrogen amount is described in Appendix 1.
Equation 5 is valid for a regular meshing with the
same cell size, z’. However, in the vicinity of the ’/

interface, due to the presence of the variable cell, the
mesh is irregular. Therefore, the relation between the N
concentrations in the last three nodes in ’is not given by
Equation 5. An equivalent equation can be obtained by
replacing partial derivatives with finite differences in
Equation 4 as follows:
[ ]
∂
∂ 2
2
1
2 2
2
N
z
N n t t
z d z
N
z z n
  
  


’ ’ ’
’ ’ ’’
( , )
= −
=
− +
+
−
−
Δ
Δ Δ
[ ]
’ ’
’ ’
’ ’
’ ’ ’
( , ) ( , )n t t
z d
N n t t
d d z

 
 
  
− +
+
+
+
1 2Δ
Δ
Δ
Δ
(6)
∂
∂
N
z
N n t t N n t
d z
z z n
    
 

’ ’ ’ ’ ’
’ ’
’
( , ) ( , )
= −
=
+ − −
+
1
2Δ
Δ
(7)
∂
∂
D
z
D n D n
d z
z z n
    
 

’ ’ ’ ’ ’
’ ’
’
( ) ( )
= −
=
− −
+
1
2
Δ
(8)
∂
∂
N
t
N n t t N n t
t
z z n
    

’ ’ ’ ’ ’
’
( , ) ( , )
= −
=
− + − −
1
1 1Δ
Δ
(9)
By substituting these relations (6 to 9) in Equation 4,
Equation 10 can be obtained as follows:
A N n t tn    ’ ’ ’( , )− − +1 2 + B N n t tn    ’ ’ ’( , )− − +1 1 +
+ C N N n tn   
  ’ ’ ’ ’( , )− = −1 1 (10)
A
D n D n
d z
D n
zn 
   
 
 

’
’ ’ ’ ’
’ ’
’ ’( ) ( )
( )
( )
− =
− −
−1 2
2 2
+ Δ Δ ’ ’ ’( )d z
t
 +
⎡
⎣
⎢
⎤
⎦
⎥Δ
Δ
B
D n t
z dn 
 
 
’
’ ’
’ ’
( )
− = +1 1
2 Δ
Δ
C
D n D n
d z
D n
zn 
   
 
 

’
’ ’ ’ ’
’ ’
’ ’( ) ( )
( )
( )
− =
− −
+
−1 2
2 2
Δ Δ ’ ’ ’( )d z
t
 +
⎡
⎣
⎢
⎤
⎦
⎥Δ
Δ
Equations 5 and 10 represent a set of (n’ – 2) equat-
ions with (n’ – 2 ) unknowns. In order to determine the
nitrogen concentrations in the nodes at time (t + t) this
set of equations must be solved using the tridiagonal
matrix algorithm (TDMA)23.
The same analysis can be applied for the numerical
solution of the diffusion equation in the ferritic phase
(Equation 2). However, it was shown in several referen-
ces15–18 that the dependence of the diffusion coefficient
on the composition of ferrite is weak. So, in the present
paper, a constant value for diffusivity is adopted and
noted as D
.
The following relation can be established for the first
three nodes in the ferritic phase (irregular mesh, similar
to Relation 10):
A N2



’ + B N t t2 2



 ( , )+ +
+ C N t t N t2 3 2



 
( , ) ( , )+ = (11)
A
D t
d d z2
2

 


 
 

=
− Δ
+ Δ( )
, B
D t
z d2
1
2

 


 

= +
Δ
Δ
,
C
D t
z d z2
2

 


 
 

=
− Δ
Δ + Δ( )
For the other nodes (for i = 2 to n
 – 1) a relationship
similar to (5) can be given:
− − +r N i t t
 
 ( , )1 + ( ) ( , )1 2+ +r N i t t
 
  –
− + + =r N i t t N i t
 
 
’ ( , ) ( , )1 (12)
where r
D t
z





=
Δ
Δ( ) 2
As for the ’ phase, the resolution of the set of
Equations 11 and 12 indicates the nitrogen concentration
in each node of the ferritic phase at time (t + t).
4 CONSIDERING THE DISPLACEMENT OF
THE (’/) INTERFACE
After the diffusion step, the nitrogen concentration in
both phases is modified. The change in the nitrogen pro-
file leads to a displacement of the corresponding inter-
face so that the nitrogen-amount balance is conserved.
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
46 Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53
Equation 3 which gives the interface velocity must then
be included into the resolution.
The main feature associated with the present model
and compared to most of the other front-tracking
methods (given, for example, in19–21) is the inclusion of
the diffusion near the interface. The diffusion process
near the interface region was already considered in the
modelling from reference22. In the present work, the
simplest method is used to take this phenomenon into
account. For this purpose, it is supposed that the
interface is governed not only by Relation 3 but also by
Diffusion Equations 1 and 2. In fact, near the interface,
the three partial differential equations, 1, 2 and 3, are
solved simultaneously.
The technique used in the resolution consists of
replacing the partial derivatives with their finite diffe-
rences. But at this time, the sizes of variable cells are
time-dependent and are still unknown at t + t.
For example, to get the diffusion equation for ’, in
the finite-difference form near the interface, (according
to an implicit scheme), we simply substitute Relations 6
to 9 into 4 given that the term d’ is considered at the
time t + t and not at t. The following equation can be
obtained:
N n t t N n t
t
D n
N n
   
 
 
’ ’ ’ ’
’ ’
’ ’
( , ) ( , )
( )
(
− + − −
= − ⋅
1 1
2 1
Δ
Δ
[ ]
− +
+ +
⎧
⎨
⎪
⎩⎪
−
− +2 1, )
( )
( , )
’ ’ ’
’ ’t t
z d t t z
N n t t
z
Δ
Δ Δ Δ
Δ
Δ
  
 

[ ]
’ ’
’
’ ’ ’
’
( )
( ) ( )
d t t
N
d t t d t t z
N

 

  
 

+
⋅
⋅
+ + +
⎫
⎬
⎪
⎭⎪
+
−
Δ
Δ Δ Δ
[ ] [ ]
[ ]
N n t t D n D n
d t t z
     
 
’ ’ ’ ’ ’ ’
’ ’
( , ) ( ) ( )
( )
− + ⋅ − −
+ +
2 2
2
Δ
Δ Δ
(13)
It should be noted that this relationship is similar to
10, but in this case d’ is still unknown at the (t + t)
instant.
Similarly, the diffusion Equation 2 can be written
near the interface as finite differences. Then the follow-
ing relation can be established:
[ ]
N t t N t
t
D
N
d t t d t t z

 






 
 

( , ) ( , )
( ) ( )
’
2 2
2
+ −
=
=
+ + +
⎧
Δ
Δ
Δ Δ Δ
⎨
⎩
−
+
+
+
+
+
+ +
N t t
z d t t
N t t
z d t t z



 




 

( , )
( )
( , )
( )
2
3
Δ
Δ Δ
Δ
Δ Δ Δ[ ]

⎫
⎬
⎭
(14)
The thickness of the shrinking cell d
 (t + t) is not
an independent variable, but it is related to d’(t + t)
with the following relation (Figure 1):
d t t d d t t
 ( ) ( )’+ = − +Δ Δ (15)
By substituting Equation 15 into 14, the discrete form
of Equation 2 can be given near the interface as follows:
[ ]
N t t N t
t
D
N
d d t t d d t

 





 
( , ) ( , )
( ) (
’
’ ’
2 2
2
+ −
=
=
− + ⋅ − +
Δ
Δ
Δ Δ[ ]
[ ]
t z
N t t
z d d t t
N t t
z
)
( , )
( )
( , )
’
+
⎧
⎨
⎪
⎩⎪
−
−
+
− +
+
+
Δ
Δ
Δ Δ
Δ
Δ





 

2 3
[ ]
  
d d t t z− + +
⎫
⎬
⎪
⎭⎪’ ( )Δ Δ
(16)
Finite differences are also determined in the case of
Equation 3 and the following relations can then be
written:
∂
∂
N
z
N N n t t
d t t
z


 
  

’ ’ ’ ’
’
( , )
)
=
=
− − +1 Δ
( + Δ
(17)
∂
∂
N
z
N t t N
d t t
z




 


=
=
+ −( , )
)
’2 Δ
( + Δ
(18)
In order to write the derivative d/dt as a finite
difference, it is worth noting that the change in the
position of the interface is related only to the change in
the expanding cell size. So, the following relation can be
written:
d
d
d(
d
’   
t
d
t
d t t d t
t
= =
+ −’ ’( ) ( )Δ
Δ
(19)
Substituting these three last relations (17, 18 and 19)
into Equation 3, the following relation can be obtained:
d t t d t
tV
X X
D n
N n
 

 
 

 
 
’ ’
’
’ ’
’ ’
’ ’
( ) ( )
( )
( )
(
+ − =
−
⋅
⋅
Δ
Δ
− + −
+
−
− +
− +
1 2, )
( )
( , )
(
’
’
’
’
t t N
d t t
D
N N t t
d d t t
Δ
Δ
Δ
Δ
 





 

[ ])
⎧
⎨
⎪
⎩⎪
⎫
⎬
⎪
⎭⎪
(20)
The above analysis has finally led to the set of three
nonlinear equations (13, 16 and 20) with the following
unknowns: d’(t + t), N’(n’ – 1,t + t) and N’(2,t + t).
So, after the diffusion step, the consideration of the
interface displacement is performed by solving this set of
equations in order to determine the corresponding
unknowns. The resolution requires the use of an iterative
algorithm with the Newton-Raphson method.
Once the resolution is achieved (i.e., d’(t + t) is
calculated), the interface position at the time (t + t) can
be determined from Equation 21:
   ( + Δ ( ( + Δ (t t t d t t d t) ) ) )’ ’= + − (21)
The term  = d’(t + t) – d’(t) represents the
advancement of the interface during a time step.
The thickness of the shrinking cell is also calculated
by using Relation 15.
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53 47
The two steps of the resolution (i.e., the nitrogen
diffusion and the interface displacement) are afterwards
repeated for the next time step.
It should be mentioned that the nitrogen concen-
tration in the first neighbouring nodes of the interface
(i.e., N’(n’ – 1,t + t) and N’(2,t + t) is computed in
both diffusion step and interface-displacement step. For
the next time step, only the values resulting from the
interface-displacement analysis are retained.
During the simulation process, there is an expansion
of the left cell and shrinkage of the one on the right.
However, this process cannot continue indefinitely and
the remeshing may become necessary if d’ is too large or
d
 becomes too small. So, the two following criteria have
been adopted: The expanding cell splits into two cells if
its size becomes 1.5 times greater than the original size.
This division is performed by inserting a new node
inside the cell. The nitrogen concentration in the new
node is obtained with a linear interpolation from the
concentration in the neighbouring nodes. Similarly, if the
shrinking cell reaches half of its original size, a grid
point is removed from 
 in order to increase d
 to its
initial value. It can be seen that when the simulation
process is performed, there is a creation of nodes and
cells in phase ’ and a disappearance of others in 
.
5 APPLICATION OF THE MODEL FOR THE
NITRIDED Fe-M BINARY ALLOYS
In fact, the configuration presented in Figure 1 does
not appear at the first instant at the beginning of the
nitriding process, but there is a period of nitrogen
enrichment in the ferrite which precedes the formation of
the compound layer.
This incubation time is more important when the M
alloying element is present in the ferritic matrix because
of the consumption of a large amount of nitrogen in the
MN nitride precipitation. In addition, if the nitriding
potential rN is not too high, as assumed in the present
work, the incubation time is even longer25. For these
reasons, the early stage of the process, which takes place
completely in the ferritic phase, must be considered.
Before the ’ formation, the diffusion of nitrogen is
carried out exclusively in the ferrite zone according to
Equation 2. In this case, at z = 0, there is a ferrite/gas
contact. The transfer of nitrogen from the gas mixture to
the metal is expressed with the following relationship:
At z = 0,
[ ]− = −
=
D
N
z
K N N t
z






 
 

∂
∂ eq
( , )0 (22)
Equation 22 is obtained by considering a thermo-
dynamic equilibrium between the solid (the 
 phase) and
the gas mixture. The term Neq
 indicates the nitrogen
concentration in ferrite at the thermodynamic equili-
brium between 
 and the gas mixture. N
(0,t) represents
the actual concentration at the surface and K
 is the
kinetic constant. Equation 22 is regarded as an external
boundary condition when solving the diffusion problem
in the preliminary stage.
The nitriding process during the preliminary stage
can be regarded as the nitrogen diffusion in one phase
(ferrite) coupled to the precipitation of fine MN nitrides.
The details of the calculation procedure are given in1.
After the diffusion step which causes a modification of
the nitrogen profile, a test of the MN precipitation is
realized at each depth using the solubility-product
method. So, the amount of the nitrogen precipitated as
MN and the quantity which remains dissolved in the
ferrite can be calculated. It is also possible to determine
the proportion of M consumed by the precipitation and
the residual part dissolved in the matrix1.
The nitrogen transfer from the gas mixture to the
sample, according to Equation 22, means that the N
concentration in the ferrite at z = 0 increases with time.
The formation of ’ takes place when the surface concen-
tration reaches the solubility limit of the nitrogen in
ferrite. Therefore, the condition of the precipitation of ’
on the top of 
, at z = 0, is satisfied if: N
 (0,t)  N
’.
This test is performed for each time step during the
calculation process in the preliminary stage.
It should be noted that the kinetic considerations
related to the ’ nucleation have been neglected which
makes the formation of the compound layer take place as
soon as the nitrogen solubility limit is reached.
Subsequently, a small ’ layer with a thickness not
exceeding 0.1 μm is inserted at the top of the first cell of

 having an initial thickness of 1 μm.
The introduction of a small thickness of the ’ layer
is the beginning of the diffusion problem with a moving
boundary. Thereafter, the front-tracking solving process
is conducted as explained in detail in Sections 3 and 4.
However, when solving a diffusion problem with
moving boundaries for binary alloys, there is not just the
nitrogen diffusion in the phase (ferrite) on the right, but
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
48 Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53
Figure 2: Schematic diagram illustrating a full simulation of nitriding
Fe-Cr alloys. The diagram shows the formation of CrN precipitates
and evolution of the variable cell size: a) during the early stage, b)
after the formation of the compound layer (at time t1), c) for a time
t2 > t1
Slika 2: Shematski prikaz polne simulacije nitriranja zlitin Fe-Cr.
Diagram prikazuje nastanek izlo~kov CrN in razvoj spremenljivke
velikosti celic: a) v za~etni fazi, b) po nastanku spojinske plasti (v
~asu t1), c) za ~as t2 > t1
also the precipitation of MN nitrides which continues to
occur after the ’ formation. The precipitation of MN
nitrides is considered during the simulation in the same
way as in the preliminary stage.
A schematic illustration of a full simulation of the
gas nitriding of Fe-M binary alloys is shown in Figure 2.
A general scheme giving different steps of the calcul-
ation is presented in Figure 3. A computer code written
in the MATLAB language was used to achieve the
numerical resolution and the necessary data for the
implementation of the simulation are presented and
discussed in Appendix 2.
6 SIMULATION RESULT
Before presenting the results of the model for
nitriding Fe-Cr binary alloys, the model was first applied
to a Fe-N binary system. The case of nitriding pure iron
is quite interesting since the analytical solution of the
problem is already available10–12. Therefore, a compa-
rison between the results of the analytical model11 and
those obtained with the present work for Cr % = 0 is
possible. For this purpose, a constant diffusion coeffi-
cient was adopted for ’ instead of the composition-
dependent diffusivity. The same data, used in11, was
adopted for the numerical resolution and the incubation
time was also ignored.
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53 49
Figure 4: Results given by the model for the nitriding of Fe – 1 % Cr
binary alloy for the given nitriding conditions: T = 550 °C, rN = 0.7
bar–1/2, time = 16 h: a) nitrogen-concentration profile for the com-
pound layer and the beginning of the diffusion zone, b) profile of
nitrogen inside the compound layer, c) total nitrogen profile for the
diffusion zone
Slika 4: Rezultati, ki jih da model za nitriranje binarne zlitine Fe –
1 % Cr v danih razmerah nitriranja: T = 550 °C, rN = 0,7 bar–1/2, tra-
janje = 16 h: a) profil koncentracije du{ika v spojinski plasti in za~etek
difuzijske cone, b) profil koncentracije du{ika v spojinski plasti, c)
skupni profil du{ika v difuzijski coni
Figure 3: Flow chart showing a global scheme of the simulation of the
nitriding process
Slika 3: Diagram prikazuje celotno shemo simulacije poteka postopka
nitriranja
It is noticed that the numerical solution converges
towards the analytical one provided that the spatial and
temporal discretization steps are as small as possible.
Indeed, the following parameters (z’ = 0.1 μm, z
 =
0.1 μm and t = 0.1 s) provide a good accuracy. These
parameters were adopted for all the numerical calcul-
ations done in the present study. A good agreement
between the numerical and analytical results represents,
in fact, the first step in the validation of the model.
Figure 4a shows the simulated nitrogen-concen-
tration profiles through the ’ layer and the beginning of
the diffusion zone obtained after nitriding the Fe – 1 %
Cr binary alloys (for the following conditions: T = 550
°C, KN = 0.7 bar–1/2 and t = 16 h).
For the compound layer, despite the narrow com-
position range, Figure 4b shows that the profiles are not
linear, in contrast to the assumption made in15.
Figure 4c shows the total nitrogen profile within the
diffusion zone. The major part of nitrogen is in the form
of the CrN precipitate. The results are the same as those
obtained in1.
In this last reference1, the experimental nitrogen
profile fits well with the simulation results despite the
fact that some aspects of the problem were not taken into
account (for example, the nitrogen excess and the
nucleation and growth kinetic of the formation of MN
precipitates). These considerations can be seen as part of
the validation of the present model with respect to the
diffusion zone as long as the same solubility product
KCrN is used.
Figure 5 shows, for the same alloy and for two diffe-
rent temperatures (520 °C and 550 °C) the evolution of
simulated compound-layer thickness versus time. In this
figure, it is possible to deduce the corresponding values
of incubation times, and to confirm the parabolic regime
of the growth law of the ’ layer.
In the same figure, the growth of the ’ layer in the
case of the Fe-N system (w(Cr) = 0 %) is also presented.
It can be seen that the presence of alloying element
Cr slows down the growth of the compound layer. This
result can be attributed to the CrN precipitation. Indeed,
on the one hand, the precipitation reaction near the outer
surface during the preliminary phase delays the forma-
tion of iron nitrides by increasing the corresponding
incubation time. On the other hand, the CrN precipitation
accentuates the nitrogen concentration gradient in the
ferrite near the interface (on the right-hand side) because
of the consumption of a large amount of nitrogen
initially dissolved in the ferrite. Thereafter, the nitrogen
flux leaving the interface is increased which leads to a
reduction in the interface velocity.
7 EXPERIMENTAL VALIDATION OF THE
MODEL
In order to validate the model, nitriding experiments
were carried out on the Fe – 1 % Cr binary alloy. The
measurements of the compound-layer thicknesses were
subsequently realized with the purpose of comparing
them with the simulated results.
The samples of the Fe – 1 % Cr alloy were gas
nitrided in the laboratory nitriding furnace at 520 °C and
550 °C during different treatment times for the same
nitriding potential (KN = 0.7 bar–1/2). The identification of
the phases formed after nitriding was performed using
X-ray diffraction. The measurements of the thicknesses
were performed using a scanning electron microscope at
different locations of the cross-sections of the nitrided
samples. The mean value of the compound-layer thick-
ness was then taken from different measurements.
Figure 6 shows a micrograph of the nitrided zone,
where one can distinguish the compound layer and the
beginning of the diffusion zone. Figure 7 shows the
X-ray diffraction pattern of the gas-nitrided sample at T
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
50 Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53
Figure 5: Comparison between the ’ layer thicknesses obtained with
the model (curve) and those measured experimentally (marks) for rN =
0.7 bar–1/2
Slika 5: Primerjava med debelino plasti ’, dobljene z modelom (kri-
vulja) in eksperimentalno izmerjene (to~ke) pri rN = 0,7 bar–1/2
Figure 6: SEM micrograph of the cross-section of nitrided Fe – 1 %
Cr alloy for the nitriding conditions: (T = 550 °C, t = 10 h, KN = 0.7
bar–1/2)
Slika 6: SEM-posnetek prereza nitrirane zlitine Fe – 1 % Cr pri raz-
merah nitriranja: (T = 550 °C, t = 10 h, KN = 0,7 bar–1/2)
= 550 °C during 10 h of treatment. It can be seen that the
detected phase includes ’nitride and ferrite with no CrN
which is probably due to its small fraction. It can also be
seen that the nitriding conditions did not allow the
formation of  nitride.
The comparison between the experimental values of
the ’ layer thicknesses and those calculated by the
model is shown in Figure 5. It can be seen that there is
good agreement between the experimental thicknesses
and the simulated ones. This comparison leads to the
conclusion that the model predicts, in a good way, the
growth kinetics of the compound zone.
8 FURTHER POSSIBILITIES OF THE MODEL
The model provides good results concerning the
growth of the compound layer, at least for the considered
range of nitriding conditions and for low concentrations
of the Cr alloying element. The model can be extended
to perform a bilayer configuration (/’) of the com-
pound zone. The model can also be extended to simulate
the nitriding of alloyed steel.
This requires a method for handling, the presence of
several alloying elements with the possibility of different
types of MN precipitates. The modeling of nitriding
steels necessitates, especially, the consideration of the
carbon element which brings significant changes to the
microstructure before and during nitriding.
For example, it is reported in several studies5,26 that
the alloyed carbides transform into nitrides or carboni-
trides during the process and that the nitrogen interstitial
diffusion is accelerated by that of carbon.
One of the solutions that can be adopted to simulate
the nitriding of steels is to combine the present model
with a thermodynamic model using the CALPHAD
approach as shown in5 for the diffusion zone.
If the present model is interfaced with appropriate
thermodynamic and kinetic data (resulting from thermo-
dynamic calculations) it becomes a more complete
model for the nitriding of steels. Such a model is able to
predict nitrogen profiles in the diffusion zone and also in
the compound layer.
Furthermore, the present model can be extended to
study other processes involving moving boundaries, such
as phase transformations and solidification. The model
can even be extended to consider a case of a multicom-
ponent system where several elements are able to diffuse
as reported in20–22. However, it requires the use of a
comprehensive thermodynamic description necessary to
determine different diffusivities and chemical potentials
required for handling a diffusion process. However, in
the present work, the model was restricted to the
simulation of a nitriding process for binary alloys.
9 CONCLUSION
A numerical model for a simulation of the gas
nitriding of Fe-M binary alloys is presented. The model
takes into account both the diffusion precipitation in the
diffusion zone and the nitrogen within the iron nitride ’,
also considering the (’/
) interface displacement. The
diffusion problem with a moving boundary formulated in
this way is solved through the front-tracking approach.
Because of the presence of the M alloying element, the
model also considers the precipitation of the MN nitrides
in the diffusion zone. The numerical solving was
achieved using the finite-difference method via an
implicit scheme.
The numerical resolution allows a determination of
the nitrogen profile along the depth as well as the inter-
face position which makes it possible to follow the
growth kinetics of the compound layer. There is a good
concordance between the numerical calculations and the
experimental results in the case of nitriding the Fe – 1 %
Cr alloy.
Despite the fact that the model is limited to the case
of a monolayer compound zone, it can be used as a tool
for controlling the nitriding process and optimizing the
desired properties.
Moreover, the applicability of the model can be ex-
tended to simulate the nitriding of steels using an appro-
priate thermodynamic and kinetic database in the
simulation.
APPENDIX 1: VARIATION IN THE DIFFUSION
COEFFICIENT IN ’ WITH THE COMPOSITION
The implementation of the model requires a determi-
nation of the dependence of the intrinsic diffusion
coefficient on the nitrogen amount in the ’ phase. On
the basis of thermodynamic and kinetic considerations,
the following expression can be established:
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53 51
Figure 7: X-ray diffraction pattern of the surface of gas-nitrided Fe –
1 % Cr alloy at T = 550 °C during 10 h of treatment
Slika 7: Rentgenska difrakcija povr{ine plinsko nitrirane zlitine Fe –
1 % Cr 10 h pri T = 550 °C
D N D
a
N   
’ ’ ’
*
’
( )
ln
ln
=
d
d
N
(A1)
where aN is the nitrogen activity in ’ and D*’ is the
nitrogen-tracer diffusion coefficient which is considered
independent of the composition. The term 

=
d
d
Nln
ln ’
a
N
is often called the thermodynamic factor.
During the gas nitriding process with the presence of
a gas mixture (NH3/H2), it is assumed there is a thermo-
dynamic equilibrium between the solid (’) and the
atmosphere. It was shown that the activity of nitrogen is
proportional to the nitriding potential of rN = PNH3/PH2
15,16. The thermodynamic factor can then be rewritten as
follows:


=
d
d
Nln
ln ’
r
N
(A2)
Using the absorption-isotherm theory, Somers et al.15
gave the relation between the nitrogen amount in ’and
the nitriding potential, expressed with the following
equation:
y
T
r
r
r
N,

’
’
exp .= + − +⎛⎝
⎜ ⎞
⎠
⎟⎡
⎣⎢
⎤
⎦⎥
⋅ −
1
4
1
7558
2 978
N
N,
0
N, ’
0
Nr
⎡
⎣
⎢
⎤
⎦
⎥
⎧
⎨
⎩
⎫
⎬
⎭
(A3)
where yN,’ is the nitrogen amount expressed as a site
fraction. The meanings and values of the constants
appearing in A3 are given in the same reference15. In
order to obtain the nitrogen concentration (the number
of moles per volume), the following relationship can be
used:
N
y
V y

 
’
’
’ ’( )
=
+
N,
N,1
(A4)
For each composition, Relations A3 and A4 can be
combined and derived in order to calculate the thermo-
dynamic factor  (according to A2) and then the corres-
ponding diffusivity (according to A1). In this way, the
composition dependence of diffusivity was taken into
account in the model and implemented as explained
above.
Figure 8 shows the evolution of the diffusion coeffi-
cient D’ versus the mass fraction of N, at T = 550 °C and
rN = 0.7 bar–1/2. A strong dependence on the concentra-
tion can be observed.
APPENDIX 2: DATA USED IN THE
SIMULATION
The data required for the simulation are gathered in
Table 1.1,13,15,16,18,27 It is worth noting that when consi-
dering the nitrogen concentrations in both the external
surface (Ns) and at the interface (N’
, N
’), the Fe-N
binary phase diagram was used instead of an isothermal
section of the Fe-N-Cr phase diagram.
Table 1: Different constants used for implementing the model with
their source references
Tabela 1: Razli~ne konstante, uporabljene v modelu in reference
njihovega izvora
Constant Value at 520 °C Value at 550 °C References
K
/(cm/s) 5.3332 E–6 7.6047 E–6 27
w(Neq
)/% 0.1807 0.2742 15,16
w(Ns)/% 5.8630 5.8621 15,16
x(X’
)/% 0.1973 0.1964 15,16
x(X
’)/% 0.0026 0.0032 15,16
w(KCrN)/%–2 0.0019 0.0033 1
D
/(cm2/s) 4.9183 E–8 7.5636 E–8 15,16
D’/(cm2/s) 8.3870 E–12 1.3303 E–11 18
V
/(cm3/mol) 7.14 7.14 13
V’/(cm3/mol) 8.22 8.22 13
mass fraction, w/%
amount (= mol) fraction, x/%
This was possible because during the preliminary
phase the CrN precipitation near the external surface
consumes practically all the dissolved Cr so that ’
nitride precipitates, thereafter, from the ferrite containing
almost no alloying element. The CrN nitride is thermo-
dynamically stable and assumed to be stoichiometric.
This simplification was already adopted in1.
In addition, in the Fe-N system the boundaries bet-
ween phases were considered as a function of the
nitriding potential instead of the total nitrogen amount.
This phase diagram is called the Lehrer diagram and the
corresponding thermodynamic description can be found
in15,16.
10 REFERENCES
1 Y. Sun, T. Bell, Materials Science and Engineering A, 224 (1997),
33–47
2 M. Gouné, T. Belmonte, J. M. Fiorani, S. Chomer, H. Michel, Thin
Solid Films, 377–378 (2000), 543–549
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
52 Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53
Figure 8: Evolution of the diffusion coefficient D’(N) versus the
nitrogen concentration in ’ (T = 550 °C, rN = 0.7 bar–1/2)
Slika 8: Razvoj koeficienta difuzije D’(N) v odvisnosti od koncen-
tracije du{ika v ’ (T = 550 °C, rN = 0,7 bar–1/2)
3 H. P. Van Landeghem, M. Gouné, A. Redjaimia, Journal of Crystal
Growth, 341 (2012), 53–60
4 R. Schacherl, P. C. J. Graat, E. J. Mittemeijer, Metall. Mater. Trans.,
35A (2004), 3387
5 S. Jegou, PhD thesis, ENSAM, Aix En Provence, France, 2009
6 J. D. Kamminga, G. C. A. M. Janssen, Surface and Coatings Tech-
nology, 200 (2006), 5896–5901
7 R. Kouba, M. Keddam, M. E. Djeghlal, Journal of Alloys and Com-
pounds, 536 (2012), 124–131
8 H. C. F. Rozendaal, P. F. Coljin, E. J. Mittemeijer, Surface Engineer-
ing, 1 (1985), 30–43
9 J. Ratajski, R. Olik, Advanced Materials Research, 83–86 (2010),
1025–1034
10 M. Keddam, M. E. Djeghlal, L. Barralier, Materials Science and
Engineering A, 378 (2004), 475–478
11 L. Torchane, P. Bilger, J. Dulcy, M. Gantois, Metallurgical and
Materials Transactions A, 27A (1996), 1823–1834
12 M. Keddam, M. E. Djeghlal, L. Barralier, E. Salhi, Annales de
Chimie-Science des Matériaux, 28 (2003), 53–61
13 H. Du, J. Ågren, Metallurgical and Materials Transactions A, 27A
(1996), 1073–1080
14 H. Du, J. Ågren, Z. Metallkunde, 86 (1995), 522–529
15 M. A. Somers, E. J. Mittemeijer, Metallurgical and Materials Tran-
sactions A, 26A (1995), 57–74
16 E. J. Mittemeijer, M. A. J. Somers, Surface Engineering A, 13
(1997), 483–497
17 P. B. Friehling, M. A. J Somers, HTM J, 57 (2002), 415–420
18 T. Belmonte, M. Gouné, H. Michel, Materials Science and Engineer-
ing A, 302 (2001), 246–251
19 T. C. Illingworth, I. O. Golosnoy, Journal of Computational Physics,
209 (2005), 207–225
20 S. Crusius, G. Inden, U. Knoop, L. Höglund, J. Ågren, Zeitschrift für
Metallkunde, 83 (1992), 673–678
21 H. Larsson, R. C. Reed, Acta Materialia, 56 (2008), 3754–3760
22 J. Svoboda, E. Gamsjager, F. D. Fischer, Y. Liu, E. Kozeschnik, Acta
Materialia, 59 (2011), 4775–4786
23 L. V. Fausett, Applied Numerical Analysis Using Matlab, Prentice
Hall, New Jersey 1999
24 J. W. Thomas, Numerical partial differential equations: Texts in
Applied Mathematics, vol. 33, Springer, New York 1999
25 P. B. Friehling, F. W. Poulsen, M. A. J. Somers, Z. Metallkd, 92
(2001), 589–593
26 S. Jegou, L. Barralier, R. Kubler, Acta Materialia, 58 (2010),
2666–2676
27 H. C. F. Rozendaal, E. J. Mittemeijer, P. F. Colijn, P. J. Van Der
Schaff, Metallurgical and Materials Transactions A, 14A (1983),
395–399
R. KOUBA, M. KEDDAM: NUMERICAL PREDICTION OF THE COMPOUND LAYER GROWTH ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 43–53 53

J. SEDLÁK et al.: ANTIBACTERIAL COMPOSITE BASED ON NANOSTRUCTURED ZnO MESOSCALE PARTICLES ...
ANTIBACTERIAL COMPOSITE BASED ON NANOSTRUCTURED
ZnO MESOSCALE PARTICLES AND A POLY(VINYL CHLORIDE)
MATRIX
PROTIBAKTERIJSKI KOMPOZIT NA OSNOVI
NANOSTRUKTURNIH DELCEV ZnO IN OSNOVE IZ POLIVINIL
KLORIDA
Jakub Sedlák1,2, Pavel Ba`ant1,2, Jiøí Klofá~1,2, Miroslav Pastorek1,3, Ivo Kuøitka1,2
1Centre of Polymer Systems, University Institute, Tomas Bata University in Zlin, Nad Ovcirnou 3685, 760 01 Zlin, Czech Republic
2Polymer Centre, Faculty of Technology, Tomas Bata University in Zlin, Namesti T.G. Masaryka 275, 762 72 Zlin, Czech Republic
3Department of Polymer Engineering, Faculty of Technology, Tomas Bata University in Zlin, Namesti T.G. Masaryka 275, 762 72 Zlin, Czech
Republic
j1sedlak@ft.utb.cz
Prejem rokopisa – received: 2013-10-04; sprejem za objavo – accepted for publication: 2014-03-10
The microwave-assisted solvothermal synthesis of inorganic ZnO nanoparticles is a facile method yielding a broad variety of
active fillers with specific properties. The synthesis of hierarchical nanostructured zinc oxide mesoscale particles was carried
out under continuous microwave irradiation from soluble zinc acetate as the precursor in a diethylene glycol solvent. The pre-
pared ZnO particles were characterized with X-ray diffractometry, scanning electron microscopy and UV-VIS spectrometry. A
composite was obtained by mixing these particles with the softened medical-grade poly(vinyl chloride) as a model polymer
matrix to develop an antibacterial polymer system with a possible application for plastic medical devices. The testing of the
antibacterial activity of the composite confirmed an excellent performance against Escherichia coli (gram-negative) and
sufficient activity against Staphylococcus aureus (gram-positive) bacteria according to ISO 22196: 2007. In addition, no adverse
effects of the filler on the mechanical properties of the composite were observed in comparison with the neat PVC resin. There-
fore, the prepared composites can be considered as suitable candidates for application in plastic medical devices and other
industries.
Keywords: antibacterial, PVC, composite, ZnO, nanoparticles
Solvotermi~na sinteza anorganskih delcev ZnO v mikrovalovni pe~ici je lahka metoda za izdelavo razli~nih polnil s posebnimi
lastnostmi. Sinteza delcev cinkovega oksida s hierarhi~no nanostrukturo je bila izvr{ena s kontinuirnim mikrovalovnim obse-
vanjem v mikrovalovni pe~ici iz predhodno raztopljenega cink acetata v raztopini dietilen glikola. Pripravljeni delci ZnO so bili
karakterizirani z rentgensko difrakcijo, vrsti~no elektronsko mikroskopijo in z UV-VIS-spektrometrijo. Kompozit je bil izdelan
z me{anjem teh delcev v zmeh~an medicinski polivinil klorid kot modelno polimerno osnovo za razvoj protibakterijskega
polimernega sistema z mo`nostjo uporabe za plasti~ne medicinske naprave. Preizku{anje protibakterijske aktivnosti kompozita
je pokazalo odli~no vedenje pri Escherichia coli (gram-negativne) in zadovoljivo aktivnost proti bakterijam Staphylococcus
aureus (gram-pozitivne) skladno z ISO 22196: 2007. Dodatno pa v primerjavi s ~isto PVC-smolo ni bil ugotovljen ne`elen
u~inek polnila na mehanske lastnosti kompozita. Zato je pripravljeni kompozit primeren kandidat za uporabo v plasti~nih
medicinskih napravah in v drugih industrijah.
Klju~ne besede: protibakterijski, PVC, kompozit, ZnO, nanodelci
1 INTRODUCTION
Nosocomial infections have become a serious
problem of today’s hospital environment. The prevalence
of medical-devices-related infections (MDIs) rapidly
grows among hospital-acquired infections and high
health risks are especially associated with the devices
which are in close contact with the inner parts of a
patient’s body.1 As the polymers are the most frequently
used materials in medical devices, the prevention against
bacterial adhesion and growth on their surfaces is an
absolute imperative.2 There are two principally different
strategies of how to resolve this problem. The first
approach uses organic antimicrobial additives to plastics.
They are very efficient; however, their mechanism of
action requires diffusion out of the plastic surface and
they can be classified as migratory additives. In contrast,
the second approach uses inorganic additives with a very
low solubility which ensures sufficient surface antibacte-
rial activity and a minimum leakage of the active species
into the patients body.3 Therefore, a utilization of inor-
ganic antibacterial nanopowders by compounding them
with commercial polymer resins seems to be the simplest
way leading to an enhanced antibacterial activity of a
polymer composite system.
In recent years, a great effort has been made by the
community of researchers for the synthesis and utiliza-
tion of nanomaterials in the field of medical materials.4
Zinc oxide is one of the promising candidates from the
group of metal oxides (e.g., CaO, MgO, TiO2, Cu2O) that
can be obtained with various properties depending on the
nanoparticle size and shape as the crucial parameters.5
ZnO is a biocompatible material exhibiting antimicrobial
properties against the gram-negative as well as gram-
positive bacteria. The mechanisms of the antibacterial
Materiali in tehnologije / Materials and technology 49 (2015) 1, 55–59 55
UDK 678.7:66.017:620.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)55(2015)
action are still not perfectly understood, but there are
several proposals for the explanation of the microbial
inhibition such as the generation of reactive oxygen
species (ROS) like hydroxyl radicals and singlet oxygen,
or the release of Zn2+ cations.6,7
Up to now, a large variety of physical and chemical
methods have been developed for the ZnO-nanoparticle
preparation such as thermal decomposition, hydrother-
mal, solvothermal or microwave-assisted syntheses. The
last mentioned microwave energy exclusively reduces
the reaction time as well as the energy demands of the
process.8–12
Here we report on a fast and simple solvothermal
microwave-assisted synthesis of hierarchical nanostruc-
tured mesoscale ZnO aggregates, their compounding into
medical-grade plasticized PVC as a model polymer
matrix suitable for utilization in medical devices and an
evaluation of the surface antimicrobial and mechanical
properties of the final composites.
2 EXPERIMENTAL WORK
2.1 Materials
All the chemical reagents used in the experiment
were of the analytical grade and used as received without
further purification: zinc acetate dihydrate, p.a., (a purity
of 99 %, Penta, The Czech Republic), diethylene glycol,
p.a., (a purity of 99 %, Penta, The Czech Republic).
Azeotropic denatured ethanol (The Czech Republic) was
used for washing the product. The medical-grade
softened PVC resin RB-3 (Modenplast Medical, Italy)
was used as the polymer matrix.
2.2 Synthesis of the ZnO antibacterial filler
At first 2.195 g of Zn(CH3COO)2 · 2H2O was com-
pletely dissolved in 100 mL of diethylene glycol at a
temperature of about 100 °C during a 30 min agitation
on a magnetic stirrer. The obtained solution was trans-
ferred to a 250 mL boiling flask and irradiated under
reflux for 5 min using microwaves. The product was
cooled down naturally after turning off the microwave
oven. The powder was collected with centrifugation,
washed with ethanol and air-dried up to the constant
mass. The microwave reaction was performed in a
microwave open-vessel system MWG1K-10 (Radan, The
Czech Republic, 800 W, 2.45 GHz), based on modifying
a domestic microwave oven by drilling a hole in the
ceiling for external cooling and equipped with an exter-
nal source of microwave energy that allows a control of
the duty cycle. The reaction mixture was heated in a
quasi-continuous mode at the maximum heating power.
2.3 Compounding the ZnO filler with the PVC matrix
The PVC-filler compounds were prepared with a
melt-mixing process using a Brabender Plasti-Corder
machine (Brabender, Germany) equipped with a 50 cm3
mixing chamber and the mixing elements working in a
contra-rotating mode. In this way, four mixtures differing
from each other with respect to the load of the filler in
the polymer matrix were compounded with the compo-
sitions as follows: w = (0.5, 1, 2 and 3) %. The com-
pounding process consisted of three continuous phases.
The first started by feeding the mixing chamber with
PVC pellets and synthesized filler in the form of a fine
powder at 20 r/min for two minutes. In the next step,
after the feeding of the chamber was finished, the engine
speed was continuously increased from 20 r/min up to
50 r/min within 1 min. The last part of the mixing was
done at 50 r/min for 5 min. The mixing process was
monitored by measuring the torque of the drive engine.
The processing time of 8 min was enough to obtain a
constant torque value for all the samples. No resin
discoloration or other signs of a polymer deterioration
connected with the use of the ZnO filler were observed
under the used mixing conditions. The prepared com-
pounds were compression molded to square-shaped
sheets with the dimensions of 50 mm × 50 mm × 1 mm
at 180 °C for 2.5 min. Blank samples without any filler
were prepared in the same way as the filled ones.
2.4 Characterization
The synthesized powder was characterized with
X-ray powder diffraction (XRD) using an X´Pert PRO
(PANalytical, The Netherlands) with a Cu K
 radiation
source ( = 0.1540598 nm). The micrographs of the
prepared powder and cross-section surfaces of the com-
posites were taken with a scanning electron microscope
Vega II (Tescan, The Czech Republic) and the UV-VIS
absorption spectra were collected with an Avaspec
UV-VIS spectrometer (Avantes, The Netherlands) with
an AvaLight-DHS-DUV source type and an integrating
sphere (BaSO4 coated) was used for the diffuse reflec-
tance measurement. The mechanical properties of the
composites were evaluated with the tensile tests per-
formed on a Testometric universal testing machine
M350-5CP (LABOR machine Ltd., The Czech Republic)
according to the ISO 37: 2005 standard with a type-2
testing specimen. The speed of the moving clamps was
250 mm/min. All the sample measurements were repli-
cated six times.
2.5 Antibacterial testing
An evaluation of the surface antibacterial activity of
the composites against bacterial adherence and growth
was performed according to the ISO 22196: 2007
(formerly known as JIS Z-2801) standard. Gram-positive
bacteria were represented by Staphylococcus aureus
ATCC 6538P and gram-negative by Escherichia coli
ATCC 8739, both obtained from The Czech Collection
of Microorganisms (The Czech Republic). The size of
the test specimens was 50 mm × 50 mm × 1 mm. A
HERAcell 150i incubator (Thermo Scientific, USA) was
J. SEDLÁK et al.: ANTIBACTERIAL COMPOSITE BASED ON NANOSTRUCTURED ZnO MESOSCALE PARTICLES ...
56 Materiali in tehnologije / Materials and technology 49 (2015) 1, 55–59
used for the cultivation. The antibacterial activity R was
calculated using Equation 1:
R = (Ut – U0) – (At – U0) = Ut – At (1)
where R is the antibacterial activity; U0 is the average of
the common logarithm of the number of viable bacteria
cells in 1/cm2, recovered from the untreated test speci-
mens immediately after the inoculation; Ut is the ave-
rage of the logarithm of the number of viable bacteria
cells in 1/cm2, recovered from the untreated test speci-
mens immediately after 48 h; and At is the average of
the logarithm of the number of viable bacteria cells in
1/cm2, recovered from the treated test specimens
immediately after 48 h. The colonies were counted after
24 h and checked after 48 h for a presence of slowly
growing colonies on both untreated and treated test
specimens, improving the original standard protocol. All
the tests were repeated in triplicates.
3 RESULTS AND DISCUSSION
The crystalline-phase structure of the synthesized
filler is shown on the diffractogram in Figure 1. By com-
parison with JCDD PDF-2 entry 01-079-0207, diffrac-
tion peaks were labelled as (100), (002), (101), (102),
(110), (103), (200), (112), (201), (004), (202) and (104)
and assigned to the hexagonal ZnO wurtzite structure.
Moreover, the evident broadness of diffraction lines
implies that the synthesized particles are of nano-scale
dimensions.
Figure 2 shows the diffuse reflectance UV-VIS
(DR-UV-VIS) spectrum of the prepared ZnO nanocry-
stals that exhibit a strong absorption with the maximum
below 380 nm. This peak is slightly blue-shifted from
the typical position of bulk ZnO13 which points towards
the nanosized particles, similarly as with the XRD anal-
ysis. The morphology of the prepared filler is presented
in the SEM image in Figure 3a. As can be seen, the
material consists of globular-like particle aggregates with
a diameter ranging from 200 nm up to 1 μm. According
to the XRD analysis and DR-UV-VIS measurement it
can be concluded that these aggregates are assembled
from much smaller nanocrystals.
The filler dispersion in the prepared composites was
achieved at the aggregate level and its good distribution
in the PVC matrix is shown in the SEM image in Figure
3b, for the sample with w = 2 % of the filler. There are
no visible agglomerates of the globular aggregates within
the polymer matrix. On the other hand, the aggregates
were not disaggregated during the mixing procedure.
The mechanical properties of the prepared compo-
sites were characterized with tensile tests, and the
Young’s modulus, elongation at break and tensile
strength were chosen as the three representative quanti-
ties for evaluating the filler-amount effects on the com-
posite performance. As summarized in Table 1, no signi-
ficant changes in the mechanical properties were
observed across all the samples. Thus, it can be
concluded that (i) the incorporation of the inorganic ZnO
filler up to w = 3 % did not change the properties of the
J. SEDLÁK et al.: ANTIBACTERIAL COMPOSITE BASED ON NANOSTRUCTURED ZnO MESOSCALE PARTICLES ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 55–59 57
Figure 3: SEM images of: a) prepared ZnO filler and b) cross-section
of the prepared composite material with w = 2 % of ZnO filler
Slika 3: SEM-posnetka: a) pripravljeno ZnO-polnilo in b) prerez
pripravljenega kompozitnega materiala z masnim dele`em ZnO-pol-
nila 2 %
Figure 1: X-ray diffractogram of the prepared ZnO filler
Slika 1: Rentgenski difraktogram pripravljenega ZnO-polnila
Figure 2: UV-VIS diffuse-reflectance spectrum of the prepared ZnO
filler. Relative reflectance is plotted versus wavelength.
Slika 2: UV-VIS-spekter razpr{ene odbojnosti pripravljenega ZnO-
polnila. Relativna odbojnost je prikazana v odvisnosti od valovne
dol`ine.
neat medical-grade plasticized PVC matrix which was
already designed by its producer to be an optimally
flexible material for medical devices such as urinal
catheters or blood bags, and that (ii) the filler had no
adverse influence on the application potential of the
prepared antibacterial composite.
Table 1: Summary of the selected mechanical properties of the
prepared composites and their standard deviations
Tabela 1: Povzetek izbranih mehanskih lastnosti pripravljenih kom-
pozitov in njihov standardni odmik
Concentration
of ZnO filler
(w/%)
Young’s
modulus
(MPa)
Elongation at
break
(%)
Tensile
strength
(MPa)
0 9.9 ± 0.2 501 ± 17 19.1 ± 0.9
0.5 9.8 ± 0.2 494 ± 30 19.0 ± 1.1
1 10.3 ± 0.3 500 ± 20 19.4 ± 0.8
2 10.1 ± 0.3 497 ± 9 19.4 ± 0.8
3 9.2 ± 1.0 520 ± 20 18.4 ± 0.4
The quantitative evaluation of the surface antibacte-
rial activity of the as-prepared ZnO/PVC composites is
summarized in Table 2. The obtained R values represent-
ing the antibacterial activity of the composites against
both bacteria are high enough for the composites filled
with w = 2 % and 3 %. The R value should not be lower
than 2 for the materials that can be classified as the ones
exhibiting an antibacterial surface suitable, for instance,
for hygienic applications or less demanding applications.
However, medical devices are more demanding with
respect to the antibacterial performance of the used
materials. The R-values of 5 or even 6 are expected for
current high-end commercial materials, especially those
with organic additives intended to be used in indwelling
medical devices.3 The surface antibacterial activity of the
composites containing w = 2 % or 3 % of the filler
against E. coli meet these highest requirements, while
the activity against S. aureus is not so high. A similar
antibacterial performance was recently obtained for the
antibacterial systems containing hybrid Ag/ZnO filler
employing nanosilver in addition to nanostructured
ZnO.14 On the other hand, S. aureus is a very vital and
resistant bacterium, so it is difficult to suppress it and
inhibit its growth. Concurrent studies reported reductions
with respect to the control samples by about 70–95 % for
much higher nano-ZnO loadings in PVC matrices, which
means that the R values of less than 2 were claimed as
success by their authors.4 However, other strains were
used throughout these studies and the surface-activity-
estimation methods were similar but not identical to
ISO 22196: 2007. From this point of view, and with a
full awareness of the peculiar comparability among
different published results, the performance of our
prepared composites seems to be sufficiently high. Next,
as the bacteria gain antibiotic resistance, the organic
molecular additives in antimicrobial polymer systems
may be found to be inefficient or ineffective and the inor-
ganic nanofillers will serve as the last line of defense.
4 CONCLUSIONS
The present study describes a microwave-assisted
synthetic route leading to the development of a hierar-
chical nanostructured ZnO mesoscale filler. The syn-
thesized powder was clearly identified to have a ZnO
hexagonal wurtzite structure and globular morphology of
the aggregated nanocrystals. The composite materials
compounded of medical-grade softened PVC and an
as-prepared filler showed excellent antibacterial-activity
values against E. coli and satisfactory antibacterial-acti-
vity values against S. aureus. Moreover, the obtained
composites kept the mechanical properties of the neat
PVC resin suitable as they were deteriorated neither by
the processing nor by the effects of the nanofiller. These
facts suggest that the prepared nanostructured mesoscale
ZnO/PVC composite has an application potential in
medicine as the material for medical devices being in the
direct contact with the human body, besides many other
possible utilizations.
Acknowledgment
The authors wish to thank for the internal grant of
TBU in Zlin, No. IGA/FT/2013/026, funded from the
resources for specific university research.
This article was written with the support of the
Operational Program "Research and Development for
Innovations" co-funded by the European Regional
Development Fund (ERDF) and the national budget of
the Czech Republic, within the Centre of Polymer
Systems project (reg. number: CZ.1.05/2.1.00/03.0111).
This article was written with the support of the
Operational Program "Education for Competitiveness"
J. SEDLÁK et al.: ANTIBACTERIAL COMPOSITE BASED ON NANOSTRUCTURED ZnO MESOSCALE PARTICLES ...
58 Materiali in tehnologije / Materials and technology 49 (2015) 1, 55–59
Table 2: Summary of the surface-antibacterial-activity-evaluation results obtained for the prepared composites
Tabela 2: Povzetek rezultatov ocene povr{inske protibakterijske aktivnosti za pripravljene kompozite
Concentration of
ZnO filler
(w/%)
Treated specimens after
48 h E. coli
N/(cfu cm–2)
Treated specimens after
48 h S. aureus
N/(cfu cm–2)
Antibacterial activity
(lg CFU) E. coli,
R = Ut – At
Antibacterial activity
(lg CFU) S. aureus,
R = Ut – At
0 4.0 × 106 7.5 × 104 Ut = 6.6 Ut = 4.9
0.5 1.5 × 105 8.0 × 104 1.4 0
1 1.7 × 105 1.0 × 100 1.4 4.9
2 < 1 2.2 × 100 > 6.6 4.5
3 < 1 1.3 × 100 > 6.6 4.8
co-funded by the European Social Fund (ESF) and the
national budget of the Czech Republic, within the
Advanced Theoretical and Experimental Studies of
Polymer Systems project (reg. number: CZ.1.07/2.3.00/
20.0104).
5 REFERENCES
1 P. M. Sivakumar, S. Balaji, V. Prabhawathi et al., Carbohydrate
Polymers, 79 (2010), 717–723
2 X. Y. Ma, W. D. Zhang, Polymer Degradation and Stability, 94
(2009), 1103–1109
3 A. Jones, Plastics Engineering, 64 (2008), 34–40
4 B. M. Geilich, T. J. Webster, International Journal of Nanomedicine,
8 (2013), 1177–1184
5 J. Sawai, Journal of Microbiological Methods, 54 (2003), 177–182
6 N. Padmavathy, R. Vijayaraghavan, Science and Technology of Ad-
vanced Materials, 9 (2008) 3, 035004, 7 pp
7 K. H. Tam, A. B. Djurisic, C. M. N. Chan et al., Thin Solid Films,
516 (2008), 6167–6174
8 Y. Yang, H. Chen, B. Zhao et al., Journal of Crystal Growth, 263
(2004), 447–453
9 H. Lu, S. Wang, L. Zhao et al., Journal of Materials Chemistry, 21
(2011), 4228–4234
10 P. Tonto, O. Mekasuwandumrong, S. Phatanasri et al., Ceramics
International, 34 (2008), 57–62
11 J. Zhu, J. Zhang, H. Zhou et al., Trans. Nonferrous Met. Soc. China,
19 (2009), 1578–1582
12 A. Phuruangrat, T. Thongtem, S. Thongtem, Materials Letters, 63
(2009), 1224–1226
13 N. S. Pesika, K. J. Stebe, P. C. Searson, J. Phys. Chem. B, 107
(2003), 10412–10415
14 P. Bazant, I. Kuritka, O. Hudecek, M. Machovsky, M. Mrlik, T. Sed-
lacek, Polymer Composites, 35 (2014) 1, 19–26
J. SEDLÁK et al.: ANTIBACTERIAL COMPOSITE BASED ON NANOSTRUCTURED ZnO MESOSCALE PARTICLES ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 55–59 59

E. ADÁMKOVÁ et al.: WATER-SOLUBLE CORES – VERIFYING DEVELOPMENT TRENDS
WATER-SOLUBLE CORES – VERIFYING DEVELOPMENT
TRENDS
JEDRA, TOPNA V VODI – PREVERJANJE SMERI RAZVOJA
Eli{ka Adámková, Petr Jelínek, Jaroslav Beòo, Franti{ek Mik{ovský
V[B-Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of Metallurgy and Foundry, Czech
Republic
eliska.adamkova@vsb.cz
Prejem rokopisa – received: 2013-10-07; sprejem za objavo – accepted for publication: 2014-01-17
Application of pure inorganic salt-based cores has been known since the end of the 20th century, especially in the field of gravity
and low-pressure die casting. The contemporary trend in technology leads to the use of the cores in the field of non-ferrous-alloy
high-pressure die casting. The main methods of the core production include high-pressure squeezing and shooting (warm core
box). During research processes it was shown that pure-salt application is not very suitable for high-pressure casting. That is
why a composite salt-based matrix of defined properties was started to be used. The aim of the paper is to verify the influences
of the chemical composition, shape and morphology of the grains of various NaCl compounds on the mechanical properties (the
bending strength) of water-soluble salt cores used for Al-alloy high-pressure die casting and to evaluate their properties
resulting from the squeezing and shooting methods.
Keywords: salt cores, inorganic salts, die casting, warm box, non-ferrous metals
Uporaba jeder iz ~istih anorganskih soli je poznana `e od konca 20. stoletja, {e posebno na podro~ju gravitacijskega in
nizkotla~nega litja. Podobne usmeritve in tehnologije vodijo k uporabi jeder na podro~ju visokotla~nega litja ne`eleznih zlitin.
Med glavnimi metodami izdelave jeder sta visokotla~no iztiskanje in vbrizganje (segreti jedrnik). Raziskava je pokazala, da
uporaba ~iste soli ni najbolj primerna za visokotla~no litje. To je razlog za za~etek uporabe kompozitnih jeder na osnovi soli.
Namen ~lanka je preveriti vpliv kemijske sestave, oblike in morfologije zrn razli~nih vrst NaCl na mehanske lastnosti (upogibna
trdnost) in uporabnost v vodi topnih jeder za visokotla~no litje Al-zlitin ter oceniti njihove lastnosti po stiskanju ali
vbrizgavanju.
Klju~ne besede: jedra iz soli, anorganske soli, tla~no litje, segret jedrnik, ne`elezne kovine
1 INTRODUCTION
Together with the development in various technical
fields (the automotive industry) the demand for more and
more difficult and challenging castings increases; these
are mechanically cleanable only with large difficulties.
The application of the technology of disposable, inor-
ganic, water-soluble salt cores is one of the solutions for
the difficulty of a core removal from the areas that are
hardly accessible for mechanical cleaning.1 A reverse
crystallization of the salt from a water solution is
enabled due to the core water solubility, which is a
requirement for creating an environmentally friendly
closed circuit of the core production.
The usage of water-soluble salt cores has been so far
known in the field of non-ferrous-alloy gravity and
low-pressure die casting.2 Promising possibilities can be
created by focusing the investigation on the development
of water-soluble salt-core application technology in the
field of Al-alloy high-pressure die casting.3,4 At present,
two salt-core manufacturing techniques are being
developed: high-pressure squeezing with the use of the
recrystallization process and shooting with the use of
inorganic binders, e.g., alkaline silicates.5,6 Considering
the material-purchase costs for the production of salt
cores from chemically pure salts, it is necessary to look
for more suitable solutions for crating the basic salt
matrix.
The aim of this paper is to verify the applicability of
NaCl (from the vast range of the common salts
commercially available in the Czech market), with the
main focus on the possibility to replace the more
expensive chemically pure salt with a cheaper common
salt. The influences of various salt origins (rock, Alpine
or sea salt), chemical compositions, shapes and surface
morphologies of the grains on the mechanical properties
(the bending strength) of the salt cores manufactured
with the high-pressure squeezing and shooting methods
are investigated.
2 SELECTION OF THE USED SALTS
To evaluate the cores manufactured with the
squeezing and shooting methods, the tested cores were
prepared from technical (common) and chemically pure
(standard) salts. The salts for the core production were
chosen according to the chemical composition declared
by the manufacturer on the packaging (Table 1). From
the wide range of commercially available salts six types
of salts were finally chosen for the observation per-
formed in this experiment. Each of the salts was of a
different chemical composition, grain shape and surface
morphology. These aspects and the selected
salt-core-production technologies had significant
influences on the final mechanical properties of the
cores.
Materiali in tehnologije / Materials and technology 49 (2015) 1, 61–67 61
UDK 621.74.043:621.743 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)61(2015)
The chemical composition of the selected types of
salts can be evaluated from two important viewpoints:
• Health
• Technology
With respect to health, salts contain healthy elements
– iodine, fluorine – in the form of chemical compounds
(KIO3, KI, NaF) in the amount of 15–58 (max. 250) mg/kg.
However, more important for the salt-core production is
the technology (the chemical composition) that affects
the quality of the manufactured cores. Some of the
commercially available salts contain the additives that
prevent the grains from sticking, i.e., the anti-sintering
additives. These are K2CO3, CaCO3, MgCO3, SiO2,
K4[Fe(CN6)] · 3H20. Alpine salts feature substantial
amounts of carbonates. These components added for the
purposes of trouble-free storage, transport and manipu-
lation have significant influences on the core-production
technology and strength properties. The mechanisms of
grain sticking and further recrystallization are disturbed
and the strength of squeezed cores is decreased. Simi-
larly, in the case of shot cores a decrease in the bending
properties of the salt grains and alkaline silicate occurs.
3 MECHANICAL PROPERTIES OF SALT CORES
The criterion for evaluating the mechanical properties
of squeezed and shot cores was the bending strength
which was measured as follows:
• after 24 h in the air (the primary strength)
• after drying at 160 °C 1 h (105 °C 1 h)
E. ADÁMKOVÁ et al.: WATER-SOLUBLE CORES – VERIFYING DEVELOPMENT TRENDS
62 Materiali in tehnologije / Materials and technology 49 (2015) 1, 61–67
Table 1: Chemical compositions of the selected salts as declared by the manufacturers
Tabela 1: Kemijska sestava izbranih soli po navedbi proizvajalca
Salt
sample Commercial name Additives Presence of compounds
1 Rock salt with iodine and fluorine, edible(fine ground)
F
I
max. 250 mg/kg
27 ± 7 mg/kg
NaF
KIO3
2 Rock salt with iodine (fine ground) I 20–35 mg/kg KIKIO3
3 Alpine salt with iodine (vacuum packed)
NaCl
CaCO3
MgCO3
KIO3
I
98.8 %
min. 0.9 %
min. 0.2 %
33–58 mg/kg
20–34 mg/kg
CaCO3
MgCO3
KIO3
4 Alpine salt with fluorine and iodine(vacuum packed)
NaCl
CaCO3
MgCO3
KIO3
I
F
98.0 %
min. 0.7 %
min. 0.1 %
33–58 mg/kg
20–34 mg/kg
max. 250 mg/kg
KIO3
CaCO3
MgCO3
NaF
5 Fine sea salt with iodine (evaporatedfrom sea water)
I 15–35 mg/kg KIO3
6 NaCl, p.a. (chemically pure)
Fe
Heavy metals (Pb)
SO4
Ca
Mg
I
Br
max. 0.0003 %
max. 0.0005 %
max. 0.005 %
max. 0.005 %
max. 0.002 %
max. 0.008 %
max. 0.01 %
Figure 1: Comparison of the mean bending strengths of 6 types of
shot salt: a) after drying and b) after 24 h in the open air (the mean
value for 6 cores; A – original granularity, B – fraction of 0.063–1.0
mm)
Slika 1: Primerjava povpre~ne upogibne trdnosti 6 vrst vbrizganih
soli: a) po su{enju in b) po zadr`anju 24 h na zraku (srednja vrednost 6
jeder; A – originalna zrnatost, B – frakcija od 0,063 mm do 1,0 mm)
3.1 Shot cores
For the shot-core production a mixture of the
originally supplied granulometry and a mesh fraction of
0.063–1.0 mm was used. The conditions for the shot-
core production are as follows:
• Mixture composition: 100 mass parts of NaCl; 5
mass parts of Na – sodium silicate (M = 1.85)
• Hardening time: 50 s
• Core-box temperature: 190 °C
• Shooting speed: 7.5 s
• Air pressure: 7.5–8 bar
From Figure 1 it is evident that for all the salts,
except the sea salt with iodine (sample No. 5), the remo-
val of dust proportion (< 0.063 mm) led to a bending-
strength decrease. The final drying of the cores
(160 °C/1 h) led to a bending-strength increase only for
the NaCl sample (No. 5). Strengths comparable to the
standard strength were achieved for rock salts with I and
IF (Nos. 1, 2). The lowest strengths were achieved for
the cores made of Alpine salt with I and IF (samples
Nos. 3, 4). It was shown that the salt grain shape and the
grain-surface quality are of the highest importance.
The apparent porosity m was calculated with the
following relation:
m =
− 

NaCl v
NaCl
· 100 (%) (1)
where:
v = bulk density of cores (g/cm3)
The apparent porosity of shot cores ranged from
35 % to 45 %. After sieving the dust (< 0.063 mm) off
the apparent porosity, an increase (a decrease in the
volume mass v and strength) occurred for all the salt
samples except for sample No. 5 (sea salt with I), which
underwent an extreme increase in the bending strength,
especially after being dried at 160 °C 1 h (Figure 2).
3.2 Squeezed cores
For the production of squeezed cores, a mixture
granulometry adjustment was carried out: the fraction of
0.063–1.0 mm. The conditions for the squeezed-core
production are as follows:
• NaCl humidity: 0.65–1.04 %
• squeezing force: 200 kN
• squeezing tension: 104 MPa
• load rate: 9 kN/s
After the squeezing significantly higher strengths
were achieved for all the salt samples (Figure 3). The
increase in the strength (together with a decrease in the
apparent porosity) was evident for dried cores (105 °C
1 h). It was shown that the lowest strengths were
achieved for the cores made of the salts with regular,
E. ADÁMKOVÁ et al.: WATER-SOLUBLE CORES – VERIFYING DEVELOPMENT TRENDS
Materiali in tehnologije / Materials and technology 49 (2015) 1, 61–67 63
Figure 3: Comparison of: a) the mean bending strengths and b)
apparent porosities of 6 types of squeezed salts (the mean value for 6
cores; A – cores additionally dried at 105 °C 1 h, B – 24 h in the open
air)
Slika 3: Primerjava: a) povpre~ne upogibne trdnosti in b) navidezne
poroznosti 6 vrst iztiskanih soli (povpre~na vrednost 6 jeder; A – jedra
dodatno su{ena 1 h pri 105 °C, B – 24 h na zraku)
Figure 2: Comparison of: a) the mean bending strengths and b)
apparent porosities of 6 types of shot salts after 24 h in the open air
(the mean value for 6 cores; A – original granularity, B – fraction of
0.063–1.0 mm)
Slika 2: Primerjava: a) povpre~ne upogibne trdnosti in b) navidezne
poroznosti 6 vrst vbrizganih soli po 24 h zadr`anju na zraku (srednja
vrednost 6 jeder; A – originalna zrnatost, B – frakcija od 0,063 mm do
1,0 mm)
cubic grain shape (Alpine salt, samples Nos. 3, 4 –
recrystallized salts), whereas the highest strengths were
achieved for the cores made of crushed-rock salts
(samples Nos. 1, 2) and sea salt (sample No. 5). How-
ever, this does not apply to the standard salt (NaCl, p. a.,
sample No. 7). Regular, dipyramidal grain shape allows
high strengths of the squeezed cores as well as of the
shot ones. This shape is the closest to the globular
silica-sand grain shape.
4 GRAIN SHAPE AND GRAIN-SURFACE
MORPHOLOGY (SEM AND EDX ANALYSES)
A comparison of the chemical compositions declared
on salt packagings and the real compositions determined
with the use of an EDX microprobe is shown in Table 2.
In most cases, the grain shapes and surface morpho-
logies of the selected salts vary significantly. The diffe-
rences are due to different salt-production technologies,
e.g., crushing and milling of the rock salt or the salt
recrystallized from a solution – Alpine or sea salt.
Figures 4 to 7 below show grain shapes and surface
morphologies, together with an EDX analysis of the
investigated salts.
The differences between the salt-core strengths are
explained with a detailed investigation of the salts using
the SEM and EDX analyses. The low strengths of Alpine
salts are caused by their lower quality of grain surfaces
catching the blocking additives of MgCO3 and CaCO3;
their presence was proven with the EDX analysis
(Figure 5).
The dipyramidal grain shape of the chemically pure
NaCl, p.a., is distinctive for its high grain coordination
number, resembling the grain shape of SiO2 sand grains,
and being different from the regular, cubic Alpine-salt
grains. This shape is more suitable for both core squeez-
ing and core shooting.
E. ADÁMKOVÁ et al.: WATER-SOLUBLE CORES – VERIFYING DEVELOPMENT TRENDS
64 Materiali in tehnologije / Materials and technology 49 (2015) 1, 61–67
Figure 4: Shattered surface of crushed-rock salts (samples Nos. 1, 2) and EDX analysis of its chemical composition
Slika 4: Razbita povr{ina drobljene kamene soli (vzorca {t. 1 in 2) in EDX-analiza kemijske sestave
Table 2: Comparison of declared chemical compositions and EDX microanalysis of the investigated real compositions
Tabela 2: Primerjava deklarirane kemijske sestave in EDX-mikroanaliza preiskovanih realnih sestav
Type Anti-sticking additives F-additive I-additive
C
O
M
P
O
SI
T
IO
N
D
E
C
L
A
R
E
D
B
Y
T
H
E
M
A
N
U
FA
C
T
U
R
E
R
1 - Rock salt with iodine and fluorine (fine ground) NaF KIO3
2 - Rock salt with iodine, edible (fine ground) KI, KIO3
3 - Alpine salt with iodine (vacuum packed) CaCO3 min. 0.9 % KIO3MgCO3 min. 0.2 %
4 - Alpine salt with fluorine and iodine (vacuum
packed)
CaCO3 min. 0.7 % NaF KIO3MgCO3 min. 0.1 %
5 - Fine sea salt with iodine (evaporation of sea
water) KIO3
7 - NaCl, p.a. (chemically pure)
Elements observed with the EDX microprobe
E
D
X
A
N
A
L
Y
SI
S 1 - Rock salt with iodine and fluorine (fine ground) Mg, C, O F
2 - Rock salt with iodine, edible (fine ground) Ca, C, O K
3 - Alpine salt with iodine (vacuum packed) Mg, Ca, C, O K
4 - Alpine salt with fluorine and iodine (vacuum
packed) Mg, Ca, C, O F K
5 - Fine sea salt with iodine (evaporation of sea
water) Mg, Ca, Si K
7 - NaCl, p.a. (chemically pure) K
5 HYGROSCOPICITY OF CORES
Hygroscopicity of common-salt cores and chemically
pure salt cores was investigated using the weight-change
measurement. The cores were stored in various climate
conditions including the humidity of 35–58 % RV, the
temperature of 20.7–24.9 °C and normal pressure.
Hygroscopicity was observed for 23 d.
The results of the measurements given in Figure 8
show different hygroscopic behaviours of squeezed and
shot cores. The porosity and, at the same time, the
strength of the cores are considerably influenced by the
used manufacturing technology. It is evident that the
hygroscopicity increases with the growing porosity (shot
cores) while the strength is on a decrease. In the case of
a lower porosity, hygroscopicity is less evident for the
squeezed cores. But in both cases the growth of weight is
not considerable, only sufficient for short-time storing
E. ADÁMKOVÁ et al.: WATER-SOLUBLE CORES – VERIFYING DEVELOPMENT TRENDS
Materiali in tehnologije / Materials and technology 49 (2015) 1, 61–67 65
Figure 7: Regular dipyramidal form of NaCl, p. a. – standard (sample No. 7) and EDX analysis of its chemical composition
Slika 7: Dvopiramidna pravilna oblika NaCl, p. a. – standard (vzorec {t. 7) in EDX-analiza kemijske sestave
Figure 6: Oval form of sea-salt grains (sample No. 5) and EDX analysis of its chemical composition
Slika 6: Ovalna oblika zrn morske soli (vzorec {t. 5) in EDX-analiza kemijske sestave
Figure 5: Regular cubic grains of Alpine salts (samples Nos. 3, 4) and EDX analysis confirming the presence of anticaking agents on the salt
grain surface (MgCO3, CaCO3)
Slika 5: Pravilna kockasta zrna soli Alpine (vzorca {t. 3 in 4) in EDX-analiza, ki potrjuje prisotnost sredstva proti sprijemanju na povr{ini zrn soli
(MgCO3, CaCO3)
the salt cores destined for immediate use in the manu-
facturing process of castings.
6 INFLUENCE OF PRODUCTION
TECHNOLOGY ON THE SALT-CORE
POROSITY
The actual porosity, the pore size and pore distri-
bution, was measured with a Micromeritics Instrument
Corporation mercury porosimeter investigating the salt
cores, based on NaCl and squeezed with a pressure of
200 kN (104 MPa) or shot with a pressure of 7.5–8 bar.
The results are shown in Table 3.
Table 3: Properties of different salt cores
Tabela 3: Lastnosti razli~nih jeder iz soli
Sample
Total
volume of
pores
(cm3/g)
Mean
diameter
of pores
(μm)
Porosity
(%)
Bulk
density
(g/cm3)
Shot
2 0.1482 0.3240 24.20 1.6329
5 0.1443 0.4207 23.52 1.6299
7 0.2278 0.2409 19.10 0.8383
Squeezed
2 0.0320 0.2416 6.51 2.0351
5 0.0354 0.2216 6.98 1.9734
7 0.0635 0.5525 12.06 1.8977
The porosity of the squeezed cores is 6–12 % (a
correlation with the apparent-porosity calculated value)
while the porosity of the shot cores is 19–24 %, i.e.,
about four times higher. The overall volume of pores is
one order lower for the squeezed cores.
From the resulting values it is evident that the applied
production technology has a significant influence on the
core actual porosity. This is important for the penetration
of metal into the cores, especially in the case of
high-pressure die casting, where the cores are extremely
stressed by the metal pressure.
7 DISCUSSION
On the basis of the results achieved during the
investigation of commercial salts (different origins,
granulometries, grain shapes and chemical compositions)
the following can be stated:
• In all the cases the cores made with the squeezing
method achieved 2–3 times higher bending strengths
than the shot cores (Figure 4). They had a lower
hygroscopicity (lower porosity) and lower water
solubility so that only simple geometrical forms can
be manufactured with this method.
• Shot cores are highly water soluble. It is possible to
utilize the existing shooting-core foundry machines.
However, there is a negative aspect of the high hygro-
scopicity correlating with the porosity and it is, there-
fore, necessary to protect the core surfaces against
the penetration with coatings and lubricants.
• Due to the removal of the parts < 0.063 mm in the
case of all the cores, the apparent porosity increased
and the strength decreased, with the exception of
sample No. 5 – the sea salt. An explanation is offered
involving the granulometry of the bought sea salt that
contained, unlike the other salts, almost no dust parts.
The grains are highly oval, having no sharp edges.
• Dipyramid (sample No. 7 – standard) seems to be the
most suitable grain shape giving high strengths to
both the shot cores and squeezed ones. It enables the
highest coordination number of the grains and, most
of all, it is close to the shape of the SiO2 base sand.
For that reason this shape seems to be advantageous
for squeezing and shooting the cores.
• An adverse influence of the anticaking agents
(CaCO3, MgCO3, SiO2, K4[Fe(CN6)]·3H2O) in com-
bination with the regular, cubic shape of Alpine salts
resulted in a strength decrease for both methods of
the core production (squeezing and shooting). The
E. ADÁMKOVÁ et al.: WATER-SOLUBLE CORES – VERIFYING DEVELOPMENT TRENDS
66 Materiali in tehnologije / Materials and technology 49 (2015) 1, 61–67
Figure 8: Different hygroscopicity trends for squeezed and shot cores
Slika 8: Razli~na usmeritev higroskopi~nosti iztisnjenega in vbrizganega jedra
presence of undesirable additives on the grain
surfaces does not enable binding with alkali silicate
(shooting) and sticking of the grains with the
following recrystallization along the grain boundaries
(squeezing).
8 CONCLUSION
The aim of the research was to investigate the
possibility of replacing the expensive chemically pure
salt with common salt for the production of foundry
cores applied in gravity, low-pressure, and possibly also
high-pressure die casting of non-ferrous alloys.
The best strength results were achieved for crushed
salts (rock salt) used for squeezing and shooting. The
occurrence of fine-dust fractions contributes to achieving
higher core strengths. The salts with recrystallized grains
(Alpine salts) of a regular, cubic shape are less suitable,
with the exception of dipyramidal shapes (chemically
pure salts). Additions of the anti-sintering additives
(SiO2, MgCO3, CaCO3, K4[Fe(CN6)]·3H2O) (for Alpine
salts) deteriorate the strength properties.
A comparison of the core-production methods shows
that the squeezed cores have higher mechanical proper-
ties (the bending strength) and the minimum hygroscopi-
city (decreased actual porosity).
The investigation proved that chemically pure salt
can be adequately replaced with rock salt that is up to
50-times cheaper.
Acknowledgement
The research was performed with the financial sup-
port of the Technology Agency of the Czech Republic
within the Alfa programme, TA 02011314.
9 REFERENCES
1 H. Michels et al., Suitability of lost cores in rheocasting process,
Trans. Nonferrous Met. Soc. China, 20 (2010), s948–s953
2 P. Stingl, G. Schiller, Leichte und rückstandfreie Entkernung,
Giesserei Erfahrungsaustausch, 6 (2009), 4–8
3 P. Jelínek et al., Salt cores in high-pressure die-casting technology,
In: 5. Hole~kova’s conference, 1 edition, Czech Foundry Society,
Brno 2013, 63–67 (in Czech)
4 E. Adámková et al., Technology of Water Soluble Cores for Foundry
Applications, Proceedings of XX. International Student’s Day of
Metallurgy, Cracow, Poland, 2013, 8 p. (CD-ROM)
5 P. Jelínek et. al., Development of technology of salt cores manufac-
ture, Slévárenství, LXI (2013) 1–2, 28–31 (in Czech)
6 P. Jelínek et. al., Influencing the strength characteristics of salt cores
soluble in water, Slévárenství, LX (2012) 3–4, 85–89 (in Czech)
E. ADÁMKOVÁ et al.: WATER-SOLUBLE CORES – VERIFYING DEVELOPMENT TRENDS
Materiali in tehnologije / Materials and technology 49 (2015) 1, 61–67 67

E. KALINOWSKA-OZGOWICZ et al.: MATHEMATICAL MODELLING AND PHYSICAL SIMULATION ...
MATHEMATICAL MODELLING AND PHYSICAL SIMULATION
OF THE HOT PLASTIC DEFORMATION AND
RECRYSTALLIZATION OF STEEL WITH MICRO-ADDITIVES
MATEMATI^NO MODELIRANJE IN FIZIKALNA SIMULACIJA
VRO^E PLASTI^NE PREDELAVE IN REKRISTALIZACIJE JEKLA
Z MIKRODODATKI
El¿bieta Kalinowska-Ozgowicz1, Wojciech Wajda2, Wojciech Ozgowicz1
1Institute of Engineering Materials and Biomaterials, Silesian University of Technology, ul. Konarskiego 18a, 44-100 Gliwice, Poland
2Institute of Metallurgy and Materials Science, Polish Academy of Science, ul. Reymonta 25, Kraków, Poland
wojciech.ozgowicz@polsl.pl
Prejem rokopisa – received: 2013-10-07; sprejem za objavo – accepted for publication: 2014-02-06
This paper deals with the possibility of optimizing the parameters for a thermomechanical treatment of structural steel with
micro-additives, particularly in the processes of the controlled rolling of economical profiles type [240E by means of
mathematical modelling of the flow stress (p) and the physical simulation of this process based on the results of investigations
obtained by plastometric hot-torsion tests. In the description of the flow stress the rheological model suggested by C. M. Sellars
is applied in the paper, in the form p = f(,,T). Based on this model, the cause of the experimental and theoretical flow
stress-strain curves was verified, applying the minimum of the goal function in order to determine most accurately the matching
of the analysed curves of the investigated steels. It was found that the applied rheological model achieves a good matching of
the experimental and theoretical curves, whereas a physical simulation permits a complementary verification of the optimal
parameters of the controlled rolling of the manufactured products.
Keywords: flow stress, plastometric torsion test, rheological model, physical simulation, thermomechanical controlled
processing (TMCP)
^lanek obravnava mo`nost optimizacije parametrov termomehanske obdelave konstrukcijskega jekla z mikrododatki, posebno
procese kontroliranega valjanja ekonomskih profilov vrste [240E, z uporabo matemati~nega modeliranja napetosti te~enja (p)
in fizikalne simulacije tega procesa na osnovi rezultatov raziskav, dobljenih s preizkusi na vro~em torzijskem plastometru. V
tem ~lanku je bil za opis napetosti te~enja uporabljen reolo{ki model, ki ga je predlagal C. M. Sellars, v obliki p = f(,,T). Na
podlagi tega modela so bile preizku{ene eksperimentalne in teoreti~ne krivulje te~enje – raztezek z uporabo minimuma ciljne
funkcije, da bi dolo~ili najbolj{e ujemanje analiziranih krivulj preiskovanih jekel. Ugotovljeno je, da uporabljeni reolo{ki model
zagotavlja dobro ujemanje eksperimentalnih in teoreti~nih krivulj, medtem ko fizikalna simulacija dopolnjuje preverjanje
optimalnih parametrov kontroliranega valjanja proizvodov.
Klju~ne besede: napetost te~enja, plastometri~ni torzijski preizkus, reolo{ki model, fizikalna simulacija, termomehansko
kontrolirana predelava (TMCP)
1 INTRODUCTION
Mathematical modelling is an essential and econo-
mical tool in the application of new techniques for the
hot plastic deformation of metals and alloys. It makes it
possible not only to reduce the number of technical
experiments, but also to assess the interaction of the
technological parameters of the investigated processes
determining the usability of the final products. There-
fore, it integrates, in most cases, the thermal, mechanical
and microstructural processes. In the future modelling
will constitute the basis of a system for controlling most
technological processes. The constructed models will be
applied simultaneously in an off-line simulation and
on-line control of the processes1–3.
The determination of the parameters of force in
thermomechanically controlled processing (TMCP),
particularly in the controlled rolling of new profiles,
requires a knowledge of the values of the flow stress of
the steel under the specific conditions of the deforma-
tion. One of numerous methods for determining these
stresses is the hot plastometric torsion test, which also
permits a physical simulation of the sequential rolling
deformation. Thus, it becomes possible to determine the
values of the flow stress in successive roll passes, sepa-
rated by interpass times in which thermally activated
processes occur, removing the results of the strain
hardening4.
The paper presents possibilities for the mathematical
modelling of hot plastic deformation concerning some
selected products obtained by rolling structural steels
with micro-additives in the range of technological
parameters simulating such processes in the plastometric
torsion test. Special attention was paid to the verification
of the rheological model developed by C. M. Sellars et
al.5,6 describing the flow stress as a function of the
deformation and temperature, as well as strain rate and
the effect of dynamic recrystallization, which dominates
in the hot deformation for the investigated structural
steels with micro-additives of Nb, V, Ti and N.
Materiali in tehnologije / Materials and technology 49 (2015) 1, 69–74 69
UDK 519.61/.64:621.7.016.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)69(2015)
2 EXPERIMENTAL PROCEDURE
2.1 Materials
The tests were carried out on industrial melts of
structural steels with micro-additives, the chemical com-
position of which is shown in Table 1.
Plastometric tests of the steel were performed apply-
ing the hot-torsion method. The tested samples had a
diameter of 6 mm and a basis length of 10 mm. They
were cut out from slabs with dimensions of 200 mm ×
220 mm × 5300 mm provided after the hot rolling. By
means of continuous torsion, the characteristics of
plasticity for the investigated steels were determined in
the  –  system, depending on the temperature of auste-
nitization (1150–1200 °C) and deformation (800–1050
°C) as well as on the strain rate within the range 0.15 s–1
to 10 s–1. The testing was conducted on a torsion plasto-
meter that was designed at the Institute of Iron Metal-
lurgy in Gliwice, Poland. The conversion of torque-twist
to stress-strain was generally based on a Fields and
Backofen7 analysis, commonly quoted in8.
2.2 Rheological model
In constitutive equations, actually applied for the pur-
pose of modelling the processes of hot plastic working,
the effects of dynamic recovery and dynamic recrystalli-
zation are not always conveyed explicitly. The equation
in which both components are distinctly separated,
elaborated at the University of Sheffield by C. M. Sellars
et al.5,6, takes the following form:
   


= + − −
⎛
⎝
⎜
⎞
⎠
⎟
⎡
⎣⎢
⎤
⎦⎥
−0 0
1 2
1( ) exp
/
SS(e)
r
R (1)
where the respective variables are defined as follows:
R
x
=
− − −
⎛
⎝
⎜
⎞
⎠
⎟
⎡
⎣
⎢
⎤
⎦
⎥
⎧
⎨
⎩
⎫
⎬
⎭
0
1
2
( ) exp 

 SS(e) SS
C
r C
⎧
⎨
⎪
⎩
⎪
⎫
⎬
⎪
⎭
⎪
≤
>
for C
C
 
 
(1a)


0 0
1
0
1
1 0
=
⎛
⎝
⎜
⎞
⎠
⎟−sinh
Z
A
n
(1b)


SS(e) SS(e) SS(e)
SS(e)
=
⎛
⎝
⎜
⎞
⎠
⎟−
1 1
1
sinh
Z
A
n
(1c)


SS SS SS
SS
=
⎛
⎝
⎜
⎞
⎠
⎟−
1 1
1
sinh
Z
A
n
(1d)
[ ] r SS(e)= +031 1 2 2. ( )q q (1e)
 
 
x
x
r C
s C− =
−
198.
(1f)

C SS(e)
2=
⎛
⎝
⎜
⎞
⎠
⎟C
Z
Ne
0 (1g)
 
x x
Nx
C
Z
s C
SS(e)
2− =
⎛
⎝
⎜
⎞
⎠
⎟ (1h)
Z
Q
RT
=
⎛
⎝
⎜
⎞
⎠
⎟ exp def
def
(1i)
The variables in Equation (1) are as follows:
p – flow stress
0 – the maximum stress when the plastic strain  = 0
SS(e) – the onset of steady-state conditions in the extra-
polated curve
SS – the onset of steady-state conditions on the experi-
mental flow stress curve
 – plastic deformation
C – strain for the onset of dynamic recrystallization
r – the "transient strain constant" and effectively defines
the curvature of the flow-stress curve between p and
SS(e) where the equation saturates
xr – the strain required to reach a fixed amount of soften-
ing, measured in terms of S. This term effec-
tively defines the rate of softening as a result of the
dynamic recrystallization
xs – the strain at the "onset" of steady state when dyna-
mic recrystallization occurs
 – strain rate
Z – Zener – Hollomon parameter
R – the universal gas constant
Rx – a term expressing the dynamic recrystallization
Qdef – the activation energy for the deformation
A, 
, n, q, C, N – constants for each characteristic stress
p.
E. KALINOWSKA-OZGOWICZ et al.: MATHEMATICAL MODELLING AND PHYSICAL SIMULATION ...
70 Materiali in tehnologije / Materials and technology 49 (2015) 1, 69–74
Table 1: Chemical composition of the tested steels
Tabela 1: Kemijska sestava uporabljenih jekel
Steel
Concentration of the element in mass fractions, w/%
C Mn Si Cr P S Nb V Ti N Al
B1 0.15 1.03 0.25 – 0.018 0.009 0.017 0.05 – 0.0070 0.040
B2 0.16 1.25 0.32 – 0.023 0.019 0.030 0.01 – 0.0060 0.042
S9 0.38 1.52 0.63 0.56 0.027 0.012 0.030 – 0.11 – –
S1 0.66 1.02 0.26 0.30 0.029 0.021 – 0.10 – 0.0080 –
The presented model illustrates more distinctly the
behaviour of the material in the course of dynamic
recrystallization, because it makes it possible to describe
the point of deflection (peak) on the curves  –  and the
procedure of flow stress beginning at the peak strain up
to the achievement of the value of the stresses in the
steady state (SS). The coefficients of this constitutive
equation are usually determined based on the results of
plastometric tests within a wide range of temperature and
strain rates.
This allows us to model various structural materials
over a wide range of conditions for hot plastic deforma-
tion. The main difficulty in applying this model lies in
the large number of parameters, which have to be
unambiguously identified. One of the main factors
limiting the application of the mode simulating the pro-
cesses of plastic working is the difficulty in determining
these coefficients, the values of which depend on the
kind of applied material. Nevertheless, Equation (1) has
been used in this study for the investigated microalloyed
steels.
For the optimization of the parameters of the applied
model of C. M. Sellars5,6, describing stress–strain curves,
making use of the goal function  in the form:

 

=
−⎛
⎝
⎜
⎞
⎠
⎟
=
∑ 1
2
1 N
i i
ii
N
p
m C
m
p
(2)
where:
m – measured stress
c – calculated stress
Np – number of measurement points of all the curves in
the i-th experiment
For the purpose of the minimum of Equation (2), the
Nelder and Mead Simplex algorithm was used. The
calculations were performed by means of Scilab’s calcul-
ation packet.9 In results of the optimization for the appro-
priate parameters of Equation (1) were found.
3 EXPERIMENTAL RESULTS AND DISCUSSION
The results of the mathematical modelling of the
process of high-temperature plastic deformation for the
investigated micro-alloyed steels allowed us to verify
analytically the assumed Eqation (1) of the type p =
p(, , T) describing the flow stress (p) on the experi-
mental flow-stress curves, making use of the plasto-
metric method for metals and alloys with a low stack-
ing-fault energy (SFE), which in the course of hot
deformation display the phenomenon of dynamic recry-
stallization. Modelling was applied for some selected
microalloyed steels of the type HSLA with various
contents of the micro-additive Nb (steel B1 and B2), as
well as an average-carbon structural steel (0.38 % C)
with a binary system of micro-additives of Nb and Ti and
a structural rail steel containing about 0.66 % carbon
with the micro-additive V (0.10 %). The performed
numerical calculations were based on the results of a
hot-torsion test, analysing about 80 stress-strain curves
recorded with a wide range of external variables, parti-
cularly the temperature and the strain rate. A comparison
of the experimental flow curves with those in the model
concerning the investigated steels is presented in the
diagrams in Figures 1 to 4. In order to obtain a universal
description of the flow curves, the complete data of
measurements p(, , T) concerning the respective kinds
of steel were optimized. The influence of the varying
distribution of strains on the radius of the twisted
sample, as well as the changes in the temperature at the
cross-section of a massive sample have, in the process of
optimization, not been taken into account. This might
perturb the obtained results, particularly the identified
parameters of the rheological model, as has been shown
in10. The elimination of the effect of the homogeneity of
deformation and other errors in the measurements
ensures the identification of the parameters of the model,
applying the inverse method11,12. The paper10 analyses,
however, only the application of this method in the case
of various types of hot-compression tests. Future plans
will include the application of the inverse method in
dealing with the results of hot torsion tests. The numeri-
cally determined values of the coefficients assumed in
the model, taking into account the Simplex algorithm
E. KALINOWSKA-OZGOWICZ et al.: MATHEMATICAL MODELLING AND PHYSICAL SIMULATION ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 69–74 71
Figure 1: Comparison of the experimental and theoretical flow curves
of microalloyed steel B1, hot deformed in compliance with the torsion
method at a strain rate of 8.8 s–1 and at: a) Tdef = 850 °C and b) Tdef =
1000 °C
Slika 1: Primerjava eksperimentalne in teoreti~ne krivulje te~enja
vro~e deformiranega mikrolegiranega jekla B1 v primerjavi z metodo
torzije pri hitrosti deformacije 8,8 s–1 pri: a) Tdef = 850 °C in b) Tdef =
1000 °C
concerning the analysed stress-strain curves are shown in
Table 2, which also contains the values of the energy of
activation of the process of deformation and the final
values of the goal function , representing the accuracy
of the modelling solution. Due to measurement errors
and limitations with the model, different values were
obtained for the goal function (Equation (2)) concerning
the particular types of investigated steels. However, the
median square error of the calculation remains within
2–5 % limits. This proves that the assumed model fits
well with the experimental results. The flow curves of
the investigated steels indicate principally the
characteristic procedure of the continuous dynamic
recrystallization. Their analysis permits us to state that
this process is essentially influenced by the temperatures
(Figure 1) and the strain rate (Figure 2), as well as the
chemical composition of the steel (Figures 1 to 4). The
effect of these factors concerns, first of all, changes in
the value of the maximum flow stress and the stresses in
the steady state (ss), and also the values of critical
deformations (f), generally assumed to be in compliance
with the criterion of the plasticity of steel in the process
of hot deformation. It has been generally assumed that
the stresses in the flow of the investigated steels increase
with the drop in the temperature and the increasing strain
rate. It has also been found that the best compatibility in
matching the experimental and theoretical curves is
achieved in the case of steel S9 (Figure 3), with the
E. KALINOWSKA-OZGOWICZ et al.: MATHEMATICAL MODELLING AND PHYSICAL SIMULATION ...
72 Materiali in tehnologije / Materials and technology 49 (2015) 1, 69–74
Figure 4: Experimental and modelled flow curves of steel S1, hot
twisted at: a) Tdef = 900 °C,  = 1 s–1, b) Tdef = 1000 °C,  = 10 s–1
Slika 4: Eksperimentalna in modelirana krivulja te~enja jekla S1, vro~e
zvijanega pri: a) Tdef = 900 °C,  = 1 s–1, b) Tdef = 1000 °C,  = 10 s–1
Figure 3: Comparison of experimental and modelled flow curves of
steel S9 in a hot torsion test: a) Tdef = 900 °C,  = 1 s–1, b) Tdef = 1000
°C,  = 10 s–1
Slika 3: Primerjava eksperimentalne in modelirane krivulje te~enja
jekla S9 pri preizkusu vro~e torzije: a) Tdef = 900 °C,  = 1 s–1, b) Tdef
= 1000 °C,  = 10 s–1
Figure 2: Experimental and modelled flow curves of microalloyed
steel B2, hot twisted at Tdef = 900 °C and a strain rate of: a)  = 0.8 s–1
and b)  = 4.45 s–1
Slika 2: Eksperimentalna in modelirana krivulja te~enja jekla B2,
vro~e zvijanega pri Tdef = 900 °C in hitrosti deformacije: a)  = 0,8 s–1
in b)  = 4,45 s–1
values of goal function  = 0.0336, within the entire
range of the temperature and deformations.
In the case of simulative investigations, some consti-
tutive equations were applied, concerning the modelling
of the microstructure13,14, as well as the kinetics of
dynamic recrystallization, developed based on data
obtained in plastometric tests for a varying size of the
primary austenite grain. As in the rheological model, the
kinetics of dynamic recrystallization was taken into
account, and it was included in the function of the yield
stress, expressed by Equation (1). Thus, in the Sellars
model the stress depends on the size of the austenite
grain. This model predicts a more accurate flow of the
material, particularly in the processes of reiterated
deformations in the case of hot rolling.
The rolling of channel iron [240E of type B2 steel
was physically simulated based on the parameters of the
rolling mill collected in Table 3. These parameters
served as the output data for the physical simulation
realized on a torsional plastometer in a test of hot tor-
sion, applying the sequential method at a strain rate
amounting to about 4.0 s–1. The analysed flow curves
were compared with the  –  curves that were deter-
mined in the course of continuous torsion tests in order
to determine the effect of global cycles of deformation
on the initiation and progress of the activated thermal
processes occurring during the intervals between the roll
passes. The results of the kinetic investigations, includ-
ing an analysis of the share of the decay of strain
hardening during the interpass times, are presented in the
E. KALINOWSKA-OZGOWICZ et al.: MATHEMATICAL MODELLING AND PHYSICAL SIMULATION ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 69–74 73
Table 2: Optimal coefficients of the rheological model obtained as a result of Simplex optimization concerning the investigated steels with
micro-additives
Tabela 2: Optimalni koeficienti reolo{kega modela, dobljeni kot rezultati optimizacije Simplex preiskovanih jekel z mikrododatki
Steel No.
Rheology – o Activation
energy
Q/(J/mol)
Goal
function

Coefficients
Ao no 
o
B1 4.1200 · 1011 0.0525 2.2969 278323.4 0.0489
B2 1.7800 · 1013 0.1223 0.8931 330559.2 0.0399
S9 2.2300 · 108 0.0535 36.5915 297570.7 0.0336
S1 3.7700 · 1010 0.0193 21.2038 296973.7 0.0518
Steel No.
Rheology – strain hardening and dynamic recovery
Asse nsse 
sse q1 q2
B1 3.19 · 1013 4.2182 0.0028 0.1668 0.0067 · 10–2
B2 6.38 · 1019 4.3216 0.0039 0.0024 ·10–2 0.0075 · 10–2
S9 2.81 · 1013 5.3884 0.0049 0.6791 0.0026 · 10–2
S1 1.86 · 1011 2.4098 0.0203 1.2176 0.0001 · 10–2
Steel No.
Rheology – strain hardening and dynamic recrystallization
Ass nss 
ss Cc Nc Cx Nx
B1 1.80 · 109 2.4377 0.029 0.000036 0.0244 0.0234 0.279958
B2 1.22 · 1010 2.5560 0.0365 0.0433 0.0079 0.0812 0.1763
S9 5.99 · 108 3.6426 0.0279 0.0944 0.0378 0.6839 0.0957
S1 1.21 · 109 3.9124 0.0302 0.0018 · 10–2 0.0427 1.4927 0.0670
Table 3: Characteristics of the process of rolling the profile [240E
Tabela 3: Zna~ilnosti postopka valjanja profila [240E
Table of roll passes Profile [240E
Slab: 200 mm × 200 mm × 6000 mm; So = 382 cm2
Stand RollPass No.
Working pass
No. s/% h/%
Pass input
temp. T/°C Sk/cm
2 v1/(m/s) /s–1
BD 1 4 2.9 2.5 1150–1180 371.0 2.52 1.32
Z1 2 1 17.0 24.6 1140–1170 307.9 4.54 10.30
Z1 3 2 29.4 34.3 – 217.3 5.35 17.56
Z1 4 3 29.3 30.6 – 153.6 6.12 21.86
Z1 5 4 28.3 29.9 – 110.2 6.84 28.36
Z1 6 5 19.8 21.3 1050–1080 88.4 6.89 25.54
Z2 7 6 26.2 27.0 1020–1055 65.2 9.01 47.80
Z2 8 7 25.9 25.9 – 48.3 9.09 54.36
Z2 9 8 19.0 21.5 950–1020 39.1 8.83 52.60
D1 10 9 15.3 12.7 960–1050 33.9 6.97 34.57
D2 11 10 10.0 8.8 920–950 30.5 6.93 29.16
s, h indices for the deformation of the rolled channel section [240E
Sk – running cross-section of the rolled band
vl – linear velocity of the rolling
 – strain rate of the rolled band
diagrams in Figure 5. The deformations in the respective
sequential roll passes were found to be less than the
critical values required for the initiation of dynamic
recrystallization determined by continuous  –  curves,
similar to the case in tests of physical simulations,
whereas global deformations in the respective roll passes
exceeded the values cd. Thus we can assume that in the
case of the last roll passes on the flushing stand D1 and
D2 the strain hardening is accumulated. This character of
the deformation changes determines the occurrence of
both static and meta-dynamic recrystallization, and con-
sequently a grain refinement of the investigated steels.
4 CONCLUSIONS
The mathematical modelling of flow stress in the
course of high-temperature deformation allows us to
obtain the optimal parameters for the thermomechani-
cally controlled processing of the tested structural steels
with micro-additives.
Sellars’s rheological model provides good matching
of the experimental and theoretical flow curves in the
investigated steels, determined on the basis of hot pla-
stometric torsion tests.
A correct model describing the behaviour of HSLA
steel type B2 during high-temperature plastic deforma-
tion process ensured an accurate simulation of the
sequential rolling schedule for a selected economical
section of the type [240E.
A physical simulation of the rolling schedule of a
channel section [240E by the means of hot torsion tests
ensures the yield stress values p in consecutive roll
passes and calculates the force and energy parameters in
the controlled rolling process.
A kinetic analysis of thermally activated phenomena
occurring in the course of high-temperature plastic
deformation and interpass times ensures the possibility
of shaping the structure of the investigated steels and to
determine the mechanical properties of the final pro-
ducts.
5 REFERENCES
1 P. D. Hodgson, PhD Thesis, Univ. of Queensland, 1993
2 D. Martin, Mater. Sci. Technol., 10 (1994) 10, 855–861
3 R. Kuziak, Institute for Ferrous Metallurgy, Gliwice, 2005 (in Polish)
4 Z. Kuzminski, A. Nowakowski, M. Gêbala, Hutnik – Wiadomoœci
Hutnicze, 8–9 (2006), 394–399
5 S. Davenport, N. Silk, C. Sparks, C. M. Sellars, Mater. Sci. Technol.,
16 (2000), 539–546
6 B. Kowalski, C. M. Sellars, M. Pietrzyk, ISIJ Int., 40 (2000),
1230–1236
7 D. S. Fields, W. A. Backofen, 40 (2000) 1, 35
8 F. Grosman, E. Hadasik, Silesian University of Technology, Gliwice,
2005 (in Polish)
9 http://www.scilab.org/
10 D. Szeliga, M. Pietrzyk, Arch. Civ. Mech. Eng., 7 (2007) 1, 35
11 J. Lenard, M. Pietrzyk, L. Cser, Mathematical and physical simu-
lation of the properties of hot rolled products, Elsevier, Amsterdam
1999
12 D. Szyndler, PhD Thesis, AGH, Kraków, 2001 (in Polish)
13 R. Kuziak et.al., Transaction of the Institute for Ferrous Metallurgy,
1 (2012), 17–23
14 E. Kalinowska-Ozgowicz, Structural and mechanical factors of the
strengthening and recrystalization of hot plastic deformation of steels
with microadditives, Open Access Library, 20 (2013) 2, 1–246
E. KALINOWSKA-OZGOWICZ et al.: MATHEMATICAL MODELLING AND PHYSICAL SIMULATION ...
74 Materiali in tehnologije / Materials and technology 49 (2015) 1, 69–74
Figure 5: Deformation schedules for channel section type [240E; rolling simulations of microalloyed steel B2
Slika 5: Zaporedje deformacij v kanalskem delu vrste [240E; simulacija valjanja mikrolegiranega jekla B2
P. MAJERI^ et al.: SYNTHESIS OF NiTi/Ni-TiO2 COMPOSITE NANOPARTICLES VIA ULTRASONIC SPRAY PYROLYSIS
SYNTHESIS OF NiTi/Ni-TiO2 COMPOSITE NANOPARTICLES VIA
ULTRASONIC SPRAY PYROLYSIS
SINTEZA KOMPOZITNIH NANODELCEV NiTi/Ni-TiO2 Z
ULTRAZVO^NO RAZPR[ILNO PIROLIZO
Peter Majeri~1, Rebeka Rudolf1,2, Ivan An`el1, Jelena Bogovi}3, Sre~ko Stopi}3,
Bernd Friedrich3
1Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia
2Zlatarna Celje d.d., Kersnikova 19, 3000 Celje, Slovenia
3IME Process Metallurgy and Metal Recycling, RWTH University Aachen, Intzestrasse 3, Aachen, Germany
peter.majeric@um.si
Prejem rokopisa – received: 2013-10-09; sprejem za objavo – accepted for publication: 2014-02-26
In this paper we present the production of NiTi/Ni-TiO2 composite nanoparticles via the synthesis method called ultrasonic
spray pyrolysis (USP). The precursor solution for the synthesis of spherical NiTi particles was prepared from an orthodontic
wire with a chemical composition of Ni (amount fraction x = 51.46 %) and Ti (x = 48.54 %). TEM microscopy, in combination
with EDX analyses, was used for a detailed characterization of the obtained NiTi nanoparticles. The results showed the
nanoparticle sizes ranging from 60 nm to 600 nm, depending on the parameters of the production procedure. This showed the
versatility of the new USP synthesis procedure, proving its usefulness for different materials and applications.
Keywords: ultrasonic spray pyrolysis, NiTi/Ni-TiO2 composite nanoparticles, characterization, TEM microscopy
V tem prispevku je predstavljena izdelava NiTi/Ni-TiO2 kompozitnih nanodelcev po metodi ultrazvo~ne razpr{ilne pirolize
(USP). Sferi~ni nanodelci NiTi so bili izdelani iz ortodontske `ice s kemijsko sestavo Ni (mno`inski dele` x = 51,46 %) in Ti (x
= 48,54 %). Za karakterizacijo nanodelcev je bila uporabljena TEM-mikroskopija v kombinaciji z EDX-analizo. Rezultati
meritev so pokazali velikostno porazdelitev narejenih nanodelcev v razponu od 60 nm do 600 nm v odvisnosti od parametrov
izdelave. To potrjuje vsestranskost novega izdelovalnega postopka USP in njegovo uporabnost za razli~ne materiale in
aplikacije.
Klju~ne besede: ultrazvo~na razpr{ilna piroliza, kompozitni nanodelci NiTi/Ni-TiO2, karakterizacija, TEM-mikroskopija
1 INTRODUCTION
During the recent swarm of nanosized materials in
the last decade, nanosized particles of nickel titanium, a
shape-memory alloy in the macroscopic world, have not
been overlooked. The potential of the nanoparticles of
NiTi lies in their biocompatibility, magnetic and photo-
catalytic properties.1 As such, they have a potential for
applications in the anode of solid-oxide fuel cells or in
the conductive electrolytic layer of proton-exchange-
membrane fuel cells. Replacing platinum with nickel
titanium nanoparticles in automotive catalytic converters
would also significantly reduce their costs. They also
have a potential in coatings, plastics, nanowires, nano-
fibers and textiles or as a temperature-control system.2
The most widely used method for producing nickel
titanium nanoparticles is laser ablation of the alloy in a
liquid. Laser ablation has some advantages as a produc-
tion method, such as a disperse distribution of the
colloids in the solvent, no use of toxic chemical precur-
sors, chemically pure nanoparticles and its applicability
for different materials.3 The produced NiTi nanoparticles
form a titanium oxide layer on the surface, creating a
core-shell structure of the nanoparticles. The thickness of
the layer varies, depending on the organic liquids used
for the suspension.4 Other production methods of NiTi
nanoparticles include ultrasonic electrolysis, mechani-
cally assisted synthesis (nickel and titanium powders in a
planetary ball mill),5 electro-explosion of a NiTi wire
with spark-plasma sintering,6 gas-flash evaporation,7
thin-film deposition in combination with nanosphere
lithography,8 electric-discharge plasma in liquid9 and
biosynthesis/bioreduction.10
For the purposes of positioning the teeth to a desired
location in orthodontics, NiTi wires have been proved to
be very useful. This alloy’s shape-memory effect and
superelasticity provide a good combination of care for
the patient and a reasonable treatment duration.11 Even
though nickel is considered toxic, the nickel titanium
alloy is biocompatible due to the high reactivity of
titanium in oxygenated environments. As such, a tita-
nium dioxide layer is formed on the surface of a shape-
memory alloy, providing a good corrosion resistance and
preventing nickel from migrating outwards from the
alloy, causing biocompatibility of the alloy. To further
improve the handling of such orthodontic tools, a project
involving nickel titanium nanoparticles for orthodontics
was proposed. A replacement for the bulk nickel
titanium wire with a textile or polymer fiber coated with
NiTi nanoparticles via electrospinning and then remov-
ing the fiber could produce a hollow wire for orthodontic
purposes. This wire could potentially have the same
Materiali in tehnologije / Materials and technology 49 (2015) 1, 75–80 75
UDK 620.3:66.017:669.018.9 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)75(2015)
shape-memory and superelasticity properties, while
possibly reducing the material needed for the wire
production. After producing such a hollow NiTi wire, its
properties would have to be compared with the original
NiTi orthodontic wire already in use.
In our research of the nanoparticle production, we
had some successes with the production of gold nano-
particles with the ultrasonic-spray-pyrolysis method
(USP).12 This is a simple, low-cost process, available to
produce nanoparticles from different materials. The
concept of USP is to produce aerosol droplets of a solu-
tion with a dissolved material, desired for the nanoparti-
cles. These micro-sized droplets of the precursor solution
are generated by the ultrasound. The frequency of the
ultrasound and the surface tension of the solution
determine the size of the droplets. The droplets are a few
micrometers in diameter and are produced as the
ultrasound reaches the surface of the liquid, generating
cavitation (a formation, growth and implosive collapse of
the bubbles in a liquid).13,14 The generated droplets are
then carried by a carrier gas into the furnace, where the
reactions for the nanoparticle formation occur. The gas
can be a mixture of a carrier gas and a reaction gas or, in
some cases, a single gas acting as the carrier and reaction
gas at the same time. When the droplets reach the
furnace, the heightened temperature starts to evaporate
the solvent and the droplets shrink in size. The dissolved
material required for the nanoparticles starts to diffuse
into the core of the droplet. The solvent evaporates com-
pletely and the dried porous particle then reacts with the
reduction gas. The reacted particle is then sintered into
the solid final particle, which is then carried away by the
carrier gas into the collection bottles.15 This is the prin-
ciple of the one-particle-per-droplet mechanism.
The final size of the particle is thus mainly dependent
on the size of the droplets, the precursor concentration
and the reactor temperature. The droplet size is con-
trolled with the ultrasound frequency, as higher frequen-
cies produce smaller droplets. Higher precursor concen-
trations cause a higher concentration in a droplet and a
bigger final particle size.12 The temperature must be high
enough for the reactions to occur. However, very high
temperatures evaporate the solvent faster, leaving less
time for the diffusion of the material. When this occurs,
the result is the creation of porous particles and particles
with irregular shapes. In some cases the material reaction
occurs on the surface of the droplets, creating shell struc-
tures. In the case of creating ideally spherical particles,
an ideal temperature must be found, low enough to allow
the time for diffusion, while still high enough for the
required reactions to occur. This could be done experi-
mentally or theoretically, where a lot of diffusion and
reaction data would have to be acquired. A two-step USP
process can also be employed, capable of creating
core-shell structures with two different materials.13
In this work we present the results of our first expe-
rimental work in the field of NiTi/Ni-TiO2 composite-
nanoparticle production via USP. With this research we
want to simplify the field of nanotechnology, allowing a
transfer of this production to the industrial level.
2 MATERIALS AND METHODS
For the synthesis of NiTi/Ni-TiO2 composite nano-
particles, we used an orthodontic wire composed of
amount fraction of Ni x = 51.46 % and x = 48.54 % of
Ti. To prepare the precursor for USP, the wire was
dissolved in aqua regia (12 mL or 24 mL of HNO3 +
3HCl) and the resulting solution was topped off to the
required volume with water (a total volume of 500 mL).
The resulting solution was then used for generating
aerosol droplets with a USP device assembled at the IME
Process Metallurgy and Metal Recycling, RWTH
Aachen University, Germany. The device contains an
ultrasonic atomizer with an ultrasound generator and a
reservoir for the precursor, a thermostat for keeping the
precursor solution at a constant temperature, a quartz
tube and furnace for the reaction of precursor droplets,
some tubing leading the aerosol into and out of the
furnace, and collection bottles with alcohol and water for
the nanoparticle collection (Figure 1). The ultrasonic
atomizer (Gapusol 9001, RBI/France) was producing
aerosol droplets with one transducer at the selected
frequency of 2.5 MHz.
P. MAJERI^ et al.: SYNTHESIS OF NiTi/Ni-TiO2 COMPOSITE NANOPARTICLES VIA ULTRASONIC SPRAY PYROLYSIS
76 Materiali in tehnologije / Materials and technology 49 (2015) 1, 75–80
Figure 1: Schematic presentation of the USP device
Slika 1: Shematska predstavitev USP-naprave
The experiment of the USP procedure was run in
several steps. First, nitrogen was flushed into the system
at a rate of 1 L/min to remove air from the tubes. When
the furnace reached the selected temperature, hydrogen
was added to the flow of nitrogen, at a rate of 1.5 L/min.
After a while, a steady flow of nitrogen and hydrogen
was established. At this time, the ultrasonic generator
was turned on, creating aerosol droplets from the precur-
sor solution. The droplets were then carried via gas into
the furnace, where the reactions occurred, and the final
particles were then carried out of the furnace and
collected in the collection bottles. The retention time of
the droplets/particles in the quartz tube of 1 m length and
20 mm diameter was about 30 s. The temperature of the
precursor solution was kept at 23 °C and the furnace
temperature at 900 °C. The system parameters (the solu-
tion temperature, furnace temperature, steady volume of
the solution in the ultrasonic atomizer, etc.) were being
monitored for the duration of the experiment. The
collection bottles in the first experiments contained water
and, in later experiments, alcohol for the prevention of
agglomeration and titanium oxide formation. The
experiment parameters are provided in Table 1.
Table 1: Technological parameters of the performed experiments
Tabela 1: Tehnolo{ki parametri opravljenih eksperimentov
Expe-
riment
Concen-
tration
(g/L)
Tempe-
rature
(°C)
Gas
flow N2
(L/min)
Gas
flow H2
(L/min)
Time
(h)
Collec-
tion
1 0.5 900 1 1.5 5.5 Water
2 0.25 900 1 1.5 5.5 Water
3 0.5 900 1 1.5 5.0 Ethanol
4 0.25 900 1 1.5 5.0 Ethanol
A transmission electron microscope (TEM) with an
EDX analysis was used for the characterization of the
obtained nanoparticles and analysis of their morphology.
3 RESULTS
The NiTi/Ni-TiO2 composite orthodontic wire was
dissolved in aqua regia (HNO3 + 3HCl) and diluted with
water. The oxidizing properties of chloride ions and
chlorine etch away nickel and titanium from the lattice,
creating nickel chlorides NiCl2 and titanium chlorides
TiCl4. The addition of hydrogen gas inside the furnace
then reduces the chlorides and creates particles with
nickel and titanium. The resulting sizes of the produced
nanoparticles are dependent on the precursor concen-
tration and ultrasound frequency used for the creation of
aerosol droplets, as reported previously.12 The obtained
solution was left overnight for particle sedimentation.
The excess fluid was removed and the resulting samples
were characterized with TEM and EDX analysis. The
analysis showed nearly ideally spherical nanoparticles
with different sizes (Figure 2), according to the concen-
tration and the ultrasound frequency. The obtained nano-
particles consisted of titanium, nickel and oxygen (Fig-
P. MAJERI^ et al.: SYNTHESIS OF NiTi/Ni-TiO2 COMPOSITE NANOPARTICLES VIA ULTRASONIC SPRAY PYROLYSIS
Materiali in tehnologije / Materials and technology 49 (2015) 1, 75–80 77
Figure 3: Element mapping of NiTi/Ni-TiO2 composite nanoparticles
(No. 1)
Slika 3: Porazdelitev elementov pri kompozitnih nanodelcih
NiTi/Ni-TiO2 ({t. 1)
Figure 2: Sizes of NiTi/Ni-TiO2 composite nanoparticles depending
on the precursor concentration: a) 0.5 g/L NiTi, experiment numbers 1
and 3, b) 0.25 g/L NiTi, experiment numbers 2 and 4
Slika 2: Velikosti kompozitnih nanodelcev NiTi/Ni-TiO2 pri razli~nih
koncentracijah prekurzorja: a) 0,5 g/L NiTi, {tevilki poskusov 1 in 3,
b) 0,25 g/L NiTi, {tevilki poskusov 2 in 4
ure 3). The ratio between nickel and titanium was lower
than expected; with the orthodontic wire used for the
experiments it was about 1 : 1, while with the nano-
particles the values for the nickel amount were much
lower (Figure 4 and Table 2).
Table 2: Chemical composition of the points on the NiTi nanoparti-
cles’ surface form Figure 4
Tabela 2: Kemijska sestava to~k na povr{ini nanodelcev NiTi s slike 4
Point x(Ni)/% x(Ti)/% x(O)/%
1 / 44.29 54.05
2 / 41.72 56.60
3 31.22 28.21 39.36
4 40.78 20.65 38.02
5 21.78 26.88 50.70
4 DISCUSSION
The production steps for the synthesis of nano-
particles from an orthodontic NiTi wire and their usage
with the process of electrospinning are shown in Figure
5. With electrospinning the particles are spun onto fibers
like building blocks, producing a thin shell with the
properties of the nanoparticle material. The fibers could
then be removed, for example, with chemical etching,
producing new, otherwise hard-to-obtain shapes.
The high reactivity of titanium with oxygen is a
problem for pure nickel titanium nanoparticle formation.
The results suggest that the nanoparticles are formed
from titanium oxide with smaller amounts of nickel.
Even though the reaction takes place in a hydro-
gen/nitrogen atmosphere, a high amount of oxygen also
comes from the water of the precursor, as the reaction
runs at a high temperature. Some effort was made to
reduce the amount of titanium oxide formation with the
use of alcohol in the collection bottles. A smaller amount
of oxygen in alcohol allowed this reduction to occur to a
certain degree, while the total elimination of the titanium
oxide formation from the particles was not achieved. To
obtain such particles, a different solution would be
required for the dissolution of the NiTi wire. Organic
compounds and alcohol were used previously; however,
the results of such experiments were unsatisfactory. The
use of such solutions as the precursor resulted in de-
stroyed particles or no particles being collected in the
collection bottles, or alcohol complexes being formed.
The use of organic compounds would also increase the
cost and complexity of the process, while not producing
the desired product.
The analysis of the particles also showed that the
larger particles contain more nickel than the smaller
ones. This can also be seen from the TEM figures
(Figure 2), where the darker, blackened areas represent
nickel in the particles, whereas the lighter, grayer areas
represent titanium dioxide. The different shades in the
figure are the result of different densities shown by the
TEM imagery; nickel is darker because it has about
twice the density of titanium dioxide (the nickel density
is 8.9 g/cm3 versus titanium dioxide with 4.23 g/cm3).
Smaller particles, up to about 180 nm, are grayer, which
suggests there is a smaller amount of nickel, and this is
supported by the EDX analysis. In Figure 3, less nickel
is seen in smaller particles. As such, the bigger particles
(more than 180 nm) have a core of nickel and a coating
of titanium dioxide. A further analysis is required for
determining whether the core of the bigger particles is
purely nickel or the desired nickel titanium alloy.
P. MAJERI^ et al.: SYNTHESIS OF NiTi/Ni-TiO2 COMPOSITE NANOPARTICLES VIA ULTRASONIC SPRAY PYROLYSIS
78 Materiali in tehnologije / Materials and technology 49 (2015) 1, 75–80
Figure 4: Points for the EDX analysis on the NiTi/Ni-TiO2 composite
nanoparticles’ surface (No. 1)
Slika 4: Mesta na povr{ini kompozitnih nanodelcev NiTi/Ni-TiO2 ({t.
1), kjer je bila izvr{ena EDX-analiza
Figure 5: Production steps for the synthesis of nanoparticles from an orthodontic NiTi wire and their usage in electrospinning
Slika 5: Koraki pri sintezi nanodelcev iz ortodontske NiTi-`ice in njihova uporaba pri elektropredenju
It is very difficult to obtain pure nickel titanium
particles with USP due to the high reactivity of titanium
with oxygen. Even though pure nickel titanium particles
could be obtained, they would have to be stored in a
virtually oxygen-free environment to prevent a titanium
oxide formation on the surface of pure particles. As a
source for nickel titanium used in these experiments, the
orthodontic NiTi wire already has an outside layer of
titanium dioxide. In order to use this wire as a precursor
the layer would have to be removed. This would com-
plicate the process of creating the precursor to more than
just dissolving the wire in aqua regia. At this point, using
different sources for nickel and titanium, such as nickel
and titanium powders, would probably be more reason-
able. Further research into the available precursor solu-
tions for NiTi nanoparticles and collection media has to
be performed.
Even though the resulted particles were not pure
NiTi, but rather Ni-TiO2, such particles do have some
promising properties.1,16,17 They have a photocatalytic
function and can be used as a nickel titania nanocom-
posite coating, providing a good corrosion resistance and
also a heightened hardness. The nickel containing TiO2
nanoparticles is also promising as the material for high-
temperature fuel-cell electrodes, as well as the catalyst
for hydrogen cleavage in hydrogenation processes.
Nickel titania particles were synthesized with a number
of processes, such as adsorption of Ni to TiO2, sol-gel
methods, gas condensation and photodeposition of nickel
nanoparticles on titanium dioxide.1,16–19 However, with
this experiment, the USP process was proven to be a ver-
satile process, capable of creating nanoparticles of
various materials.
5 CONCLUSION
Particles of Ni-TiO2 were synthesized with USP,
using a dissolved orthodontic NiTi wire as the precursor.
The particles were formed at 900 °C, with nitrogen as
the carrier gas and hydrogen as the reduction gas and
were collected in water or alcohol. The TEM and EDX
analysis showed that spherical particles with different
compositions were created. The particles larger than 180
nm were composed of nickel with an outside layer of
titanium dioxide, while the smaller particles were mainly
composed of titanium dioxide. Further research and
analysis of the particles obtained are needed for a better
understanding of the particle formation with the experi-
ment parameters used. With the current selection of the
precursors, reaction gas and collection medium, it is
difficult to obtain pure NiTi particles, which are desired.
For this reason, an investigation of different precursor
solutions, gases and collection media should be con-
ducted.
Acknowledgement
We would like to thank the Institute for Metallurgical
Processes and Metal Recycling (IME), RWTH Aachen
University and the Institute of Metals and Technology
Ljubljana (Dr. Darja Jenko) for their support. This
research was made possible by the Slovenian Research
Agency (ARRS) as part of the applied Project L2-4212,
Research Program P2-0120 and the Young Researcher
Program.
6 REFERENCES
1 H. Chang, Y. C. Hsu, Fabrication and material properties of NiTi
nanofluid, Journal of Vacuum Science & Technology B: Micro-
electronics and Nanometer Structures, 25 (2007) 3, 935–939
2 Nickel-Titanium, Nitinol (NiTi) Nanoparticles – Properties, Appli-
cations [Online], Available: http://www.azonano.com/article.aspx?
ArticleID=3383#1 [Accessed: 23-Sep-2013]
3 S. Barcikowski, M. Hustedt, B. N. Chichkov, Nanocomposite
Manufacturing using Ultrashort-Pulsed Laser Ablation in Solvents
and Monomers, Polimery, 53 (2008) 9, 657–662
4 M. Chakif, A. Ostendorf, A. Essaidi, M. Epple, O. Prymak, Gene-
ration of Nanoparticles by Laser Ablation in Liquids: Stoichiometry/
Composition and Characterization of NiTi-Nanoparticles, The Inter-
national Conference on Shape Memory and Superelastic Technolo-
gies (SMST), Prague, 2013, 201–202
5 D. D. Radev, Mechanical synthesis of nanostructured titanium–nickel
alloys, Advanced Powder Technology, 21 (2010) 4, 477–482
6 C. Shearwood, Y. Q. Fu, L. Yu, K. A. Khor, Spark plasma sintering
of TiNi nano-powder, Scripta Materialia, 52 (2005) 6, 455–460
7 M. Kurumada, Y. Kimura, H. Suzuki, O. Kido, Y. Saito, C. Kaito,
TEM study of early Ni4Ti3 precipitation and R-phase in Ni-rich NiTi
nanoparticles, Scripta Materialia, 50 (2004) 11, 1413–1416
8 W. Crone, W. Drugan, A. Ellis, J. Perepezko, Final Technical Report:
Nanostructured Shape Memory Alloys, 841686, 2005
9 S. Yatsu, H. Takahashi, H. Sasaki, N. Sakaguchi, K. Ohkubo, T.
Muramoto, S. Watanabe, Fabrication of Nanoparticles by Electric
Discharge Plasma in Liquid, Archives of Metallurgy and Materials,
58 (2013) 2, 425–429
10 H. B. Liu, G. Canizal, P. S. Schabes-Retchkiman, J. A. Ascencio,
Structural Selection and Amorphization of Small Ni-Ti Bimetallic
Clusters, J. Phys. Chem. B, 110 (2006) 25, 12333–12339
11 J. Fer~ec, B. Gli{i}, I. [}epan, E. Markovi}, D. Stamenkovi}, I.
An`el, J. Fla{ker, R. Rudolf, Determination of Stresses and Forces on
the Orthodontic System by Using Numerical Simulation of the Finite
Elements Method, Acta Physica Polonica A, 122 (2012) 4, 659–665
12 S. Stopic, R. Rudolf, J. Bogovi}, P. Majeri~, M. ^oli}, S. Tomi}, M.
Jenko, B. Friedrich, Synthesis of Au nanoparticles prepared with
ultrasonic spray pyrolysis and hydrogen reduction, Mater. Tehnol.,
47 (2013) 5, 577–583
13 Nanostructured Materials Through Ultrasonic Spray Pyrolysis,
Sigma-Aldrich [Online], Available: http://www.sigmaaldrich.com/
technical-documents/articles/material-matters/ultrasonic-spray-pyrol
ysis.html [Accessed: 05-Sep-2013]
14 J. Bogovic, A. Schwinger, S. Stopic, J. Schroeder, V. Gaukel, H. P.
Schuchmenn, B. Friedrich, Controlled droplet size distribution in
ultrasonic spray pyrolysis, Metall, 65 (2011) 10, 455–459
15 J. Bogovi}, S. Stopi}, B. Friedrich, Nanosized metallic oxide pro-
duced by Ultrasonic Spray Pyrolysis, Proceedings of European
Metallurgical Conference, Düsseldorf/D, 2011, 1–12
16 X. Shu, J. He, D. Chen, Visible-Light-Induced Photocatalyst Based
on Nickel Titanate Nanoparticles, Ind. Eng. Chem. Res., 47 (2008)
14, 4750–4753
P. MAJERI^ et al.: SYNTHESIS OF NiTi/Ni-TiO2 COMPOSITE NANOPARTICLES VIA ULTRASONIC SPRAY PYROLYSIS
Materiali in tehnologije / Materials and technology 49 (2015) 1, 75–80 79
17 K. E. Engates, H. J. Shipley, Adsorption of Pb, Cd, Cu, Zn, and Ni to
titanium dioxide nanoparticles: effect of particle size, solid con-
centration, and exhaustion, Environ. Sci. Pollut. Res., 18 (2011) 3,
386–395
18 F. A. Sheikh, J. Macossay, M. A. Kanjwal, A. Abdal-hay, M. A. Tan-
try, H. Kim, Titanium Dioxide Nanofibers and Microparticles
Containing Nickel Nanoparticles, ISRN Nanomaterials, 2012 (2012),
article ID 816474, 1–8
19 J. L. Rodríguez, F. Pola, M. A. Valenzuela, T. Poznyak, Photocata-
lytic Deposition of Nickel Nanoparticles on Titanium Dioxide,
Materials Research Society Symposium Proceedings, 1279 (2010),
19–25
P. MAJERI^ et al.: SYNTHESIS OF NiTi/Ni-TiO2 COMPOSITE NANOPARTICLES VIA ULTRASONIC SPRAY PYROLYSIS
80 Materiali in tehnologije / Materials and technology 49 (2015) 1, 75–80
R. RUDOLF et al.: HYDROXYAPATITE COATINGS ON Cp-TITANIUM GRADE-2 SURFACES ...
HYDROXYAPATITE COATINGS ON Cp-TITANIUM GRADE-2
SURFACES PREPARED WITH PLASMA SPRAYING
NANOS HIDROKSIAPATITA NA POVR[INO Cp-TITANA GRADE-2
Z NABRIZGAVANJEM S PLAZMO
Rebeka Rudolf1, Dragoslav Stamenkovi}2, Zoran Aleksi}2, Monika Jenko3,
Igor \ordevi}2, Aleksandar Todorovi}2, Vukoman Jokanovi}4, Karlo T. Rai}5
1University of Maribor, Faculty of Mechanical Engineering, Maribor, Slovenia
2University of Belgrade, School of Dental Medicine, Belgrade, Serbia
3Institute of Metals and Technology, Ljubljana, Slovenia
4University of Belgrade, VINCA Institute of Nuclear Sciences, Belgrade, Serbia
5University of Belgrade, Faculty of Technology and Metallurgy, Belgrade, Serbia
karlo@tmf.bg.ac.rs
Prejem rokopisa – received: 2013-10-09; sprejem za objavo – accepted for publication: 2014-03-24
Thin hydroxyapatite coatings were produced on Cp-Titanium Grade-2 samples, with new high-voltage pulse-power equipment
PJ-100 (Plasma Jet, Serbia) in order to get a more stable implant structure appropriate for further clinical applications. A com-
parative analysis of differently prepared surfaces of the Cp-Titanium Grade-2 samples was done before the hydroxyapatite was
applied.
Microstructural observation of the modified hydroxyapatite/implant surface was done using scanning-electron-microscopy
imaging and Auger electron spectroscopy, with the aim of detecting the morphology and the elements contained in the new
surfaces of the samples. The results confirmed that the surface of Cp-Titanium Grade-2 modified with hydroxyapatite is very
similar to the bone structure.
Keywords: Cp-Ti2 material, hydroxyapatite (HA), plasma-spray coating, characterization
Z novo, visokonapetostno pulzirajo~o plazemsko napravo PJ-100 (plasma Jet, Serbia) je bil pripravljen tanek nanos hidroksi-
apatita na vzorce iz Cp-Titan Grade-2, da bi dobili bolj stabilno strukturo vsadka za klini~no uporabo. Izvr{ena je bila pri-
merjalna analiza razli~no pripravljenih povr{in vzorcev iz Cp-Titana Grade-2, preden je bil izvr{en nanos hidroksiapatita.
Opazovanje mikrostrukture s hidroksiapatitom modificirane povr{ine vsadka je bilo izvr{eno z vrsti~no elektronsko mikro-
skopijo, kot tudi z Augerjevo elektronsko spektroskopijo, z namenom odkritja morfologije in razporeditve elementov na novi
povr{ini vzorca. Rezultati so potrdili, da je modificirana povr{ina Cp-Titana Grade-2 s hidroksiapatitom zelo podobna strukturi
kosti.
Klju~ne besede: material Cp-Ti2, hidroksiapatit (HA), s plazmo nabrizgan nanos, karakterizacija
1 INTRODUCTION
Titanium (Ti) has two most useful metal properties:
corrosion resistance and the highest strength-to-weight
ratio compared to other metals. Moreover, when in the
unalloyed condition, it is stronger than some other simi-
lar materials, such as contemporary steel, and it is more
than 40 % lighter. Ti can be alloyed with various ele-
ments (Al, V, Nb, Zr) that can additionally improve its
strength, while maintaining high temperature perfor-
mance, creep resistance, weldability, response to ageing
heat treatments and formability. Ti and its alloys are
resistant to corrosion because of the formation of an
insoluble and continuous titanium-oxide layer on the sur-
face having one of the highest heats of reaction: H =
–912 kJ/mol. In air, an oxide (usually TiO2) begins to
form within nanoseconds (10–9 s) and reaches a thickness
of 2–10 nm already in 1 s. This oxide is very adherent to
the parent Ti, protecting the metal from other impurities
and it is impenetrable to oxygen1. Moreover, this oxide
on the Ti-surface gives Ti its excellent biocompatibility2.
Commercially pure (Cp) Ti is available in four grades
(1–4) that vary with respect to oxygen (w = 0.18–0.40 %)
and iron (w = 0.20–0.50 %) amounts. These slight con-
centration differences have a substantial effect on the
physical properties of the metal. Oxygen, in particular,
has a great influence on the ductility and strength of Ti3.
In our investigation we used Cp-Titanium Grade-2
(Cp-Ti2) materials which have the medium value of O.
Cp-Ti2 is used widely because it combines excellent for-
mability and moderate strength with superior corrosion
resistance. This combination of properties makes Cp-Ti2
a good candidate for medical applications, i.e., implant
manufacturing.
These materials are classified as biologically inert
biomaterials. This means that they remain essentially un-
changed when implanted into human bodies4–6. The
human body is able to recognise these materials as
foreign and tries to isolate them by encasing them in
fibrous tissues. However, they do not cause any adverse
reactions and are tolerated well by the human tissues7.
Although the bulk properties of biomaterials are import-
ant with respect to their biointegration, the biological
responses of the surrounding tissues to dental implants
are controlled mostly with their surface characteristics
Materiali in tehnologije / Materials and technology 49 (2015) 1, 81–86 81
UDK 621.793/.795:669.295 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)81(2015)
(chemistry and structure) because biorecognition takes
place at the interface of an implant and the host tissue.
Biological tissues mainly interact with the outermost
atomic layers of an implant, which are about 0.1–1 nm
thick. The molecular and cellular events at the bone-
implant interface are well described8, but many crucial
aspects are still far from being understood. A justifica-
tion of the surface modification of implants is, therefore,
straightforward: to retain the key physical properties of
an implant while modifying only the outermost surface
to control the bio-interaction. As a result, a lot of re-
search work is devoted to elaborate methods of modi-
fying the surfaces of the existing implants (biomaterials)
to achieve the desired biological responses.
With respect to the presented problems of the state-
of-the-art implant surface, the main question of our
research was how to prepare the ideal micro-topography
of the selected Cp-Ti2 material. In this framework we
took into consideration that the most common physico-
chemical treatments are chemical surface reactions (e.g.,
oxidation, acid-etching), sand blasting, ion implantation,
laser ablation, coating the surface with inorganic hydro-
xyapatite, etc. These methods alter the energy, charge
and composition of the existing surface, but they can also
provide the surfaces with modified roughness and mor-
phology9.
Although more accurate knowledge is required,
typical Cp-Ti2 surfaces were treated with different high-
voltage-plasma process parameters. In this manner we
wanted to test the experimental presentation of the thin
hydroxyapatite (HA) coating on the Cp-Ti2 samples
using a specific plasma installation10. The aim of this
work was, namely, to produce a thin HA coating with the
new methods in order to get a more stable surface struc-
ture of Cp-Ti2 that could be used for further clinical
applications. Finally, a comparative analysis was done
using the classical preparation method, chemical etching,
in order to acquire the information about the new plasma
technology.
2 MATERIALS AND METHODS
The samples for the experiments were made from the
Cp-Ti2 material (ASTM F 67-100) with a chemical com-
position in mass fractions as follows: w(C)max = 0.08 %,
w(N)max = 0.030 %, w(O)max = 0.25 %, w(Fe)max = 0.30 %,
w(H) = 0.015 %. The properties of Cp-Ti2 are: the
melting point of 1668 °C, a Vickers hardness of 180 HV,
a 0.2 % yield strength of 275–450 MPa, a tensile
strength of > 350 MPa, an elongation of > 20 %, a den-
sity of 4.5 g/cm3, the coefficient of thermal expansion of
25–400 °C 9.4 × 10–6/K, and elasticity modulus of 110
GPa. The samples were produced in the shape of discs
with the dimension of  = 10 mm and a thickness of
2 mm.
It is well known that the adhesion between a ceramic
coating (HA) and a metal implant is strongly influenced
by the metal surface roughness. Therefore, the process of
roughening is adjusted to the optimum level with several
parameters: pressure of water, air flow, corundum granu-
lation and diameter of the nozzle for roughening. The
chamber for roughening was constructed in-house from
hardened steel. The tool for capturing the samples (also
our own construction) was used as a separate set of tools,
rotating at more than 1000 r/min. The roughening of the
samples was done under the following conditions: a fluid
pressure of 8 bar, a nozzle diameter of  10 mm and a
corundum granulation of 1–2 mm. The obtained surface
roughness (Ra) was in the range from 4.72 μm to 5.28 μm
(a Perthen Perthometer). After the roughening, the sam-
ples were washed with toluene and acetone to remove
the traces of oil from the compressor. The specimens
were then immersed in 5.0 mL of 5 M NaOH aqueous
solution for 24 h. Afterwards they were removed from
the solution and washed gently with distilled water,
followed by drying at 40 °C for 24 h in an air atmo-
sphere. The treated metal was then heated up to 300 °C
(Sample F1), 500 °C (Sample F2), 700 °C (Sample F3)
and 800 °C (Sample F4) at a rate of 5 °C/min in an Ni–Cr
electrical furnace in an air atmosphere over 5 h. After the
thermal treatment all the samples were cooled to room
temperature in the furnace. An untreated sample (Sample
F0) was used for obtaining a better insight into the diffe-
rences between the treated and untreated samples.
The plasma installation PJ-100 (Plasma Jet, Serbia)
was used for the plasma-spray process10,11. The basic
parameters of the installation used for applying HA
depositions were: a power of (52.0 ±1.5) kW, a voltage
of (120±2) V, a current of (430±5) A, an argon flow of
(38.5±1.2) L/min, powder carrier gas, air of 8 L/min and
a powder-mass input of (2.0±0.1) g/s. The samples were
placed on a drum (ø = 200 mm) of the plasma instal-
lation (a rotation velocity of 3.7 r/s) for the HA deposi-
tion. The complete coating process was performed in 2–3
short intervals taking 7 s to 10 s each, with a break rate
of a few minutes between the cycles of deposition. Be-
fore the plasma treatment Ti disks were heated at a
temperature of 200 °C, which was found to be the
optimum for obtaining the maximum partition of the HA
crystalline phase in the process of its deposition. Two
commercially available HA powders were used: HA
powder XPT-D-703 (Sulzer Metco, USA) and HA pow-
der Captal 90 (Plasma Biotal, UK). The granulations of
the powders were obtained from the manufacturers’ spe-
cifications.
A microstructural characterization of the samples was
carried out with scanning electron microscopy (SEM-
Sirion 400 NC), in addition to an energy-dispersive
X-ray (EDX) analysis (Oxford INCA 350). The speci-
mens were observed directly without any surface prepa-
ration.
Detailed microstructure observations of the obtained
surfaces were done using field-emission Auger electron
spectroscopy (FE-AES) – a Thermo Scientific Microlab
R. RUDOLF et al.: HYDROXYAPATITE COATINGS ON Cp-TITANIUM GRADE-2 SURFACES ...
82 Materiali in tehnologije / Materials and technology 49 (2015) 1, 81–86
310-F spectrometer equipped with a spherical-sector
analyser and a field-emission electron gun with a ther-
mally assisted Schottky field-emission source providing
a stable electron beam in an accelerating voltage range of
(0.5–25) keV.
3 RESULTS AND DISCUSSION
3.1 Microstructure observation of Cp-Ti2 surfaces
Typical microstructures of the Cp-Ti2 surfaces after
different processes of roughening are shown in Figure 1.
From the pictures it can be concluded that immersion
in NaOH represents a technique which results in obtain-
ing a high surface roughness (Figure 1a). On the other
hand, the thermal treatment at different temperatures
represents the attainable method for surface roughness.
At lower temperatures the roughness is lower (Figures
1b, 1c, 1d) and, consequently, at higher temperatures the
roughness is higher (Figure 1e).
This is very important for the quality of the HA coat-
ing textures on dental Cp-Ti2 implants. The ideal micro-
topography of a Cp-Ti surface is still unknown because it
is very difficult to associate Cp-Ti/HA interface proper-
ties with clinical results12.
3.2 Microstructure observation of HA/Cp-Ti2 surfaces
Using FE-AES spectroscopy we discovered that the
cross-sections of the HA coatings for samples F0-F4
obtained with plasma-jet deposition clearly show their
depths, as shown in Table 1.
Table 1: HA coating depth
Tabela 1: Debelina HA-nanosa
Sample
label
Average
depth,
nm
Maximum
depth,
nm
Minimum
depth,
nm
Maximum
deviation from the
average value, nm
F0 206 221.7 192.9 15.7
F1 205.8 216.2 191.4 14.4
F2 152.4 172.2 138.5 19.8
F3 156.2 167.4 140.7 15.4
F4 156.1 164.3 146.5 9.7
According to the data given in Table 1 there are no
significant differences in the depth of these coatings. For
the same sample (the same conditions of their plasma
treatment) the variety of the depths along the entire
section was in the range of ± 20 nm. These values show
a very uniform depth of the coating for each sample,
with the maximum depth deviation (from its average
value) between (9.7 and 19.8) nm.
R. RUDOLF et al.: HYDROXYAPATITE COATINGS ON Cp-TITANIUM GRADE-2 SURFACES ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 81–86 83
Figure 2: SEM micrographs of HA coating on different substrates: a)
F0, b) F1, c) F2, d) F3 and e) F4
Slika 2: SEM-posnetki HA nanosa na razli~nih podlagah: a) F0, b) F1,
c) F2, d) F3 in e) F4
Figure 1: SEM micrographs of Cp-Ti2 surfaces: a) F0, b) F1, c) F2,
d) F3 and e) F4
Slika 1: SEM-posnetki povr{ine Cp-Ti2: a) F0, b) F1, c) F2, d) F3 in
e) F4
The HA coating of sample F0 (Figure 2a) shows a
typical, very rough surface, with a very interesting mor-
phology. The smallest sizes of the mosaic details were
only a few μm, while very large and mutually inter-
connected parts were dominant. The obtained form is
typical for quickly cooled systems that were exposed to
melting at the edges of various particles. The specific
form of the surface topology seems to be very promising
for the cell adhesion. The structure is layered showing
strongly exposed hills and valleys. In some places very
deep voids can be observed. The mosaic details with an
irregular form of very large agglomerates show the
dimensions of about 50 μm.
The HA coating of sample F1 (Figure 2b) shows a
fairly homogenous surface with almost parallel patterns
that look like textile fibres. The motifs are more discreet.
The intertwined details are very ordered, with slightly
raised or lowered parts. The details have the size of
about 1 μm or less.
The HA coating of sample F2 (Figure 2c) has a very
rough surface with almost parallel patterns and it is
sprinkled with white and black motifs resembling
furrows, with the corresponding typical morphology.
Very light details show individual particle sizes between
(0.7 and 2.9) μm.
The HA coating of sample F3 (Figure 2d) shows a
specific surface morphology with a worm-like structure
and finer, more expressed grains than in the previous
cases. Individual particles are clearly visible. Their
values are between 0.3 μm and 0.6 μm. The mosaic is
very uniform. The grains are elongated and of a prisma-
tic form. There are no visible voids and it seems that the
layers of HA were sorted properly, without a build-up
causing the creation of valleys and hills, as shown for
samples F0-F2.
The HA coating of sample F4 (Figure 2e) shows a
very specific layout of cavities and elevations. It looks
like a very nicely knitted network with several visible
layers and holes with the dimensions of 7–36 μm. The
depths of the continuous, interconnected parts of the net-
work are of fairly similar dimensions. These morpholo-
gies show the first layer below the perforated layers (the
ropes in network) with a fairly flat surface and small
parallel patterns. The sizes of individual particles are
between 0.25 μm and 1.45 μm.
The EDS analysis of the HA coating (Figure 3)
shows all the elements present inside the coating from
the titanium substrate to the top of the coating. The part
of the substrate changed during the plasma-jet deposition
can be also seen on the corresponding picture.
The concentrations of various elements given in the
above diagrams show very similar data: the boundaries
between the Ti substrate and HA coatings are clearly
visible. In all the cases, small quantities of oxygen and
aluminum are noticed, while the concentration of Ti
deceases strongly, reaching a zero value. On the left side
R. RUDOLF et al.: HYDROXYAPATITE COATINGS ON Cp-TITANIUM GRADE-2 SURFACES ...
84 Materiali in tehnologije / Materials and technology 49 (2015) 1, 81–86
Figure 3: EDS analysis of the cross-sections of various samples of HA coatings: a) F0, b) F1, c) F2, d) F3 and e) F4 depending on the distance
from the left or right side of the boundary
Slika 3: EDS-analiza prereza razli~nih vzorcev HA-nanosa: a) F0, b) F1, c) F2, d) F3 in e) F4, v odvisnosti od razdalje od meje levo ali desno
of the boundary, small quantities of calcium and phos-
phorous were registered. Also, in this area, at a distance
of 3.6 μm, the highest level of Al was noticed for sample
F0 (w = 3.95 %), while for F1, F2, F3 and F4 the highest
levels of Al were noticed at the following boundaries: w
=12.8 % for F1, w = 13.7 % for F2, w = 16 % for F3 and w
= 17 % for F4. The oxygen level was very low on the left
side of the boundary for the untreated sample F0, while
for sample F1 it was between w = 3.9 % at a distance of
3.6 μm left of the substrate boundary and w = 35 % at the
boundary; for sample F2 it was w = 1.1 % at a distance of
3.6 μm left of the boundary and w = 13.7 % at the boun-
dary; for sample F3 it was w = 1.9 % at a distance of 1.8
μm left of the boundary and w = 16 % at the boundary;
and for F4 it was w = 20.45 % at a distance of 3.6 μm left
of the boundary and w = 41.5 % at the boundary. The
boundary was narrower or wider, and the mixing of ions
from the right and left sides of the boundary was clearly
visible at a distance of 1.8 μm left of the boundary (for
samples F2, F3 and F0). For samples F1 and F4 the
boundary was narrower and the decrease in Ti was very
sharp. The maximum depth of penetration of Ca and P
ions in all the cases was about 3.8 μm inside the sub-
strate depth. These values show good agreement for all
the samples. Inside the HA coatings, at a distance of
1.8 μm on the right, there was no more interference, i.e.,
the mixing of the ions from the Ti substrate with the ions
contained inside the HA coating. In deeper layers of HA
coatings the concentrations of Ca, P and O ions reached
R. RUDOLF et al.: HYDROXYAPATITE COATINGS ON Cp-TITANIUM GRADE-2 SURFACES ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 81–86 85
Figure 4: a) SEM image of the interface, line-scanned with AES, bet-
ween the substrate and HA layer for sample F1, b) AES line scan of
the interface between HA and the substrate indicating the concentra-
tions of Ca, O, Ti and P
Slika 4: a) SEM-posnetek stika, kjer je bila izvr{ena AES linijska
analiza med podlago in HA-nanosom pri vzorcu F1, b) AES linijska
analiza stika HA in podlage, ki prikazuje koncentracijo elementov Ca,
O, Ti in P
Figure 5: a) SEM image with marked places of AES, b) AE spectra
for the spots in a)
Slika 5: a) SEM-posnetek to~k AES-analize, b) AES-spektri v to~kah
iz a)
Table 2: Measured elements using AES, amount-of-substance, x/%
Tabela 2: Merjenje elementov z AES, mno`inski dele`, x/%
Concentration, x/%
Spot C Ca O Ti
P1 11.2 0.0 35.2 53.6
P2 12.2 6.5 60.4 20.9
P3 46.3 25.6 24.3 3.8
P4 17.6 3.6 31.0 47.8
saturation, becoming uniform with a further increase in
the distance.
A typical secondary electron image (SEI) of the
interface, line-scanned with Auger electron spectroscopy
(AES), between the substrate and HA coatings for
sample F1, and a spectrogram of the distributions of Ca,
O, Ti and P are given in Figure 4.
This figure clearly shows a decrease in the Ti amount
(present until the depth of the coating reaches a value of
20 nm) and an increase in the Ca in the coating between
4 nm and 17.5 nm. Both of these curves are typical S
curves, inversely relative to one another. The oxygen
amount decreases slightly, showing its highest value
inside the chemically and thermally treated part of tita-
nium (between 0 nm and 4 nm). In the area of the HA
coating its amount reduces slightly by 4 nm and 10 nm
(the part in which Ca atoms are propagated due to a high
velocity of the calcium-hydroxide particles melted on the
surfaces by the plasma jet). Afterwards, the oxygen
amount varies statistically within certain limits.
Similar results are shown in Figure 5, where the con-
centrations of various elements, measured in counts per
second (CPS), are explained with respect to the kinetic
energy of the back-scattering electrons. Particular values
of these energies prove the presence of Ca, Ti and O,
showing their changes between sites P1–P4 (the sites at
various distances from the boundary with hydroxyapa-
tite). Besides, these values are given in Table 2 for each
of the observed sites.
Finally, as reported by Kokubo13 and Kim14, the
calcium-phosphate deposition obeys the nucleation pro-
cess. The electric-charge interaction also heavily favours
this process. A higher surface energy favours the calcium
phosphate initiating nucleation and the basic TiOH with
the negative charge can adsorb the positively charged
calcium ion (Ca2+) first adsorbed to neutralize the surface
charge until too much charge accumulates producing a
positive layer of calcium titanate. Then, this positive
layer interacts further with the negatively charged
phosphate ion (PO43–) to form the amorphous calcium
phosphate, which will be crystallized further into an
OCP phase despite the fraction of the amorphous
calcium-phosphate residue.
4 CONCLUSIONS
The innovative plasma-jet process of hydroxyapatite
deposition with a previous preparation of titanium sub-
strates with NaOH etching and a subsequent thermal
treatment is shown in this paper. In addition, the as-depo-
sited coatings were composed mainly of the HA phase
with the crystallite sizes between 15.5 nm and 31 nm.
Furthermore, Auger electron spectroscopy, through
Ca-ion implantation in the depth of the oxide area of the
substrate, shows the processing advantages that, conse-
quently, led to the increased adhesion strength. There-
fore, this improved plasma-jet method, with an unusually
high kinetic energy of the plasma, seems to be promising
for the fabrication of nanostructured HA coatings with
unique microstructural features, which are desirable for
enhancing the biological properties of HA coatings.
Research confirms a high potential of a deposition of
HA on titanium substrates via a new plasma installation
for any medical purpose.
Acknowledgements
This work was supported by the Ministry of Educa-
tion, Science and Technological Development of the
Republic of Serbia (Project 172026) and EUREKA Pro-
gramme CELL-TI E!5831.
5 REFERENCES
1 E. P. Lautenschlager, P. Monaghan, Titanium and titanium alloys as
dental materials, Int Dent J, 43 (1993) 3, 245–53
2 B. D. Ratner, A. S. Hoffman, F. J. Schoen, J. E. Lemons, Bioma-
terials science: an introduction to materials in medicine, Academic
Press, New York 1996
3 H. M. Kim et al., Thin film of low-crystalline calcium phosphate
apatite formed at low temperature, Biomaterials, 21 (2000),
1129–1134
4 L. Le Guehennec, A. Soueidan, P. Layrolle, Y. Amouriq, Surface
treatments of titanium dental implants for rapid osseointegration,
Dental Materials, 23 (2007), 844–854
5 G. Daculsi, O. Laboux, O. Malard, P. Weiss, Current state of the art
of biphasic calcium phosphate bioceramics, Journal of Materials
Science: Materials in Medicine, 14 (2003), 195–200
6 Z. T. Sul et al., The bone response of oxidized bioactive and non-
bioactive titanium implants, Biomaterials, 24 (2003), 3893–3907
7 K. Turzo, Surface Aspects of Titanium Dental Implants, In: R. H.
Sammour (ed.), Biotechnology – Molecular Studies and Novel
Applications for Improved Quality of Human Life, InTech, Rijeka
2012, 237 (www.intechopen.com)
8 D. A. Puleo, A. Nanci, Understanding and controlling the bone-im-
plant interface, Biomaterials, 20 (1999), 2311–2321
9 B. Kasemo, Biological surface science, Surface Science, 500 (2002),
656–677
10 M. Vilotijevi}, Dc plasma arc generator with increasing volt-ampere
characteristic, RS Patent No 49, 7062007, 2007
11 J. Ru`ic et al., Understanding plasma spraying process and charac-
teristics of DC-arc plasma gun (PJ-100), Metallurgical & Materials
Engineering, 18 (2012), 273–282
12 I. Milinkovi} et al., Aspects of titanium-implant surface modification
at the micro and nano levels, Mater. Tehnol., 46 (2012) 3, 251–256
13 T. Kokubo, H. M. Kim, M. Kawashita, Novel bioactive materials
with different mechanical properties, Biomaterials, 24 (2003) 13,
2161–2175
14 H. M. Kim et al., Surface potential change in bioactive titanium
metal during the process of apatite formation in simulated body fluid,
Journal of Biomedical Materials Research, Part A, 67A (2003) 4,
1305–1309
R. RUDOLF et al.: HYDROXYAPATITE COATINGS ON Cp-TITANIUM GRADE-2 SURFACES ...
86 Materiali in tehnologije / Materials and technology 49 (2015) 1, 81–86
P. TERNIK et al.: HEAT-TRANSFER CHARACTERISTICS OF A NON-NEWTONIAN Au NANOFLUID ...
HEAT-TRANSFER CHARACTERISTICS OF A NON-NEWTONIAN
Au NANOFLUID IN A CUBICAL ENCLOSURE WITH
DIFFERENTIALLY HEATED SIDE WALLS
ZNA^ILNOSTI PRENOSA TOPLOTE NENEWTONSKE Au
NANOTEKO^INE V KOCKASTEM OHI[JU Z RAZLI^NO
GRETIMA STRANSKIMA STENAMA
Primo` Ternik1, Rebeka Rudolf2,3, Zoran @uni~4
1Private Researcher, Bresterni{ka ulica 163, 2354 Bresternica, Slovenia
2University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia
3Zlatarna Celje, d. d., Kersnikova ul.19, 3000 Celje, Slovenia
4AVL-AST, Trg Leona [tuklja 5, 2000 Maribor, Slovenia
pternik@pt-rtd.eu
Prejem rokopisa – received: 2013-10-14; sprejem za objavo – accepted for publication: 2014-02-17
The present work deals with the laminar natural convection in a cubical cavity filled with a homogenous aqueous solution of
carboxymethyl cellulose (CMC) based gold (Au) nanofluid obeying the power-law rheological model. The cavity is heated on
the vertical and cooled from the adjacent wall, while the other walls are adiabatic. The governing differential equations were
solved with the standard finite-volume method and the hydrodynamic and thermal fields are coupled using the Boussinesq
approximation.
The main objective of this study is to investigate the influence of the nanoparticle volume fraction on the heat-transfer
characteristics of CMC-based Au nanofluid over a wide range of the base-fluid Rayleigh number.
Accurate numerical results are presented in the form of dimensionless temperature and velocity variations, the mean Nusselt
number and the heat-transfer rate. It is shown that adding nanoparticles to the base fluid delays the onset of natural convection.
In addition, numerical simulations showed that, just after the onset of natural convection, adding nanoparticles reduces the mean
Nusselt number value for any given base-fluid Rayleigh number.
Keywords: natural convection, CMC-Au nanofluid, heat transfer, Nusselt number
Prispevek obravnava naravno konvekcijo v kockastem ohi{ju, napolnjenem s homogeno nanoteko~ino karboksimetil celuloza
(CMC)-zlato (Au), z reolo{kim vedenjem, opisanim s poten~nim zakonom. Ohi{je je greto na navpi~ni in hlajeno na prile`ni
steni, medtem ko so druge stene adiabatne. Vodilne diferencialne ena~be so bile re{ene s standardno metodo kon~nih prostornin,
pri ~emer sta hidrodinami~no in temperaturno polje sklopljena z Boussinesqovo aproksimacijo.
Glavni cilj prispevka je raziskati vpliv prostorninskega dele`a nanodelcev na zna~ilnosti prenosa toplote CMC-Au nanoteko~ine
za {iroko obmo~je vrednosti Rayleighjevega {tevila nosilne teko~ine.
Natan~ni rezultati so predstavljeni v obliki spreminjanja brezdimenzijske temperature in hitrosti, srednje vrednosti Nusseltovega
{tevila in hitrosti prenosa toplote. Pokazano je, da dodajanje nanodelcev v nosilno teko~ino zakasni za~etek naravne konvekcije.
Poleg tega so numeri~ne simulacije pokazale, da takoj za pojavom naravne konvekcije dodajanje nanodelcev zmanj{uje srednjo
vrednost Nusseltovega {tevila za katero koli vrednost Rayleighjevega {tevila nosilne teko~ine.
Klju~ne besede: naravna konvekcija, nanoteko~ina CMC-Au, prenos toplote, Nusseltovo {tevilo
1 INTRODUCTION
In recent years, nanosized particles dispersed in a
base fluid, known as nanofluid1, have been used and
researched extensively to enhance the heat transfer in
many engineering applications. While the presence of
nanoparticles shows an unquestionable heat-transfer
enhancement in forced-convection applications2 there is
still a dispute on the influence of nanoparticles on the
heat-transfer enhancement of a buoyancy-driven flow.
Natural convection (i.e., a flow caused by tempera-
ture-induced density variations) is one of the most exten-
sively analysed configurations because of its fundamen-
tal importance as the "benchmark" problem for studying
convection effects (and comparing as well as validating
numerical techniques). In addition to the obvious
academic interest, a thermally driven flow is the strategy
preferred by heat-transfer designers when a small power
consumption, a negligible operating noise and a high
reliability of a system are the main concerns (e.g., a re-
duction of the cooling time, manufacturing cost and an
improvement of the product quality in the injection-
moulding industry3). Although various configurations of
the enclosure problem are possible4–8, one of the most
studied cases (involving nanofluids) is a two-dimensio-
nal square enclosure with differentially heated iso-
thermal vertical walls and adiabatic horizontal walls9–11.
Recent numerical studies12,13 illustrated that the sus-
pended nanoparticles substantially increase the heat-
transfer rate for any given Rayleigh number. In addition,
it is shown that the heat-transfer rate in water-based
nanofluids increases with an increasing volume fraction
of Al2O3, Cu, TiO2 or Au nanoparticles. On the other
hand, an apparently paradoxical behaviour of the heat-
transfer deterioration was observed in many experimen-
Materiali in tehnologije / Materials and technology 49 (2015) 1, 87–93 87
UDK 519.61/.64:536.2:620.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)87(2015)
tal studies: e.g., Putra et al.14 reported that the presence
of Al2O3 nanoparticles in a base fluid reduces the
natural-convective heat transfer. However, they did not
clearly explain why the natural-convective heat transfer
decreases with an increase in the volume fraction of
nanoparticles. This was later explained by the work of
Ternik and Rudolf4.
Regarding the natural convection in a 3D cubic
cavity, most (if not all) of the work was done for the
Newtonian fluid. For example, Tric et al.15 studied the
natural convection in a 3D cubic enclosure using a
pseudo-spectra Chebyshev algorithm based on the pro-
jection-diffusion method with the spatial resolution
supplied by polynomial expansions. Lo et al.16 also
studied the same problem under different inclination
angles using the differential-quadrature method to solve
the velocity-vorticity formulation of the Navier-Stokes
equation employing higher-order polynomials to appro-
ximate differential operators.
The above review of the existing literature shows that
the problem of natural convection in a cubical cavity
filled with a nanofluid is an issue still far from being
tackled and solved. Framed in this general background,
the purpose of the present study is to examine the effect
of adding Au nanoparticles to a non-Newtonian base
fluid on the conduction and convection heat-transfer cha-
racteristics in a differentially heated cubical cavity
heated over a wide range of the base-fluid Rayleigh
number (101  Rabf  106) and the volume fraction of
nanoparticles (0 %    10 %).
2 NUMERICAL MODELLING
The standard finite-volume method is used to solve
the coupled conservation equations of mass, momentum
and energy. In this framework, a second-order central
differencing scheme is used for the diffusive terms and a
second-order upwind scheme for the convective terms.
The coupling of the pressure and velocity is achieved
using the SIMPLE algorithm. The convergence criteria
were set to 10–8 for all the relative (scaled) residuals.
2.1 Governing equations
For the present study, a steady-state flow of an
incompressible non-Newtonian CMC-Au nanofluid is
considered. It is assumed that both the fluid phase and
nanoparticles are in thermal and chemical equilibrium.
Except for the density, the properties of the nanoparticles
and the fluid (presented in Table 1) are taken to be con-
stant. The Boussinesq approximation is invoked for the
nanofluid properties to relate the density changes to the
temperature changes, and to couple the temperature field
with the velocity field.
The governing equations (mass, momentum and
energy conservation) of such a flow are9:
∂
∂
v
x
i
i
= 0 (1)
( ) 
 

nf nf
nf
v
v
x x
v
x
p
x
g
j
i
j j
i
j
i
∂
∂
∂
∂
∂
∂
∂
∂
−
⎛
⎝
⎜
⎞
⎠
⎟ =
= − +

( ( ( )T T
x
v
xC j
j
i
− +
⎛
⎝
⎜
⎞
⎠
⎟) 
∂
∂
∂
∂

nf
(2)
 c v
x x
k
T
xj j j j
p nf nf
∂
∂
∂
∂
∂
∂
Τ
=
⎛
⎝
⎜
⎞
⎠
⎟ (3)
where the cold-wall temperature TC is taken to be the
reference temperature for evaluating the buoyancy term
()nfg(T – TC) in the momentum-conservation equa-
tion.
In the momentum-balance law (Equation 2) the con-
stitutive (i.e., rheological) equation is required for the
viscous function ( )
 nf  , which, for the power-law
model, reads as:
( )
  nf  = −K n 1 (4)
where   = ∑∑1
2 ij jiji
is the II invariant of the symmetrical rate-of-deforma-
tion tensor with Cartesian components  ij = (vi/xj) +
(vj/xi), K is the consistency and n is the dimensionless
power-law index with experimentally determined values
(for a w = 0.4 % aqueous CMC solution) of K = 0.048
Pasn and n = 0.882.
The relationships between the properties of the
nanofluid (nf), the base fluid (bf) and pure solid (s) are
given with the following empirical models4,8–13:
• Dynamic viscosity:

 
 nf bf= −/ ( )
.1 2 5
• Density:
   nf bf s= − +( )1
• Thermal expansion:
   nf bf s= − +( )(1
• Heat capacitance:
      c c cp nf p bf p s= − +( )(1
• Thermal conductivity:
k k
k k k k
k k k knf bf
s bf bf s
s bf bf s
=
+ − −
+ + −
2 2
2


( )
( )
Table 1: Thermophysical properties of CMC-Au nanofluid9
Tabela 1: Toplotno-fizikalne lastnosti CMC-Au nanoteko~ine9
 (kg/m3) Cp (J/kg K) k (W/m K)  (1/K)
0.4 % CMC 997.1 4179 0.613 2.1 × 10–4
Au 19320 128.8 314.4 1.416 ×10–7
2.2 Geometry and boundary conditions
The simulation domain is shown schematically in
Figure 1. The two horizontal walls of the square enclo-
P. TERNIK et al.: HEAT-TRANSFER CHARACTERISTICS OF A NON-NEWTONIAN Au NANOFLUID ...
88 Materiali in tehnologije / Materials and technology 49 (2015) 1, 87–93
sure are kept at different constant temperatures (TH > TC),
whereas the other boundaries are considered to be
adiabatic. All the velocity components (i.e., vx, vy and vz)
are identically zero on each boundary because of the
no-slip condition and the impenetrability of the rigid
boundaries.
To study the heat-transfer characteristics due to the
natural convection in nanofluids, the local Nusselt num-
ber (along the vertical hot wall) is defined as15:
Nu y z
Q
Q
h L
k
L
T T
T y z
x
x
( , )
( , )
= = = −
−
nf, conv
nf, cond
nf
nf H C
∂
∂ =0
(5)
where Nu(y,z) presents the ratio of the heat-transfer rate
by convection to that by the conduction in the nanofluid
in question, knf is the thermal conductivity and hnf is the
convection heat-transfer coefficient of the nanofluid:
h k
T y z
x T T
x
nf nf
H C
= −
−=
∂
∂
( , )
0
1
(6)
Finally, in the present study, the heat-transfer charac-
teristics are analysed in terms of the mean Nusselt
number:
Nu
L
Nu y z y z
LL
= ∫∫
1
2
00
( , )d d (7)
and the ratio of the nanofluid heat-transfer rate to the
base-fluid one:
Q
Q
k Nu
k Nu
nf
bf
nf nf
bf bf
= (8)
at the same value of the base-fluid Rayleigh number.
In order to investigate the influence of solid-particle
volume fraction  on the heat-transfer characteristics, the
Rayleigh number (Ranf) and Prandtl number (Prnf) of the
CMC-based nanofluid (obeying the power-law viscous
behaviour) are expressed as follows:
Ra
c g TL
k K
K k
c
nf
nf p nf
n
nf
n
nf
n
nf
nf p
=
=
⎛
⎝
⎜
⎞
⎠
+   
 
( Δ 2 1
Pr ⎟ −
nf
n 2
n
–
L2 2
(9)
Using Equation 9 we show that Ranf < Rabf (Figure
2a) and Prnf < Prbf (Figure 2b) for all the values of .
The ratio of the CMC-based nanofluid Rayleigh and
Prandtl numbers to the base-fluid Rayleigh and Prandtl
numbers decreases with the increasing volume fraction
of the nanoparticles; i.e., for the fixed values of Rabf and
Prbf, the value of Ranf and Prnf decrease when adding
nanoparticles.
2.3 Grid refinement and numerical accuracy
The grid independence of the present results was
established with a detailed analysis using three different
non-uniform meshes (the elements were concentrated
towards each solid wall), the details of which are pre-
sented in Table 2. The table includes the numbers of the
elements in a particular direction as well as the nor-
malized minimum cell size.
P. TERNIK et al.: HEAT-TRANSFER CHARACTERISTICS OF A NON-NEWTONIAN Au NANOFLUID ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 87–93 89
Figure 2: Variation of the dimensionless numbers of the CMC-based
nanofluids with the volume fraction of nanoparticles: a) Rayleigh
number and b) Prandtl number
Slika 2: Spreminjanje brezdimenzijskih {tevil CMC nanoteko~in v
odvisnosti prostorninskega dele`a nanodelcev: a) Rayleighjevo {tevilo
in b) Prandtlovo {tevilo
Figure 1: Schematic diagram of the simulation domain
Slika 1: Shematski prikaz obmo~ja simulacije
With each grid refinement the number of the ele-
ments (i.e., the control volumes) in a particular direction
is increased and the element size is reduced. Such a pro-
cedure is useful for applying Richardson’s extrapolation
technique (a method for obtaining a higher-order estima-
te of the flow value from a series of lower-order discrete
values) encountered in many numerical studies8–12,17.
Table 2: Computational-mesh characteristics
Tabela 2: Zna~ilnosti ra~unskih mre`
Mesh I Mesh II Mesh III
Nx × Ny × Nz 40 × 40 × 40 60 × 60 × 60 90 × 90 × 90
min/L 2.250 × 10–3 1.500 × 10–3 1.000 × 10–3
For a general primitive variable  the grid-converged
value (i.e., extrapolated to the zero element size) is given
as8–12: ext = M3 – (M2 – M3)/(rp – 1) where M3 is
obtained on the basis of the finest grid and M2 is the
solution based on the next level of the coarse grid; r =
1.5 indicates the ratio between the coarse- and fine-grid
spacings and p = 2 indicates the order of accuracy.
The numerical error e = ( /  M2 ext ext− for the mean
Nusselt number Nu is presented in Table 3. It can be
seen that the differences in the grid refinement are
exceedingly small and the agreement between Mesh II
and the extrapolated value is extremely good (the
discretisation error is below 0.20 %). Based on this
estimation, the simulations in the remainder of the paper
were conducted on Mesh II that provided a reasonable
compromise between high accuracy and computational
effort.
Table 3: Effect of the mesh refinement upon the mean Nusselt number
( = 10 %, Rabf = 106)
Tabela 3: Vpliv zgo{~evanja mre`e na srednjo vrednost Nusseltovega
{tevila ( = 10 %, Rabf = 106)
Mesh I Mesh II Mesh III Nuext e
10.075 10.089 10.095 10.099 0.109 %
2.4 Benchmark comparison
In addition to the aforementioned grid-dependency
study, the present results were also checked against the
results of other authors15,16 for the natural convection of
air in a cubic enclosure (Pr = 0.71). The comparisons
between the present simulation results and the corres-
ponding benchmark values (Table 4) are very good and
entirely consistent with our grid-dependency studies.
Table 4: Comparison of the present results for Nu with the benchmark
results
Tabela 4: Primerjava dobljenih rezultatov za Nu z referen~nimi
rezultati
Ra = 103 Ra = 104 Ra = 105 Ra = 106
Present study 1070 2.048 4.327 8.627
Tric et al.15 1.070 2.054 4.337 8.641
Lo et al.16 1.071 2.054 4.333 8.666
3 RESULTS AND DISCUSSION
3.1 Temperature and velocity-flow field
It is useful to inspect the variations in dimensionless
temperature  and dimensionless vertical velocity v z
* in
order to understand the influences of Rabf and  on the
heat-transfer characteristics during the natural convec-
tion of the CMC-Au nanofluid in the cubical enclosure.
It is evident from Figure 3a that the distributions of 
become increasingly non-linear with the increasing
values of Rabf. This statement is supported by the data
plotted in Figure 3b, which demonstrates that the magni-
tude of the velocity component increases significantly
with the increasing Rabf. For a given set of values of  an
increase in Rabf gives rise to a strengthening of the
buoyancy forces in comparison to the viscous forces,
which can be seen from Figure 3b, where the magnitude
of v z
* increases with the increasing Rabf. As the convec-
tive transport strengthens with the increasing Rabf the
distribution of  becomes significantly more non-linear
with the increasing Rabf (Figure 3a).
Furthermore, Figure 3a illustrates that the thermal-
boundary-layer thickness next to the heated wall is
influenced by the addition of nanoparticles to the base
fluid. This sensitivity of the thermal-boundary-layer
P. TERNIK et al.: HEAT-TRANSFER CHARACTERISTICS OF A NON-NEWTONIAN Au NANOFLUID ...
90 Materiali in tehnologije / Materials and technology 49 (2015) 1, 87–93
Figure 3: a) Variations in the non-dimensional temperature and
b) non-dimensional vertical velocity component
Slika 3: a) Spreminjanje brezdimenzijske temperature in b) brez-
dimenzijske navpi~ne komponente hitrosti
thickness to the volume fraction of the nanoparticles is
related to the increased thermal conductivity of the
CMC-based Au nanofluid. As a matter of fact, the in-
creased values of the thermal conductivity are accom-
panied by higher values of the thermal diffusivity and
these higher values of the thermal diffusivity result in a
reduction in the temperature gradients and, accordingly,
an increase in the thermal-boundary thickness as
demonstrated in Figure 3a. Finally, this increase in the
thermal-boundary-layer thickness reduces the Nusselt-
number values.
3.2 Mean Nusselt number
Figure 4 presents the variation in the mean Nusselt
number with the Rayleigh number. For smaller values of
the base-fluid and nanofluid Rayleigh numbers there is
no convection in the nanofluid or the base fluid, and the
heat transfer occurs due to pure conduction; the mean
Nusselt number equals 1.0 and its value is independent
of the Rayleigh number. As the value of Rabf increases,
the nanofluid remains in the conductive-heat-transfer
regime, while the convective-heat-transfer regime
appears in the base fluid.
The onset of the heat convection4 (i.e., the critical
value of the base-fluid Rayleigh number, at which the
mean Nusselt number equals 1.01) depends on the vo-
lume fraction of the Au nanoparticles. The higher is the
value of , the more delayed is the occurrence of the
convective-heat-transfer regime (Figure 5a). When the
nanofluid is in the convective-heat-transfer regime, the
mean Nusselt number monotonically increases with the
Rayleigh number (Figures 4 and 5), attaining lower
values for the higher nanoparticle volume fraction (Fig-
ure 4) at a given base-fluid Rayleigh number.
On the other hand, it is interesting to notice that the
onset of the convection occurs at the same critical value
of the nanofluid Rayleigh number, i.e., Ranf,cr ≅ 281
(Figure 5b). Moreover, the value of the mean Nusselt
number at a given nanofluid Rayleigh number is prac-
tically independent of the nanoparticle volume fraction
(Figure 4b), which is consistent with the earlier findings
in the context of the natural convection of generalized
Newtonian fluids18 as well as Au nanofluids4 in a side
wall, differentially heated, 2D square enclosure.
This finding is a reflection of the nanofluid Prandtl-
number values considered in the present study and for
this range of Prandtl-number values (i.e., Prnf > 1) the
P. TERNIK et al.: HEAT-TRANSFER CHARACTERISTICS OF A NON-NEWTONIAN Au NANOFLUID ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 87–93 91
Figure 5: Onset of the heat-transfer convection as a function of the:
a) base-fluid Rayleigh number and b) nanofluid Rayleigh number
Slika 5: Nastop konvektivnega prenosa toplote v odvisnosti od:
a) Rayleighjevega {tevila nosilne teko~ine in b) Rayleighjevega {tevila
nanoteko~ine
Figure 4: Variation in the mean Nusselt number with the: a) base-fluid
Rayleigh number and b) nanofluid Rayleigh number
Slika 4: Spreminjanje srednjega Nusseltovega {tevila v odvisnosti z:
a) Rayleighjevim {tevilom nosilne teko~ine in b) Rayleighjevim {tevi-
lom nanoteko~ine
hydrodynamic-boundary-layer thickness remains much
greater than the thermal-boundary-layer thickness, and
thus the transport characteristics are driven primarily by
the buoyancy and viscous forces, which is reflected in
the weak Prandtl-number dependence on the mean
Nusselt number.
3.3 Heat-transfer rate
Figure 6 shows the effect of the base-fluid Rayleigh
number Rabf on the ratio of the heat-transfer rates for the
CMC-based Au nanofluid for different values of the vo-
lume fraction.
In the range of Rabf  281 the heat transfer occurs
due to pure conduction and the mean Nusselt number
equals Nu = 1. Consequently, the ratio of the heat
transfer is equal to the ratio of thermal conductivities and
it is constant and independent of the base-fluid Rayleigh
number. For Ranf  281 and Rabf > 281 the nanofluid
remains in the conductive-heat-transfer regime, while the
convection appears in the base fluid. The heat transfer is
more important in the base fluid than in the nanofluid
and the ratio of the heat-transfer rates is on a decrease.
After the onset of the convective-heat-transfer regime
in the nanofluid, the ratio of the heat-transfer rates is on
an increase but remains lower than the ratio obtained
when both the nanofluid and the base fluid are in the
conductive regime.
Last but not least, the present conclusions can be
extrapolated to the other CMC-based nanofluids (e.g.,
Cu, TiO2 and Al2O3) since their nanofluid Rayleigh- and
Prandtl-number values are within the range of the
present (i.e., CMC-based Au nanoparticles) Rayleigh-
and Prandtl-number values.
4 CONCLUSIONS
In the present study, a steady laminar natural convec-
tion of a non-Newtonian (i.e., CMC-based Au) nanofluid
obeying the power-law rheological model in a cubical
enclosure with differentially heated side walls, subjected
to constant wall temperatures was analysed with nume-
rical means. The effects of the base-fluid Rayleigh
number (101  Rabf  106) and the solid volume fraction
(0 %    10 %) on the momentum and heat-transfer
characteristics were systematically investigated in detail.
The influence of the computational grid refinement
on the present numerical predictions was studied
throughout the examination of the grid convergence at
Rabf = 106 and  = 10 %. By utilizing extremely fine
meshes, the resulting discretisation error for Nu is well
below 0.2 %.
The present numerical method was validated for the
case of the natural convection of air in a cubical cavity,
for which the results of other authors are available in the
available literature. Remarkable agreement of the present
results with the benchmark results yields sufficient con-
fidence in the presented numerical procedure and results.
Highly accurate numerical results allow some im-
portant conclusions such as:
Just after the onset of the convection, there is more
heat transfer in the base fluid than in the nanofluid. At a
fixed value of the base-fluid Rayleigh number Rabf, the
nanofluid Rayleigh number Ranf decreases with the vo-
lume fraction of the nanoparticles. Thus, the nano-
particles delay the onset of the convection.
In the convective-heat-transfer regime the mean
Nusselt number Nu is found to increase with the increas-
ing values of the base-fluid Rayleigh number Rabf, but
the Nu values obtained for the higher values of the nano-
particle volume fraction  are smaller than those ob-
tained in the case of the base fluid ( = 0 %) at the same
nominal values of Rabf.
The transition from the conductive to the convective
heat-transfer regime occurs at the same value of the
nanofluid Rayleigh number, i.e., Ranf  281.
The values of the mean Nusselt number at a given
Ranf are practically independent of the nanoparticle vo-
lume fraction.
The heat-transfer rate can decrease or increase
depending on the value of the Rayleigh number. So, an
addition of nanoparticles increases the heat transfer only
for the given values of the temperature difference.
Acknowledgements
The research leading to these results was carried out
within the framework of research project "Production
technology of Au nano-particles" (L2-4212) and it re-
ceived the funding from the Slovenian Research Agency
(ARRS).
P. TERNIK et al.: HEAT-TRANSFER CHARACTERISTICS OF A NON-NEWTONIAN Au NANOFLUID ...
92 Materiali in tehnologije / Materials and technology 49 (2015) 1, 87–93
Figure 6: Heat-transfer rate variation with the base-fluid Rayleigh
number and nanoparticle volume fraction
Slika 6: Spreminjanje razmerja prenosa toplote z Rayleighjevim
{tevilom nosilne teko~ine in prostorninskim dele`em nanodelcev
5 REFERENCES
1 S. U. S. Choi, Enhancing thermal conductivity of fluids with
nanoparticles, Developments Applications of Non-Newtonian Flows,
66 (1995), 99–105
2 W. Daungthongsuk, S. Wongwises, A critical review of convective
heat transfer in nanofluids, Renewable & Sustainable Energy
Reviews, 11 (2009), 797–817
3 F. H. Hsu, K. Wang, C. T. Huang, R. Y. Chang, Investigation on con-
formal cooling system design in injection molding, Advances in
Production Engineering & Management, 8 (2013), 107–115
4 P. Ternik, R. Rudolf, Conduction and convection heat transfer cha-
racteristics of water-based Au nanofluids in a square cavity with
differentially heated side walls subjected to constant temperatures,
Thermal Science, 18 (2014), 189–200
5 G. Huelsz, R. Rechtman, Heat transfer due to natural convection in
an inclined square cavity using the lattice Boltzmann equation
method, International Journal of Thermal Sciences, 65 (2013),
111–119
6 M. N. Hasan, S. C. Saha, Y. T. Gu, Unsteady natural convection
within a differentially heated enclosure of sinusoidal corrugated side
walls, International Journal of Heat and Mass Transfer, 55 (2012),
5696–5708
7 G. Yesiloz, O. Aydin, Laminar natural convection in right-angled
triangular enclosures heated and cooled on adjacent walls, Interna-
tional Journal of Heat and Mass Transfer, 60 (2013), 365–374
8 P. Ternik, R. Rudolf, Z. @uni~, Numerical study of Rayleigh-Benard
natural convection heat transfer characteristics of water-based Au
nanofluids, Mater. Tehnol., 47 (2013) 2, 211–215
9 P. Ternik, R. Rudolf, Laminar natural convection of non-Newtonian
nanofluids in a square enclosure with differentially heated side walls,
International Journal of Simulation Modelling, 12 (2013), 5–16
10 O. Turan, R. J. Poole, N. Chakraborty, Influences of boundary
conditions on laminar natural convection in rectangular enclosures
with differentially heated side walls, International Journal of Heat
and Fluid Flow, 33 (2012), 131–146
11 P. Ternik, R. Rudolf, Heat transfer enhancement for natural convec-
tion flow of water-based nanofluids in a square enclosure, Inter-
national Journal of Simulation Modelling, 11 (2012), 29–39
12 P. Ternik, R. Rudolf, Z. @uni~, Numerical study of heat transfer
enhancement of homogeneous water-Au nanofluid under natural
convection, Mater. Tehnol., 46 (2012) 3, 257–261
13 H. F. Oztop, E. Abu-Nada, Y. Varol, K. Al-Salem, Computational
analysis of non-isothermal temperature distribution on natural con-
vection in nanofluid filled enclosures, Superlattices and Microstruc-
tures, 49 (2011), 453–467
14 N. Putra, W. Roetzel, S. K. Das, Natural convection of nano-fluids,
Heat and Mass Transfer, 39 (2002), 775–784
15 E. Tric, G. Labrosse, M. Betrouni, A first incursion into the 3D
structure of natural convection of air in a differentially heated cubic
cavity, from accurate numerical simulations, International Journal of
Heat and Mass Transfer, 43 (2000), 4034–4056
16 D. C. Lo, D. L. Young, K. Murugesan, GDQ method for natural con-
vection in a cubic cavity using velocity-vorticity formulation,
Numerical Heat Transfer, Part B, 48 (2005), 363–386
17 I. Bilu{, M. Morgut, E. Nobile, Simulation of Sheet and Cloud
Cavitation with Homogenous Transport Models, International
Journal of Simulation Modelling, 13 (2013), 94–106
18 O. Turan, A. Sachdeva, N. Chakraborty, R. J. Poole, Laminar natural
convection of power-law fluids in a square enclosure with diffe-
rentially heated side walls subjected to constant temperatures,
Journal of Non-Newtonian Fluid Mechanics, 166 (2011), 1049–1063
P. TERNIK et al.: HEAT-TRANSFER CHARACTERISTICS OF A NON-NEWTONIAN Au NANOFLUID ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 87–93 93

Z. LA[OVÁ, R. ZEM^ÍK: AMPLITUDE–FREQUENCY RESPONSE OF AN ALUMINIUM CANTILEVER BEAM ...
AMPLITUDE–FREQUENCY RESPONSE OF AN ALUMINIUM
CANTILEVER BEAM DETERMINED WITH PIEZOELECTRIC
TRANSDUCERS
AMPLITUDNO-FREKVEN^NI ODZIV KONZOLNEGA NOSILCA IZ
ALUMINIJA, UGOTOVLJEN S PIEZOELEKTRI^NIMI
PRETVORNIKI
Zuzana La{ová1, Robert Zem~ík2
1University of West Bohemia in Pilsen, Department of Mechanics, Univerzitní 22, 306 14, Plzeò, Czech Republic
2European Centre of Excellence NTIS – New Technologies for Information Society, Faculty of Applied Sciences, University of West Bohemia,
Univerzitní 22, 306 14, Plzeò, Czech Republic
zlasova@kme.zcu.cz
Prejem rokopisa – received: 2013-10-16; sprejem za objavo – accepted for publication: 2014-02-13
This work is focused on the creation of an appropriate finite-element model of an aluminum cantilever beam using a pair of
piezoelectric patch transducers. Thanks to the reversible behavior of the piezoelectric effect each patch transducer can represent
either an actuator or a sensor. For a precise prediction of the amplitude values in the numerical simulations each transducer is
calibrated before being attached to the beam with strain gauges. From these experiments piezoelectric properties of each
piezoelectric patch are obtained. The cantilever beam is actuated with a voltage signal applied to one of the patches. The signal
is a linear chirp (sine wave with a swept frequency) with a sufficient range to affect the selected natural frequencies. The time
response of the beam from the piezoelectric sensor and, alternatively, from the laser position sensor is transformed with the
STFT algorithm to obtain the characteristics of the time-frequency domain (spectrogram). The finite-element model of the
cantilever beam with the piezoelectric patches was created using 3D solid structural and piezoelectric bricks in Ansys. The time
response of the model to the chirp voltage signal was determined with a transient analysis. The amplitude/frequency characte-
ristics are compared with the experimental results.
Keywords: piezoelectric materials, frequency spectrum, finite-element analysis
To delo obravnava izdelavo primernega modela z metodo kon~nih elementov konzolnega nosilca iz aluminija z uporabo para
piezoelektri~nih pretvornikov v obliki obli`a. Zaradi reverzibilnega vedenja piezoelektri~nega pojava je lahko vsak obli`ast
pretvornik aktuator ali senzor. Za natan~no napovedovanje vrednosti amplitude pri numeri~nih simulacijah je bil vsak pretvornik
pred namestitvijo na nosilec kalibriran z napetostnimi listi~i. Iz teh preizkusov so dobljene piezoelektri~ne lastnosti vsakega
piezoelektri~nega obli`a. Konzolni nosilec je bil aktiviran s signalom elektri~ne napetosti, uporabljene na enem od obli`ev.
Signal je linearno cvr~anje (sinus s {ablonirano frekvenco) s primernim obmo~jem, da se vpliva na izbrane naravne frekvence.
^asovni odziv nosilca iz piezoelektri~nega senzorja in alternativno s polo`aja laserskega senzorja je pretvorjen s STFT-algo-
ritmom, da se dobi zna~ilnosti vedenja ~as – frekvenca (spektrogram). Izdelan je bil model s kon~nimi elementi konzolnega
nosilca s piezoelektri~nimi obli`i. Z uporabo 3D strukturnih in piezoelektri~nih opek v Ansys in z uporabo kon~nih elementov
je bil izdelan model konzolnega nosilca s piezoelektri~nimi obli`i. ^asovni odziv modela na cvr~e~ signal napetosti je bil
dolo~en s prehodno analizo. Zna~ilnosti amplitude in frekvence so primerjane z eksperimentalnimi rezultati.
Klju~ne besede: piezoelektri~ni material, spekter frekvenc, analiza kon~nih elementov
1 INTRODUCTION
Automatic detection of impacts and hidden defects in
structures (denoted as structural health monitoring,
SHM) is the present-day trend in the non-destructive
testing. The data obtained from the sensors applied to a
structure are transmitted to the control system, which
evaluates the state of the structure and responses appro-
priately (e.g, it enables the warning system). Especially
the structures made of composite materials affected by
hidden defects such as fibre debonding or delamination1
call for an integration of the novel SHM methods.
Two approaches are usually distinguished in SHM:
the passive and active ones. In a passive system "natural"
impulses like impacts or crack propagation create stress
waves that are sensed by a grid of sensors. The source
can be then localized and reconstructed.2 In an active
SHM system the health of a structure is assessed by eva-
luating its response to specific actuating signals. One
way of establishing damage is by detecting the changes
in a structure’s modal properties (particularly the basic
natural frequency),3 providing the global information on
its state. Another way, denoted as the pitch-catch tech-
nique, uses the scattering of stress waves when the
actuating signal approaches a structural defect.4
The sensors and actuators in SHM systems are made
of smart materials that have a capability to convert vari-
ous kinds of energy, e.g., piezoelectric materials (Ro-
chelle salt, tourmaline or artificially produced ceramics)
respond to a mechanical deformation by generating elec-
tric voltage and vice versa. Therefore, piezoelectric
materials can be used as both sensors and actuators.
This work is focused on creating a reliable FE model
of a piezoelectric transducer used in SHM. The model is
then tested in the case when two patches are applied to a
Materiali in tehnologije / Materials and technology 49 (2015) 1, 95–98 95
UDK 519.61/.64:620.179.1:681.586.773 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)95(2015)
cantilever beam. An appropriate FE model of a piezo-
electric patch is the key part for designing a future SHM
system.
2 MODEL OF THE PIEZOELECTRIC
MATERIAL
The piezoelectric material can be described with a set
of constitutive equations:



D
C e
e E
⎡
⎣⎢
⎤
⎦⎥
=
−⎡
⎣⎢
⎤
⎦⎥
⋅
⎡
⎣⎢
⎤
⎦⎥
T
(1)
where  [Pa] is stress vector 6 × 1, C [Pa] is the matrix
of elastic coefficients 6 × 6,  is strain vector 6 × 1, D
[C/m2] is electric displacement vector 3 × 1, E [V/m] is
electric-field intensity vector 3 × 1, μ is dielectric matrix
3 × 3 (with electric permittivity constants on its diago-
nal) and e [C/m2] is piezoelectric stress matrix 3 × 6:
e
e
e
e e e
=
⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥
0 0 0 0 0
0 0 0 0 0
0 0 0
15
24
31 32 33
(2)
The rows of matrix e denote the direction of the elec-
tric field and the columns refer to the strain components,
e.g., constant e32 of the piezoelectric actuator quantifies
strain 2 induced by the electric field in transversal
direction 3.
In the producer’s datasheet piezoelectric strain matrix
d is listed instead of e. These matrices are related using
the stiffness matrix:
e CdT T= (3)
A transversally isotropic material model is consi-
dered for the piezoelectric ceramic with the main direc-
tion of anisotropy identical with the direction of polarity.
Stiffness matrix C is obtained with the inversion of
compliance matrix S:
S
E
v
E
v
E
v
E E
v
E
v
E
v
E=
− −
− −
− −
1
0 0 0
1
0 0 0
1
21
2
31
3
12
1 2
32
3
13
1
23
2
1
0 0 0
0 0 0
1
0 0
0 0 0 0
1
0
0 0 0 0 0
1
3
23
13
12
E
G
G
G
⎡
⎣
⎢
⎢
⎢
⎢
⎢
⎢
⎢
⎢
⎢
⎢
⎢
⎢
⎤
⎦
⎥
⎥
⎥
⎥
⎥
⎥
⎥
⎥
⎥
⎥
⎥
⎥
(4)
where E1 is the Young’s modulus in the direction per-
pendicular to the plane of isotropy, while E2 = E3 are the
Young’s moduli in the plane of isotropy. G12 = G13 are
the shear moduli in the planes perpendicular to the plane
of isotropy, G23 = G32 are the shear moduli in the plane
of isotropy and are defined with:
G
E
v
E
v23
2
23
3
322 1 2 1
=
+
=
+( ) ( )
(5)
Poisson’s ratios v12 = v13 give a measure of the exten-
sion (compression) in the plane of isotropy due to the
extension (compression) in the direction perpendicular to
this plane and vice versa. v23 = v32 are the Poisson’s ratios
in the plane of isotropy.
3 DETERMINATION OF PIEZOELECTRIC
COEFFICIENTS OF THE PATCHES
The piezoelectric patches used in the experiments are
DuraAct P-876.A12 with a layer of piezoelectric ceramic
(type PIC-255 5) with the dimensions of 50 mm × 30 mm
× 0.2 mm. The top and bottom surfaces of the ceramic
are silvered and connected to the soldering pads. Due to
the high fragility the ceramic is embedded in a protective
polymeric foil with the dimensions of 61 mm × 35 mm ×
0.5 mm. The mechanical and electrical properties of the
constituent materials are presented in Tables 1 and 2.
Material properties of the piezoelectric ceramic were
determined by the producer.5
Table 1: Material properties
Tabela 1: Lastnosti materiala
Units PZT Foil Al
Young’s modulus E [GPa] See
Tab.2
3 68
Poisson’s ratio v [ – ] 0.3 0.3
Density  [kg/m3] 7800 1580 2777
Relative electric
permittivity
μ1/μ0* [ – ] 1650
– –μ2/μ0 [ – ] 1650
μ3/μ0 [ – ] 1750
Piezoelectric
coefficients
e31,e32 [C/m] 6.4 – –
e33 [C/m] –20.5 – –
* μ0 = 8.85418 × 10–12 F/m is vacuum permittivity
Table 2: Elastic parameters of piezoelectric ceramic PIC-255
Tabela 2: Parametri elasti~nosti piezoelektri~ne keramike PIC-255
E1 [GPa] 62.1
E2 = E3 [GPa] 48.3
v12 = v13 [ – ] 0.34
v23 = v32 [ – ] 0.34
G12 = G13 [GPa] 23.1
G23 = G32 [GPa] 17.9
Although geometric and material properties of the
piezoelectric patches are supposed to be similar for one
type and production set, various results with differences
up to 20 % were obtained using two patches of the same
type and set. For this reason the piezoelectric properties
of the patches used in following experiment need to be
determined.
In the first experiment two pairs of strain gauges
(HBM rosettes 6/650 RY91) were glued on the two
Z. LA[OVÁ, R. ZEM^ÍK: AMPLITUDE–FREQUENCY RESPONSE OF AN ALUMINIUM CANTILEVER BEAM ...
96 Materiali in tehnologije / Materials and technology 49 (2015) 1, 95–98
piezoelectric patches (Figure 1), one on each side of a
patch, to eliminate the influence of the minor patch
curvature. The strain response to the applied static-elec-
tric voltage was measured and these values are presented
in Table 3. The difference between these two patches
loaded with 100 V was 14.4 %.
Table 3: Measured strains of the two patches (loaded with 100 V)
Tabela 3: Izmerjene napetosti dveh obli`ev (obremenjenih s 100 V)
Patch 1 / sensor 2 / actuator
11 by SG 1 10.27 × 10–5 11.74 × 10–5
11 by SG 2 9.38 × 10–5 11.23 × 10–5
11 – averaged 9.83 × 10–5 11.49 × 10–5
The piezoelectric-matrix coefficients were identified
using a FE model in Ansys v.14. The model was created
using 3D hexagonal elements: 20-node piezoelectric
bricks (SOLID 226) for PZT and 20-node structural
bricks (SOLID 186) for the protective foil. The piezo-
electric elements have an additional degree of freedom
for the electric potential in each node. One layer of these
elements was used, which was sufficient for an approxi-
mation of the electric potential across the thickness of a
patch. The nodes of the piezoelectric elements in the top
and bottom surfaces are represented by the silver elec-
trodes, where the relevant electric potential is applied.
The model parameter e31 was optimized to match the
experimental data. Coefficients e32 and e33 were calcul-
ated as e32 = e31 and e33 was chosen to maintain the
mutual ratio of e33/e31  –2.5. The resulting values are
presented in Table 4.
Table 4: Identified piezoelectric coefficients of the two patches
Tabela 4: Ugotovljeni piezoelektri~ni koeficienti pri dveh obli`ih
Patch 1 / sensor 2 / actuator
e31, e32 [C/m] 7.1 8.3
e33 [C/m] –17.8 –20.8
4 FREQUENCY RESPONSE OF THE
CANTILEVER BEAM
The calibrated patches were glued to an aluminium
beam with the dimensions of 1000 mm × 30 mm × 3 mm,
each on one side. The beam was clamped at 100 mm of
its length and 10 mm from the patches.
The patches were connected to the National Instru-
ments data acquisition system (NI CompactDAQ)
supplied with an actuating module NI 9215, sensing mo-
dule NI 9236 and an amplifier. The beam was actuated
by one of the patches and the vibration was sensed by
another patch and laser position sensor OptoNCDT. The
experimental set-up is presented in Figure 2.
The actuating signal was a linear chirp (sine wave
with a linearly swept frequency) defined with the follow-
ing equation:
x t X f
k
t( ) sin= ⋅ + +⎛⎝
⎜ ⎞
⎠
⎟⎡
⎣⎢
⎤
⎦⎥
0 0
22
2
π (6)
where X is the amplitude, 0 is the initial phase, f0 is the
initial frequency, k is the chirp rate defined by:
k
f f
t
=
−1 0
1
(7)
where f1 denotes the final frequency and t1 is the final
time.
Table 5: Parameters of the actuating signal
Tabela 5: Parametri vzbujevalnega signala
Amplitude X [V] 75
Initial phase 0 [rad] 0
Initial frequency f0 [Hz] 0
Final frequency f1 [Hz] 100
Final time t1 [s] 20
Sampling frequency fs [Hz] 50000
The properties of the actuating chirp signal are pre-
sented in Table 5. The range of frequencies was chosen
to contain the lowest natural frequencies of the beam for
the relevant two out-of-plane bending modes.
The time-voltage responses of the piezoelectric sen-
sor and laser sensor were recorded (Figure 3). A short-
time Fourier transform (STFT) was performed to obtain
a spectrogram (Figure 4). The length of the Hanning
window in STFT was set to 105 samples providing a
frequency precision of 0.5 Hz. The two lowest natural
Z. LA[OVÁ, R. ZEM^ÍK: AMPLITUDE–FREQUENCY RESPONSE OF AN ALUMINIUM CANTILEVER BEAM ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 95–98 97
Figure 2: Experimental set-up
Slika 2: Eksperimentalni sestav
Figure 1: Piezoelectric patch with the applied strain gauges, one on
the front side and one on the back side
Slika 1: Piezoelektri~ni obli` z napetostnimi listi~i: eden spredaj in
eden zadaj
frequencies can be detected in the spectrogram. In Table
6 the experimental results are compared with the nume-
rically calculated natural frequencies.
Table 6: Comparison of natural frequencies
Tabela 6: Primerjava naravnih frekvenc
FEM [Hz] Experiment [Hz]
1st 9.9 10.0
2nd 60.5 60.0
5 AMPLITUDE RESPONSE OF THE
CANTILEVER BEAM TO THE HARMONIC
LOADING
To obtain the amplitude response the structure was
loaded with a harmonic sine wave with a low frequency
(1 Hz) with a duration of 40 s to approach steady
oscillations. The amplitude was 100 V and the sampling
frequency was 25 kHz. The deflection of the beam tip
was measured with the laser sensor and compared with
the results of the FE static analysis (Table 7). The
difference between the experiment and FE was 2.4 %.
Table 7: Comparison of the experimental and numerical results
Tabela 7: Primerjava eksperimentalnih in numeri~nih rezultatov
Beam-tip deflection uz
Experiment 0.206 mm
FEA 0.201 mm
Difference 2.4 %
6 CONCLUSION
A numerical model of a piezoelectric transducer was
created in Ansys using three-dimensional piezoelectric
and structural finite elements. The piezoelectric coeffi-
cients of each patch were calibrated using strain gauges
and it was found that they differ by 14 %.
The FE model of the piezoelectric transducer was
tested on a problem of bending an aluminum cantilever
beam. The two lowest natural frequencies were deter-
mined experimentally and compared with the results of
the FE modal analysis with sufficient match for the given
frequency precision.
The amplitude of deflection of the beam loaded with
the low-frequency voltage was measured with the laser
sensor and compared to the result of the FE static
analysis with a difference of 2.4 %.
This numerical model proved to be suitable for
designing the SHM systems based on the change in the
structure’s natural frequencies. The reliability of the
model will be further tested for the case of transient
problems such as those used in the pitch-catch SHM
systems.
Acknowledgement
The work was supported by projects SGS-2013-036
and GA P101/11/0288.
7 REFERENCES
1 V. La{, R. Zem~ík, Progressive damage of unidirectional composite
panels, Journal of Composite Materials, 42 (2008) 1, 25–44
2 R. Zem~ík, V. La{, T. Kroupa, J. Barto{ek, Reconstruction of Impact
on Textile Composite Plate Using Piezoelectric Sensors, Proceedings
of the Ninth International Workshop on Structural Health Moni-
toring, Stanford, 2013, 393–400
3 R. Zem~ík, R. Kottner, V. La{, T. Plundrich, Identification of mate-
rial properties of quasi-unidirectional carbon-epoxy composite using
modal analysis, Mater. Tehnol., 43 (2009) 5, 257–260
4 V. Giurgiutiu, Structural Health Monitoring with Piezoelectric wafer
active sensors, Academic Press, Burlington 2008, 468–474
5 Piezo Material Data, Physik Instrumemte (PI), [cited 2013-09-25]
Available online on www.piceramic.com
Z. LA[OVÁ, R. ZEM^ÍK: AMPLITUDE–FREQUENCY RESPONSE OF AN ALUMINIUM CANTILEVER BEAM ...
98 Materiali in tehnologije / Materials and technology 49 (2015) 1, 95–98
Figure 4: Spectrogram of the cantilever beam, the two lowest natural
frequencies are marked with arrows
Slika 4: Spektrogram konzolnega nosilca; pu{~ici prikazujeta dve
najni`ji naravni frekvenci
Figure 3: Time response of the cantilever beam measured with PZT
Slika 3: ^asovni odziv konzolnega nosilca, izmerjen s PZT
T. KROUPA et al.: MICROMECHANICAL MODEL OF THE SUBSTITUENTS OF A UNIDIRECTIONAL ...
MICROMECHANICAL MODEL OF THE SUBSTITUENTS OF A
UNIDIRECTIONAL FIBER-REINFORCED COMPOSITE AND ITS
RESPONSE TO THE TENSILE CYCLIC LOADING
MIKROMEHANSKI MODEL NADOMESTKOV KOMPOZITOV,
OJA^ANIH Z ENOSMERNIMI VLAKNI, IN NJIHOV ODGOVOR NA
CIKLI^NO NATEZNO OBREMENJEVANJE
Tomá{ Kroupa1, Hana Srbová2, Robert Zem~ík1
1University of West Bohemia in Pilsen, NTIS – New Technologies for the Information Society, Univerzitní 22, 306 14, Plzeò, Czech Republic
2University of West Bohemia in Pilsen, Department of Mechanics, Univerzitní 22, 306 14 Plzeò, Czech Republic
kroupa@kme.zcu.cz
Prejem rokopisa – received: 2013-11-08; sprejem za objavo – accepted for publication: 2014-02-13
A micromechanical model respecting the behavior of a long-fiber unidirectional composite subjected to the cyclic tensile
loading was proposed. The so-called unit cell representing the structure of the composite material was created using the
finite-element-analysis software. Material models of the fibers and matrix proposed in this work are based on the non-linear
elastic and elasto-plastic behavior of real materials. The material and geometric parameters of the model are found by
minimizing the difference between the numerical and experimental results.
Keywords: unidirectional fiber composite, cyclic tensile test, micromodel, material parameters, non-linear, elasto-plastic
Predlagan je mikromehanski model, ki obravnava vedenje kompozita z dolgimi enosmernimi vlakni pri cikli~nemu nateznemu
obremenjevanju. Osnovna celica, ki pomeni zgradbo kompozita, je bila izdelana s programsko opremo za analizo kon~nih
elementov. Predlagani model materiala vlaken in osnove temelji na nelinearnem elasto-plasti~nem vedenju realnih materialov.
Material in geometrijski parametri modela so dobljeni z zmanj{evanjem razlike med numeri~nimi in eksperimentalnimi rezul-
tati.
Klju~ne besede: kompozit z enosmernimi vlakni, cikli~ni natezni preizkus, mikromodel, parametri materiala, nelinearno,
elasto-plasti~no
1 INTRODUCTION
Due to advantages, such as stiffness-to-weight and
strength-to-weight ratios, carbon-fiber-reinforced com-
posites are widely used for load-carrying structures in
various kinds of industries such as aerospace, auto-
motive, etc.
The behavior of composite materials can be modeled
on the micro-, meso- and macro-scale levels. The
micromechanical approach gives the most detailed
information about the constituents. The constituents and
the resulting composite exhibit, in general, a signi-
ficantly non-linear response. In this work, the material
models of the constituents include non-linear elastic and
elasto-plastic effects. The goal is to capture the response
of a unidirectional composite to the cyclic loading.
2 EXPERIMENT
In this work carbon-epoxy unidirectional conti-
nuous-fiber-composite-coupons were subjected to the
tensile cyclic loading. Rectangular specimens (coupons)
were cut with a water jet from a composite plate made of
eight equally oriented layers, the so-called prepregs,
HexPly 913C with fiber Tenax HTS 5631. With respect
to the loading direction five specimens were tested for
each of the following fiber orientations: 10°, 20°, 30°,
40°, 50°, 60°, 70°, 80° and 90° (Figure 1). Only two
specimens were successfully tested for the fiber orien-
tation of 0°. All the specimens were loaded in tension
Materiali in tehnologije / Materials and technology 49 (2015) 1, 99–102 99
UDK 66.017:669.018.9:620.178.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)99(2015)
Figure 1: Geometry of the coupons (mm), real and idealized parts of
the composite cross-section with a unit cell
Slika 1: Geometrija vzorcev (mm), realen in idealiziran del prereza
kompozita z osnovno celico
with the constantly increasing cycles up to failure. The
test control parameter was the applied load.
The force-displacement dependencies obtained from
the experiment were measured on the coupons with the
glass-textile (0°) or aluminum tabs (10°, 20°, 30°, 40°,
50°, 60°) bonded with epoxy high-strength and tough
adhesive Spabond 345. In the cases of the specimens
with the fiber orientations of 70°, 80° and 90° a non-
bonded emery cloth was used as the interface between
the grip and the coupon. The tab configuration sum-
marized in Table 1 produced the required gage-section
tensile failure (Figure 2).1
Table 1: Experiment configuration
Tabela 1: Konfiguracija eksperimenta
Fiber
orientation (°) Tab material
Tab length
(mm)
Extensometer
length (mm)
0 glass textile 60.0 10.0
10, 20, 30, 40,
50, 60 aluminum 25.0 60.0
70, 80, 90 emery cloth 25.0 60.0
2.1 Micromodel (finite-element model)
A three-dimensional idealized micromechanical
model of the composite material was proposed in the
form of a unit cell representing the periodical array of
the composite. The unit cell consists of two parts, the
part of the matrix and the part of the fibers. The fibers
and the matrix are perfectly bonded.
A perfect honeycomb distribution of the fibers is
assumed and the ratio of r:a:b (Figure 1) is given by the
fiber volume ratio Vf. The finite-strain theory was used.
The global coordinate system xyz is given by the
loading direction x and the direction perpendicular to the
composite surface z. The local coordinate system 123 is
defined by the unit-cell edges, where the axis direction 1
corresponds to the fiber direction, direction 3 is identical
to the axis direction z and 2 is perpendicular to axes 1
and 3 (Figure 3).
Assuming a uniaxial stress across the whole speci-
men, the behavior of the material can be simulated by
loading the unit cell with a normal stress:
 x
F
A
= (1)
where F is the external force and A is the cross-section
of the specimen.
The effect of the loading force is transformed to the
local coordinate system using the transformation:


 
    
  
1
2
12
2 2
2 2
⎡
⎣
⎢
⎢
⎢
⎤
⎦
⎥
⎥
⎥
=
cos sin sin cos
sin cos sin cos
sin cos sin cos cos sin
 
     

−
⎡
⎣
⎢
⎢
⎢
⎤
⎦
⎥
⎥
⎥
⋅
2 2
0
0
x⎡
⎣
⎢
⎢
⎢
⎤
⎦
⎥
⎥
⎥
(2)
where  is the angle of rotation between the local and
global coordinate systems.2
The results from the finite-element analysis (strains)
are transformed back to the global coordinate system
using the transformation:



    
 
x
y
xy
⎡
⎣
⎢
⎢
⎢
⎤
⎦
⎥
⎥
⎥
=
cos sin sin cos
sin cos sin
2 2
2 2  
     



cos
sin cos sin cos cos sin2 2 2 2
1
2
1−
⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥
⋅
2
⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥
(3)
Periodical boundary conditions were applied to
ensure that there was no separation or overlap between
the neighboring unit cells after the deformation.2,3
The micromodel was built in Abaqus/CAE (Com-
plete Abaqus Environment), consisting of eight-node
brick elements. The finite-element analysis was solved as
a quasistatic process using the Abaqus/standard solver.
The loading respected the experimentally obtained enve-
lope; however, only four cycles were prescribed to
reduce the computational time.
2.2 Material model of the fibers
The fibers were modeled as a transversely isotropic,
elastic material, respecting the non-linear behavior in the
axis direction. The nonlinearity is ensured with the form
of the stress-strain relationships:
T. KROUPA et al.: MICROMECHANICAL MODEL OF THE SUBSTITUENTS OF A UNIDIRECTIONAL ...
100 Materiali in tehnologije / Materials and technology 49 (2015) 1, 99–102
Figure 3: Rotated coordinate systems and loading of the unit cell
Slika 3: Zavrteni koordinatni sistemi in obremenitev osnovne celice
Figure 2: Cracked specimens. Final crack direction is parallel to the
fibers in the directions of 10°–90° and perpendicular for the direction
of 0°.
Slika 2: Vzorci z razpoko. Kon~na smer razpoke je vzporedna z
vlakni pri smereh 10°–90° in pravokotna pri smeri 0°.
 !   
  
11 11
1
2 11 11 12 22 33
22 12 11 22 22
1= + + +
= +
C C
C C
( ) ( )
+
= + +
=
=
C
C C C
G
G
23 23
33 12 11 23 22 22 33
12 12 12
13 12 1

   
 
  3
23 23 23 = G
(4)
where ! is the coefficient describing the measure of
nonlinearity with the growing strain 11 in the fiber
direction, while 22 and 33 are the strains in the direc-
tions transverse to the fibers. G12 and G13 are the shear
moduli in planes 12 and 13, respectively. The remaining
elastic constants are:
C
E v
v v v v v
C
E v v
11
1 23
2
12 21 23
2
21
2
23
12
1 21 2
1
1 2
1
=
−
− − −
=
+
( )
( 3
12 21 23
2
21
2
23
13 12
22
2 31 12
1 2
1
1
)
( )
− − −
=
=
−
−
v v v v v
C C
C
E v v
v12 21 23
2
21
2
23
23
2 23 21 12
12 21 23
2
2
1
v v v v
C
E v v v
v v v
− −
=
+
− −
( )
−
=
2 21
2
23
33 22
v v
C C
(5)
where E1 is the Young’s modulus in the fiber direction,
E2 is the Young’s modulus in the direction transverse to
the fibers and v12, v21 and v23 are Poisson’s ratios.4
2.3 Material model of the matrix
The matrix was modeled as an isotropic, elasto-
plastic material. The elastic parameters were Young’s
modulus Em and Poisson’s ratio vm. Von Mises plasticity
was used with the isotropic-hardening and work-hard-
ening functions proposed in the form of:
    
y
y= +0 ( )p (6)
where y is the yield stress, p is the equivalent plastic
strain,  0
y
is the initial yield stress, and  and 
 are the
shape parameters.
2.4 Identification process
The experimental results from the tensile cyclic tests
were processed using the Python programming language.
For each set of experiments belonging to the fiber direc-
tion , the arithmetic mean of displacement values at the
maximum point of each cycle, was found. By connecting
these averaged points with lines, the envelope curve was
obtained. Tangent k of a particular cycle was determined
by approximating the peak points with a linear function:
F k l q= +Δ (7)
individually for each experiment (Figure 4).
During the process of material-parameter optimiza-
tion, the residuum describing the difference between the
experimentally obtained envelope curve consisting of N
points and the numerical force-displacement data for the
cyclic tensile test:
r
l F l F
l
i i
Ni
N
=
−⎛
⎝
⎜
⎞
⎠
⎟
=
∑∑ Δ Δ
Δ
exp
exp
( , ) ( , )
( )
 

num
1
2
(8)
is minimized.
3 RESULTS
Material parameters were found for Vf = 0.55 and
r:a:b = 1:1.28:2.22 (Tables 2 and 3). Moreover, the
inaccuracy of the cutting orientation of the samples from
the plate was determined to be  = 1.7°.
The work-hardening function is shown in Figure 5.
Note that the force-displacement dependencies resulting
from the numerical analysis are in good agreement with
the dependencies obtained experimentally (Figures 6
and 7). However, there is not sufficient agreement
between the cycle tangents (Figure 8).
Table 2: Material parameters of fibers
Tabela 2: Parametri materiala vlaken
Parameter Units Value
E1 (GPa) 155.77
! (-) 14.44
E2 (GPa) 16.41
v12 (-) 0.30
v23 (-) 0.40
G12 (GPa) 50.35
T. KROUPA et al.: MICROMECHANICAL MODEL OF THE SUBSTITUENTS OF A UNIDIRECTIONAL ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 99–102 101
Figure 4: Principle of processing the data experimentally obtained
from multiple experiments into envelope curve F(l) and cycle
tangents k(l)
Slika 4: Na~in pridobitve eksperimentalnih podatkov iz ve~kratnih
preizkusov v skupno krivuljo F(l) in tangento ciklov k(l)
Table 3: Matrix-material parameters
Tabela 3: Parametri osnovnega materiala
Parameter Units Value
Em (GPa) 3.97
vm (–) 0.40
0
y (MPa) 10.14

 (–) 0.23
 (MPa) 182.71
4 CONCLUSION
Finite-element analyses of the cyclic tensile tests
were performed on unidirectional composite specimens
with various fiber orientations. Material parameters of
the orthotropic non-linear elastic material of the fibers
and of the non-linear elasto-plastic material of the matrix
were determined and good agreement with the
experiments was achieved.
In future work a degradation of the matrix will be
included into the material model in order to achieve
better agreement between the determined and experi-
mentally obtained cycle tangents.
Acknowledgement
The work was supported by projects SGS-2013-036
and European Regional Development Fund (ERDF),
project "NTIS - New Technologies for the Information
Society", European Centre of Excellence, CZ.1.05/
1.1.00/02.0090.
5 REFERENCES
1 ASTM Standard D 3039/D 3039M - 08, 2008, Standard Test Method
for Tensile Properties of Polymer Matrix Composite Materials,
ASTM International, West Conshohocken, PA 19428-2959, 2003,
www.astm.org
2 H. Srbová, T. Kroupa, R. Zem~ík, Identification of the Material Para-
meters of a Unidirectional Fiber Composite Using a Micromodel,
Mater. Tehnol., 46 (2012) 5, 431–434
3 S. W. Tsai, P. D. Shah, Strength & Life of Composites, Department
of Aeronautics & Astronautics, Stanford University, 2008
4 V. La{, Mechanics of Composite Materials, University of Bohemia,
Plzeò, 2004 (in Czech)
T. KROUPA et al.: MICROMECHANICAL MODEL OF THE SUBSTITUENTS OF A UNIDIRECTIONAL ...
102 Materiali in tehnologije / Materials and technology 49 (2015) 1, 99–102
Figure 8: Tangents of the force-displacement cycles (black – deter-
mined, gray – experiments)
Slika 8: Tangente ciklov sila – raztezek (~rno –dolo~eno, sivo – eks-
perimenti)
Figure 7: Force-displacement diagrams for 10° to 90° specimens
(black – determined, gray – experiments)
Slika 7: Diagrami sila – raztezek pri vzorcih od 10° do 90° (~rno –
dolo~eno, sivo – eksperimenti)
Figure 6: Force-displacement diagrams for the 0° specimen (black –
determined, gray – experiments)
Slika 6: Diagrami sila – raztezek pri vzorcu 0° (~rno – dolo~eno, sivo
– eksperimenti)
Figure 5: Matrix work-hardening function
Slika 5: Funkcija utrjevanja osnove
J. KORTNIK, B. MARKOLI: DRY-CUTTING OPTIONS WITH A CHAINSAW AT THE HOTAVLJE I NATURAL-STONE QUARRY
DRY-CUTTING OPTIONS WITH A CHAINSAW AT THE
HOTAVLJE I NATURAL-STONE QUARRY
MO@NOSTI SUHEGA REZANJA Z VERI@NO @AGO V
KAMNOLOMU NARAVNEGA KAMNA HOTAVLJE I.
Jo`e Kortnik1, Bo{tjan Markoli2
1University of Ljubljana, Faculty of Natural Sciences and Engineering, Department of Geotechnology and Mining, A{ker~eva 12, Ljubljana,
Slovenia
2University of Ljubljana, Faculty of Natural Sciences and Engineering, Department of Materials and Metallurgy, A{ker~eva 12, Ljubljana,
Slovenia
joze.kortnik@guest.arnes.si
Prejem rokopisa – received: 2013-12-10; sprejem za objavo – accepted for publication: 2014-01-17
This paper presents the results of a study on the possibility of replacing cutting inserts and introducing dry-cutting mode
(without water rinsing) with a chainsaw machine in the underground structures of the Hotavlje I quarry. In the experiment a
column-driven chainsaw machine Fantini G.70 was used with tungsten-carbide-based cutting inserts. The article presents the
physico-mechanical properties of the natural stone called the Hotavelj~an limestone and the state of discontinuities, which has a
significant impact on the efficiency of stone cutting. In the experiment the cutting characteristics of two types of inserts, type
H-13A and type H6T, were compared. The specimens of both types of the used inserts were subjected to a detailed metallo-
graphic investigation to evaluate the changes in the surface structure and hardness, the results of which are presented in detail.
Keywords: tungsten-carbide inserts, chainsaw machine, dimensional stone
V ~lanku so podani rezultati raziskave mo`nosti zamenjave rezalnih plo{~ic in vpeljave suhega na~ina rezanja (brez izpiranja z
vodo) naravnega kamna z veri`no `ago v podzemnih prostorih kamnoloma Hotavlje I. V poskusu je bila uporabljena veri`na
`aga Fantini G.70 na hidravli~nih stojkah. V ~lanku so predstavljene tudi fizikalno-mehanske lastnosti naravnega kamna, t. i.
pisanega apnenca "Hotavelj~ana", in stanje diskontinuitet, kar pomembno vpliva na u~inkovitost rezanja naravnega kamna. Pri
poskusu so bile primerjane rezalne karakteristike dveh vrst rezalnih plo{~ic tipa H-13A in H6T. Materiala obeh tipov rezalnih
plo{~ic sta bila po poskusu detajlno metalur{ko preiskana glede spremembe trdote povr{ine, rezultati pa so tudi podrobneje
prikazani.
Klju~ne besede: rezalne plo{~ice, veri`ni zasekovalni stroj, okrasni kamen
1 INTRODUCTION
The efforts to study and understand the cutting
mechanisms applied to hard limestone using chainsaw
machines are not new in Slovenia. The first chainsaw
machines were introduced in Slovenian natural-stone
quarries in the late nineties; they were originally
equipped with carbide cutting tools and later progres-
sively replaced with the new generation of tools. The
new generation of tools for metal machining included
inserts of tungsten carbide coated with a thin film of a
hard material that was chemically deposited on the
substrate and welded to the metal shank.
Until now, no research project aiming at a better
understanding of the stone-cutting mechanisms using a
chainsaw has been initiated in Slovenian natural-stone
quarries. For example, in Belgium the first results
obtained with this technique were widely published in
specialized literature1–9: Mingels (1971)10 and Focant
(1977)11 described sawing applications in Belgian red
marble; Boxho (1971)12, Brych (1975)13 and Neerdael
(1975)14 worked with Belgian blue stone. Later, very few
scientific publications were dedicated to the chain-
saw-cutting problems.
Research objectives were at first aimed at decreasing
the exploitation costs of the chainsaw machines and then
at optimizing the technological parameters in order to
correctly set up the cutting machines. The very first
studies tried to investigate the machine’s performance
with respect to the intrinsic properties of the stone being
cut, comparing it with the results of the cutting tests
performed in a laboratory. These studies highlighted the
considerable losses of the machines due to the friction
generated at the contact between the chain and the arm
of the machine, and also the losses in the hydraulic
system driving the machine.
Although the sawing techniques have been particul-
arly improved in the last 20 years, they are still proble-
matic when hard and/or abrasive stones are cut. The
main problems encountered with this type of materials
are a low productivity, a high consumption of cutting
tools and, thus, the related high production costs.
2 EXTRACTION OF NATURAL STONE FROM
THE HOTAVLJE I QUARRY
For the hard materials such as Hotavlje calcareous
limestone, brittle breakage is typical. This means that
this kind of material, when burdened with an additional
load, can only sustain a minor degree of plastic strains
(approximately 5 %). Thus, it can instantly break in a
Materiali in tehnologije / Materials and technology 49 (2015) 1, 103–110 103
UDK 621.93:621.936.6:622.35 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)103(2015)
plastic zone8. Because of these characteristic, the areas
with plastic zones are considered unstable and dangerous
for the stability of the underground structure.
The occurrence of cracking requires a consideration
of the rocks with cracks and their statuses. Thus, an
observed rock is called a mound. It is the cracks that
cause local differences in the mechanical properties of
the rock which differ considerably from those measured
in the laboratory. Examinations of mechanical and physi-
cal properties of natural stone were carried out in accord-
ance with European Union standards EN 12058:2004
(Natural stone products – Slabs for floors and stairs –
Requirements). Table 1 shows the physico-mechanical
properties of three color variations of natural decorative
stone called Hotavelj~an.15
Table 1: Mechanical properties of natural stone called Hotavelj~an15
Tabela 1: Mehanske lastnosti naravnega kamna, imenovanega "Hota-
velj~an"15
Type of
investigation Unit
Hotavlje
red
Hotavlje
gray
Hotavlje
gray-pink
Density without
pores and cavities t m
–3 2.90 2.74 2.73
Density t m–3 2.74 2.71 2.71
Coefficient of
density 0.94 0.99 0.99
Porosity % 0.06 0.01 0.70
Water absorption % 0.36 0.23 0.25
Grinding wear cm3 50 cm–2 29.5 27.1
Modulus of
elasticity GPa 25.0 25.0 25.0
Angle of internal
friction ° 22–35 – –
Bending tensile
strength MPa – 14.7
Bending tensile
strength after 48
freeze/thaw cycles
MPa – 13.6
Average compres-
sive strength: 90.2
– dry MPa 137.8 114.0 –
– wet MPa 150.8 156.0 –
– after 25 freeze/
thaw cycles MPa 187.8 118.0 –
Since the knowledge on rock-mass fracturing is very
important in our case in order to determine the geome-
chanical properties of a rock mass, it has to be described
by means of a statistical examination of the dispersion of
the measured cracks in the quarry and wells, and a
description of its occurrence. The main statistical cha-
racteristic of a particular occurrence of a large statistical
dispersion indicates that random cracks and crevices
thicken considerably with a mild dip to the northeast.
There are two quarries of colored calcareous stone in
the Hotavlje deposit. The Hotavlje I quarry, also called
the lower quarry, is the older and larger quarry, operating
in the northern part of the deposit, 435 m to 450 m above
sea level. In the central part of the deposit a new and
smaller quarry is in its opening phase. The Hotavlje II
quarry, also called the upper quarry, lies between 478 m
and 488 m above sea level. Both quarries are situated on
the steep eastern side of the Srednje Brdo Mountain16.
In 1993, in the colored-calcareous-stone Hotavlje I
quarry, a new method of natural-stone-block production
was introduced, called the underground stopping, which,
until then, had not been used in the Republic of Slovenia.
This method was proposed because of the geological
structure and the state of the quarry, and also due to an
increasing demand for this particular stone. Extraction in
the gallery is done with a diamond-wire saw and chain-
saw combined. This method proved to be the most
efficient one (Figure 1).
For the extraction of natural stone from the Hotavlje I
underground quarry a chain machine with hydraulic/
clamping columns is used (Figure 2). Hydraulic pillars
allow the machine blade with the cutting chain to move
in both horizontal and vertical directions. The cutting
chain machine is placed directly in front of the stope-
face and is it rigidly attached. Horizontal cuts are first
made at the bottom and top and finally on the sides.
Process water is used for cooling the saw chain and
washing the stone slurry out of the cut.
J. KORTNIK, B. MARKOLI: DRY-CUTTING OPTIONS WITH A CHAINSAW AT THE HOTAVLJE I NATURAL-STONE QUARRY
104 Materiali in tehnologije / Materials and technology 49 (2015) 1, 103–110
Figure 2: Column-driven chainsaw machine Fantini G.709
Slika 2: Veri`na `aga Fantini G.70 na hidravli~nih/vpenjalnih stebrih9
Figure 1: Column-driven chainsaw machine Fantini G.70 on the
location in the Hotavlje I quarry
Slika 1: Veri`na `aga Fantini G.70 na hidravli~nih/vpenjalnih stebrih
na lokaciji poskusa v kamnolomu Hotavlje I.
Table 2 shows technical details/characteristics of the
column-driven chainsaw machine Fantini G.70 used in
underground quarries.
Table 2: Technical properties of column-driven chainsaw machine
Fantini G.709
Tabela 2: Tehni~ni podatki veri`ne `age Fantini G.70 s hidravli~ni-
mi/vpenjalnimi stebri9
Weight of the chainsaw/Weight of the
hydraulic unit 6000 kg/2000 kg
Total installed power 52.2 kW (70 HP)
Chain rotation speed 0–0.71 m s–1
Arm/blade cutting-speed rate 0–0.07 m2 min–1
Cutting width 38 mm
Arm/blade length 2900 mm
Water consumption 20 L min–1
Cutting speed 2–4 m2 h–1
As presented in Table 2 the cutting speed is between
2 m2 h–1 and 4 m2 h–1, the cut width is 38 mm and the cut
depth is 2.40 m or more. The speed of the blade move-
ment is up to 0.07 m min–1 and the chain speed is up to
0.7 m s–1. The minimum stope-face width is 5.80 m. The
cutting method is based on cutting the rock using widia
or diamond plates that are mounted on the cutting chain
machine’s blade. The cutting machine’s blade moves
along the cogged rail and guiding rails. Process water is
used for cooling the saw chain and washing the stone
slurry out of the cut.
3 THEORY OF NATURAL-STONE CUTTING
An improvement of the sawing techniques in the
Hotavlje I quarry initially requires an understanding of
the cutting responses of the tools currently available on
the market. This implies a development of the working
methodology, making it possible to correctly charac-
terize the performances of cutting configurations inde-
pendently of the machines, the operators and any other
parameters that can influence the operations of the
machines. The field experimental approach was selected
because it makes it possible to very quickly estimate the
performances of various types of cutting tools used in a
quarry.
Rake angle 
: Cutting and normal forces decrease
monotonically with the increasing rake angle as seen in
Figure 3. Most of the benefit to the insert forces is
achieved at a rake angle of 20°, beyond which a further
marginal improvement weakens the cutting insert’s
strength and its potential to survive. The rake angle can
be either positive or negative. The rake angles between
+20° and +30° can be chosen for cutting weak rocks and
coal. High rake angles may not be beneficial since the
inserts with these angles are more susceptible to a gross
failure.
Wear angle : The wear flat is almost parallel to the
cutting direction; however, it generally tends to incline in
the opposite direction forming a wear angle. This angle
is only a few degrees and it becomes smaller for the
hardest and strongest materials. The occurrence of the
wear flat changes the tool-tip geometry and, conse-
quently, results in the generation of higher tool forces.
The normal force is the component most affected by the
wear, e.g., a wear flat of around 1.0 mm can drastically
increase the Fn/Ft ratio. It is also reported that a large
clearance angle relieves the wear effect and provides a
better overall efficiency even if, as a consequence, a
small or slightly negative rake angle is introduced.
Clearance angle : The clearance angle, which is
between the lower surface of the insert and the plane
parallel to the cutting direction, also has pronounced
effects on the cutting insert’s forces. The investigations
showed that the tool forces drop sharply after a value of
around 5° and stay sensibly constant. To meet the kine-
matic needs, the clearance angle is generally designed to
be around 10°.
When cutting stone with a chainsaw two forces are
acting on the cutting inserts (Figure 3):
• Axial (thrust) force of the chainsaw arm/blade Fn,
which allows the cutting inserts to progress from the
surface of the stone into the depth d,
• Force of the chain rotation Ft, which allows the
cutting and progression of the chainsaw arm/blade to
the depth d of the cutting grooves.
The cutting depth of individual cuts depends on the
axial force Fn which is proportional to and decreases
with the increasing cutting surface and the cutting resi-
stance of the rock (strength, toughness and hardness of
the rock). Rock cutting begins when the critical pressure
is reached, which must be greater than the strength of the
rock.
A chip formation can be described as a destruction of
the stone consistency using a tool. In the literature, the
models based on simple geometries of the cutting edges
divide the process of forming a cut into two mecha-
nisms1,2. The cutting mechanism of marble is explained
as a plastic deformation (a crushed zone) and a brittle
J. KORTNIK, B. MARKOLI: DRY-CUTTING OPTIONS WITH A CHAINSAW AT THE HOTAVLJE I NATURAL-STONE QUARRY
Materiali in tehnologije / Materials and technology 49 (2015) 1, 103–110 105
Figure 3: Forces active during the chainsaw cutting of natural stone
(
 – rake angle,  – wear angle,  – clearance angle)
Slika 3: Delovanje sil pri `aganju naravnega kamna z veri`no `ago (

– kot strganja,  – stri`ni kot,  – kot praznjenja rezi)
fracture of stone2,17. The plastic deformation and the
brittle fracture are influenced by the cutting conditions
such as the depth of cut, the tip shape of the cutting tool,
and the properties of the stone. When cutting stone with
diamond tools, the mechanical interaction of the tool and
the workpiece results in the process forces, mainly
caused by the following factors:
• elastic and plastic workpiece deformation by the
cutting edges;
• friction between the stone and the matrix;
• friction between the swarf and the matrix.
In front of the grains engaged in the process, stresses
are caused by tangential forces. In this zone, the insert is
forced out through the grooves in front of and beside the
grains. While the rock shows elastic characteristics up to
its ultimate stresses, it is necessary for the cutting to
reach a certain minimum grinding thickness. The mate-
rial to be cut is deformed by the compressive stress con-
ducted below the diamond. When the load is removed, an
elastic reversion leads to critical tensile stresses, which
cause a brittle fracture. This mechanism affected by
tensile stresses is termed the secondary chip formation.
The general insert is carried away by the coolant. The
following factors influencing this process are directly or
indirectly contained in the model:
• physical material properties of the stone;
• forces between the diamonds and the materials;
• stress distribution in the rock;
• temperatures at the tool/workpiece interface.
The cutting force and energy play important roles in
all the stone-machining processes. They are functions of
the maximum chip thickness and the geometry of the
idealized sawing chip.
The axial (thrust) force of the chainsaw arm/blade,
Fn, was calculated by Purti}18:
F
l d
k l dn
u
u=
⋅ ⋅ ⋅ +
⋅
= ⋅ ⋅ ⋅
 
 

 

sin( )
cos cos
2
2 (1)
where
u (MPa) – compressive strength of the rock
l (m) – length (thickness) of the cutting inserts
d (m) – depth of the insert penetration into the rock
k (m) – coefficient of friction.
Coefficient of friction:
k =
⋅cos cos
sin

 

"
2
(2)
The value of coefficient of friction k is dependent on
the angle of the rock internal friction  and the rake
angle 
:
Coefficient of friction k 0.27 0.30 0.36 0.44
Angle of rock internal friction /° 15 17 20 22
The cutting speed vcut is usually obtained with the
following equation:
v d m vcut chain= ⋅ ⋅ (m
2 s–1) (3)
where
d (mm) – depth of cutting
m (–) – number of cutting inserts
vchain (m s–1) – speed of the chain with cutting inserts.
The depth of cutting is strongly dependent on the
compressive strength of the rock and the force of the
chainsaw arm or blade (Figure 4). Specific energy is
also a very important factor in determining the efficiency
of cutting systems and it is defined as the work required
to excavate a unit volume of a rock. Hughes19 and Mellor20
demonstrated that specific energy can be formulated in
the following way:
S
EE
u=
 2
2
(4)
where
SE – specific energy
u – compressive strength of the rock
E – elasticity modulus.
The work done with the cutting force Fn is the work
needed to excavate a unit volume of yield. Dependent on
rock strength and toughness, degree of fracturing, ma-
chine type and method of operation, cutting insert type
and condition, available tool forces (machine size and
power) and penetration depth, the specific energy
amounts to:
S
EE
u MPa
0 MPa
.7 kPa= =
⋅
=
 2 2
2
90 2
2 25 00
162
.
.
Stone cutting is the result of the interference between
the insert grits and hard stone at the stone-tool interface.
The traditional tool is constituted of diamond grits that
are joined to the metal shank with the metal matrix. The
new-generation tools in metal machining are constituted
of inserts of tungsten carbide, coated with a thin film of
diamond that was chemically deposited on the substrate
J. KORTNIK, B. MARKOLI: DRY-CUTTING OPTIONS WITH A CHAINSAW AT THE HOTAVLJE I NATURAL-STONE QUARRY
106 Materiali in tehnologije / Materials and technology 49 (2015) 1, 103–110
Figure 4: Graph of the expected depth d of the insert penetration at
different compressive strengths of natural stone and the axial (thrust)
force of the chainsaw arm/blade Fn
Slika 4: Graf pri~akovanih globin d penetracije rezalnih plo{~ic v
naravnih kamnih razli~nih tla~nih trdnosti in razli~nih v naravnega
kamna ter z razli~nimi osnimi (potisnimi) silami me~a veri`ne `age Fn
(chemical-vapor-deposition (CVD) diamond inserts) and
welded to the metal shank2–4.
4 EXPERIMENTAL WORK: IN-SITU ATTEMPTS
TO USE DIFFERENT INSERTS
The presented research undertaken at the Faculty of
Natural Sciences and Engineering, the University of
Ljubljana, in 2009/2012 in collaboration with the Mar-
mor Hotavlje company was conducted to understand the
main problems encountered while cutting hard limesto-
nes such as Hotavelj~an and to propose technical solu-
tions for increasing the competitiveness of the column-
driven chainsaw machine Fantini G.70 in the Hotavlje I
quarry (Figure 1).
The main objectives of the research were as follows:
• Increasing the machine’s performance so that it
reaches a higher productivity when cutting the hard
Hotavelj~an limestone. The operators usually empiri-
cally determined the operating parameters of the
machines on the basis of their personal experiences.
Consequently, the performance was strongly influ-
enced. The determination of the optimum parameters
that should increase the performance required a
fundamental scientific study.
• Identifying the potential implementation of dry
cutting without the use of process water for rinsing
the stone debris from the cut and cooling the inserts.
• Increasing the lifespan of the cutting tools and
reducing the maintenance costs of the chain-saw
machines.
Two different values of the side rake angle, at the
insert/tool holders from n1–n6, angle –6° and at the
insert/tool holder n0, angle –25°, were adopted. These
were used to simulate the actual engagement between the
insert and stone during sawing and, therefore, the
cracking phenomena between the insert and stone during
sawing. A scheme of the tool engaged in stone cutting is
shown in Figures 5 and 6, while Figure 7 depicts a
detail of the actual set up of the cutting machine. The
cutting force along the cutting direction Fn and the
cutting force Ft perpendicular to the cutting direction
were evaluated from Figure 3.
In the experiment, two types of inserts were used,
namely:
• Fantini Sandvik H6T CT30 insert grade (a hardness
of 1753 HV) and
• Fantini Segatrice H-13A insert grade (a hardness of
1310 HV).
The cutting force along the cutting direction Fn and
the cutting force Ft perpendicular to the cutting direction
were calculated using the data from Table 3.
Table 3: Data on the cutting parameters used for the calculations
Tabela 3: Podatki o parametrih rezanja, uporabljenih za izra~une
vchain
Rotation speed of the chain with
tungsten-carbide inserts
Hitrost verige z rezilnimi plo{~icami
0–6.000
cm min–1
Arm/blade cutting speed rate
Hitrost pomika me~a pri `aganju
0–10
cm min–1
vcut
Average chain cutting speed
Povpre~na hitrost rezanja veri`ne `age 0–8 m
2 h–1
Width of the cut
[irina reza 38–42 mm
In addition to the data from Table 3, the values of the
cutting speed (0–60 m min–1 (the average of 55 m min–1))
J. KORTNIK, B. MARKOLI: DRY-CUTTING OPTIONS WITH A CHAINSAW AT THE HOTAVLJE I NATURAL-STONE QUARRY
Materiali in tehnologije / Materials and technology 49 (2015) 1, 103–110 107
Figure 7: Detail of the chain and arm/blade of the chainsaw machine
Slika 7: Detajl verige in zobnika roke/me~a veri`ne `age
Figure 6: Schematic view of the cutting system with 7 cutting insert
holders and their positions on the chainsaw chain (on the right, a
schematic view of the cutting grooves)
Slika 6: Shematski prikaz sistema razporeda 7 nosilcev in lege
nosilcev rezalnih plo{~ic na verigi veri`ne `age (na desni shematski
prikaz rezalne brazde)
Figure 5: Cutting configuration with 13 mm wide square cutters/
inserts (a system of 7 holders (n0–n6))
Slika 5: Postavitev rezanja 13 mm kvadratnih rezalnih plo{~ic (sistem
s 7 nosilci (n0–n6))
and many different values of the depth of cut (between
0.03 mm and 0.06 mm) were also used.
The forces exerted on the cutting inserts, as presented
in Figure 8, led to the wear of these inserts during the
cutting operation, presented in Figure 9.
The chainsaw cutting machine (Fantini G-70) ope-
rates using the following stepwise procedure to cut
natural stone: It starts in Position 0 (Figure 9) where the
initial cut using only the tip of the cutting blade is made
and it is then forced into Position 1 (a blade angle of 10°
to 72°). Then the blade is repositioned into Position 2
where it starts to move towards Position 3 (a blade angle
of 72° to 81°). In this arrangement the chain with the
cutting inserts is cutting the stone piece along the whole
of the width, i.e., from 0–250, generating the highest
forces exerted on the blade and on the cutting chain.
When Position 3 is finally reached, the blade starts to
rotate slightly so as to move into Position 4 (a blade
angle of 81° to 107°). Finally, it moves into Position 5 (a
blade angle of 107° to 166°). This is done in a manner
very similar to the one at the beginning of the cutting
process. The extent of horizontal cuts is limited to a
depth of 2.50 m and the minimum length of 5.75 m.
Vertical cuts are limited to a depth of 2.5 m and a height
of 4.6 m. The construction of the cutting machine also
enables the cuts to be made under different angles,
exerting additional forces on the blade and the chain,
thus leading to premature wear of the cutting inserts.
5 RESULTS AND DISCUSSION
The results of the cutting process described in the
section about the experimental work associated with
Figure 9 are depicted in Figures 10 and 11. Two diffe-
rent cutting inserts were employed for the cutting pur-
poses. Both cutting inserts were mounted on the same
chain of the Fantini G-70 cutting machine so as to make
the comparison of the cutting inserts as reliable as
possible. Figure 10 shows the behavior of the H6T C30
Fantini Sandvik cutting inserts mounted on the cutting
chain cutting the natural stone from the Hotavlje quarry
in the horizontal manner. For this purpose the following
process parameters were used: the chain speed varied
from 53.0 m min–1 to 55.3 m min–1 and the average cutt-
ing speed increased from 3.5 m2 h–1 to 5.5 m2 h–1. The
entire duration of the pure horizontal cut with the chain-
saw took 186 minutes. Throughout the cutting process,
the consumption of time was also recorded. In the case
of the cutting inserts H-13A Fantini Segatrice (Figure
11) the horizontal cutting manner was also applied and
the chain speed varied from 51.5 m min–1 to 55.6
m min–1, while the average cutting speed varied from 6.3
m2 h–1 to 7.9 m2 h–1. The entire duration of the pure
horizontal cut with the chainsaw took 169 min. A com-
parison of the results from Figures 10 and 11 indicated
that the yield in the case of the H-13A cutting inserts was
higher in relation to the average cutting speeds, whilst
the time consumption was pretty much the same.
However, the variation in the chain speed was smaller in
the case of the H6T cutting inserts which applies to the
average cutting speeds as well. In the case of H6T, these
showed less dissipation or variation in all the positions or
incisions (Pos. 0 to Pos. 5) which inevitably led to less
pronounced wear of the cutting inserts due to frag-
mentation than in the case of the H-13A cutting inserts.
When the H-13A cutting inserts were used, the variations
in the chain speeds were much more obvious as the
J. KORTNIK, B. MARKOLI: DRY-CUTTING OPTIONS WITH A CHAINSAW AT THE HOTAVLJE I NATURAL-STONE QUARRY
108 Materiali in tehnologije / Materials and technology 49 (2015) 1, 103–110
Figure 8: a) Insert holder n5 with a rake angle 
 = 9° and b) a tung-
sten-carbide insert
Slika 8: a) Nosilec rezalne plo{~ice n5 s kotom strganja 
 = 9° in
b) volfram-karbidna rezalna plo{~ica
Figure 10: Results of a horizontal cut with the H6T cutting inserts
Slika 10: Rezultati rezanja horizontalnega reza z rezalnimi plo{~icami
H6T
Figure 9: Stages (Pos. 0 to Pos. 5) of making a horizontal cut with
two types of cutting inserts
Slika 9: Faze (Poz. 0 do Poz. 5) rezanja horizontalnega reza z raz-
li~nima tipoma rezalnih plo{~ic
speeds varied, for all the incisions (Pos. 0 to Pos. 5), with
sharp peaks and drops.
This led to less uniform wear and a more frequent
deformation of the cutting inserts due to fragmentation.
The recorded peaks and drops in the chain speeds in the
case of H-13A are visible in Figure 12 showing an array
of grooves made in a stepwise manner as the mechanical
properties of the natural stone were changing locally, as
were the mechanical properties of the cutting inserts.
The measurements of the surface and subsurface
microhardness of the cutting inserts showed that the
degradation of the surface hardness was severe but it
was, in all the cases, limited to an area within a few
tenths of a millimeter under the surface (Table 4).
A brief metallographic overview of the applied
cutting plates provided us with a definite proof of surface
and subsurface degradation. This was observed using
LM and the microhardness measurements made on the
surface and, in a stepwise manner, also under the surface.
Table 4 summarizes the Vickers hardness measurements
with the average number of measurements of 5.
Table 4: Microhardness measurements of HV 0.05 for H-13A and
H6T CT30 cutting inserts
Tabela 4: Mikrotrdote HV 0,05 rezalnih plo{~ic H-13A in H6T CT30
Cutting insert Average hardnesson the surface
Average hardness
0.1 mm under the
surface
Grade H-13A 1240.75 1471.40
Grade H6T CT30 1269.00 1614.67
It is clear that in both cases the hardness degrades
quite rapidly and dramatically from the surface towards
the inner region of a cutting plate, which is a conse-
quence of the thermomechanical load on the surface of
the cutting plate. On the micrographs of the H-13A
cutting inserts in Figures 13a and 13b severe material
wear can be observed on the cutting edges of the cutting
inserts due to a complete rounding and cracking of the
initially rectangular edges. On the top (Figure 13b), the
indents of the diamond pyramid can still be seen allow-
ing clear indications of the places where the micro-
hardness HV measurements were made. The protective
TiN layer of the surface and the edges of H6T CT30 is
still intact, which was also confirmed with the HV
measurements showing a higher hardness of the cutting
inserts after the operation (Figure 13c). It has to be
noted here that the H-13A cutting inserts were used for
the so-called dry cutting, whereas the H6T CT30 inserts
were water cooled. Namely, the H-13A cutting inserts
were constructed to be used for dry cutting although we
believe that the use of a lubricant or coolant would most
likely prolong their service time too.
J. KORTNIK, B. MARKOLI: DRY-CUTTING OPTIONS WITH A CHAINSAW AT THE HOTAVLJE I NATURAL-STONE QUARRY
Materiali in tehnologije / Materials and technology 49 (2015) 1, 103–110 109
Figure 13: LM micrographs of the cutting inserts after the operation:
a) cutting edges of H-13A cutting inserts and b) with clear positions of
the indents (circle) of the diamond pyramid used in HV measure-
ments. Uniformity in the grain size of WC can be seen in b). The
surface of H6T CT30 exhibits fewer defects and still has the protective
layer of TiN. Positions of the indents caused by the diamond pyramid
are visible in c).
Slika 13: LM-posnetki uporabljenih rezalnih plo{~ic: a) rezalni robovi
plo{~ic H-13A in b) z lepo vidnimi vtiski po uporabi diamantne pira-
mide v okviru HV-meritev. Na sliki b) je prikazana enakomerna
velikost delcev WC. Povr{ina rezalne plo{~ice H6T CT30 izkazuje
veliko {tevilo majhnih po{kodb na {e vedno prisotni za{~itni plasti iz
TiN. Na sliki c) so vidna mesta z diamantno piramido nastalih vtiskov.
Figure 12: Visible grooves or a cut profile made by the cutting inserts
Slika 12: Vidni kanal oziroma profil reza, ki nastaja med rezanjem z
rezalnimi plo{~icami
Figure 11: Results of a horizontal cut with the H-13A cutting inserts
Slika 11: Rezultati rezanja horizontalnega reza z rezalnimi plo{~icami
H-13A
6 CONCLUSIONS
The present work shows that tungsten-carbide H-13A
inserts allow an increase in the cutting efficiency by
more than 66 % and, therefore, a decrease in the amounts
of force and energy involved in stone sawing. This
means that the tool is exposed to the stresses lower than
those affecting a tungsten-carbide insert probably
causing a reduction of tool wear. Moreover, the cutting
power of the tungsten carbide H-13A inserts is lower
than that of the tungsten-carbide H6T CT30 inserts.
A single model describes well the actual cutting force
caused by the interaction between a tungsten-carbide
insert and stone. This means that a unique simple equa-
tion can represent the relationship between the cutting
force and the chip-cutting area.
An analysis of the tungsten-carbide insert wear over
time and of the relationship between the tungsten-car-
bide wear and the cutting-force value compared to the
behavior of the chemical-vapor-deposition (CVD)
diamond inserts is currently the topic of further studies
in order to verify that this type of tool is suitable for
sawing the Hotavelj~an limestone.
Acknowledgement
Special thanks go to Marmor Hotavlje; despite the
fact that the study was not formally funded, they
unselfishly shared their knowledge and contributed a lot
of voluntary work, engaging their human resources and
their hardware.
7 REFERENCES
1 F. I. Dagrain, Understanding stone cutting mechanisms for the design
of new cutting tools sequences and the cutting optimization of chain
saw machines in Belgian blue stone quarries, Diamante, Aplicazioni
& Technologia, 17 (2011) 66, 55–67
2 S. Turchetta, L. Carrino, W. Polini, CVD diamond insert in stone
cutting, Diamond and Related Materials, 14 (2005) 3–7, 641–645
3 H. K. Tönshoff, H. Hillmann-Apman, J. Asce, Diamond tools in
stone and civil engineering industry: cutting principles, wear and
application, Diamond and Related Materials, 11 (2002) 3–6,
736–741
4 Y. Li, H. Huang, J. Y. Shen, X. P. Xu, Y. S. Gao, Cost-effective
machining of granite by reducing tribological interactions, Journal of
Materials Processing Technology, 129 (2002) 1–3, 389–394
5 H. D. Jerro, S. S. Pang, C. Yang, R. A. Mirshams, Kinematics
analysis of the chipping process using the circular diamond saw
blade, Journal of Manufacturing Science and Engineering, 121
(1999) 2, 257–264
6 K. Brach, D. M. Pai, E. Ratterman, M. C. Shaw, Grinding forces and
energy, Journal of Manufacturing Science and Engineering, 110
(1988) 1, 25–31
7 H. K. Tönshoff, J. Peters, I. Inasaki, T. Paul, Modelling and
simulation of grinding processes, Annals of the CIRP, 41 (1992) 2,
677–688
8 T. W. Hwang, S. Malkin, Grinding mechanisms and energy balance
for ceramics, Journal of Manufacturing Science and Engineering,
121 (1999) 4, 623–631
9 Fantini Technical Paper, 2002, 4
10 C. Mingels, Utilisation d’une haveuse Perrier dans une carriere de
marbre, Annales des Mines de Belgique, 2 (1971) 1, 121–130
11 J. Focant, Utilisation des haveuses Perrier dans I’exploitation des
carrieres de Marbres Rouge, Annales des Mines de Belgique, 9
(1977) 2, 819–823
12 J. Boxho, Essai d’una haveuse Perrier dans le Petit Granit, Annales
des Mines de Belgique, 2 (1971) 3, 131–142
13 J. Brych, Les outils autoaffutants au service du havage des roches
dures, Annales des Mines de Belgique, 2 (1975) 1, 165–180
14 B. Neerdael, Contribution ill’etude de la destructibilite des roches
dures telles que les callcaires crino’idiques de Soignies, Faculte
Polytechnique de Mons, Travail de fin d’etudes, 1975, 5–18
15 K. Fifer, Natural stone "Hotavlje" geotechnical report / Poro~ilo o
preiskavi naravnega kamna "Hotavlje", ZAG, Department of mate-
rials, Ljubljana 2005, 35
16 U. Baj`elj, J. Kortnik, B. Petkov{ek, K. Fifer, Environment friendly
underground excavation of dimension stone, 27th International
Symposium on Computer Applications in the Mineral Industries,
London, 1998, 11–22
17 J. Kortnik, U. Baj`elj, Underground mining of natural stone in Slo-
venia, 20th World Mining Congress, Tehran, 2005, 277–286
18 N. Purti}, Drilling and Blasting/Bu{enje i miniranje, Rudarsko-geo-
lo{ki fakultet, Beograd 1991, 20–22
19 H. Hughes, Some aspects of rock machining, Int. J. Rock Mech.
Min. Sci., 9 (1972) 2, 205–211
20 M. Mellor, Normalization of specific energy values, Int. J. Rock
Mech. Min. Sci., 9 (1972) 5, 661–663
J. KORTNIK, B. MARKOLI: DRY-CUTTING OPTIONS WITH A CHAINSAW AT THE HOTAVLJE I NATURAL-STONE QUARRY
110 Materiali in tehnologije / Materials and technology 49 (2015) 1, 103–110
I. GUNES: EFFECT OF SLIDING SPEED ON THE FRICTIONAL BEHAVIOR AND WEAR PERFORMANCE ...
EFFECT OF SLIDING SPEED ON THE FRICTIONAL BEHAVIOR
AND WEAR PERFORMANCE OF BORIDED AND
PLASMA-NITRIDED W9Mo3Cr4V HIGH-SPEED STEEL
VPLIV HITROSTI DRSENJA NA VEDENJE IN OBRABO
BORIRANEGA IN V PLAZMI NITRIRANEGA HITROREZNEGA
JEKLA W9Mo3Cr4V
Ibrahim Gunes
Department of Metallurgical and Materials Engineering, Faculty of Technology, Afyon Kocatepe University, 03200 Afyonkarahisar, Turkey
igunes@aku.edu.tr
Prejem rokopisa – received: 2013-12-11; sprejem za objavo – accepted for publication: 2014-02-06
This study investigated the effect of sliding speed on the friction and wear behaviors of plasma-nitrided (PN) and borided
(W9Mo3Cr4V) steels. The PN process was carried out in a dc-plasma system at a temperature of 923 K for 6 h in a gas mixture
of 80 % N2-20 % H2 under a constant pressure of 5 mbar. The boriding process was carried out in Ekabor-II powder at a
temperature of 1223 K for 6 h. X-ray diffraction analysis on the surface of the PN and borided steels revealed the presence of
FeB, Fe2B, CrB, MoB, WB, FeN, Fe2N, Fe3N and Fe4N compounds. The wear tests were carried out in a ball-disc arrangement
under dry friction conditions at room temperature with an applied load of 10 N and with sliding speeds of (0.1, 0.3 and 0.5) m/s
at a sliding distance of 1000 m. The wear surfaces of the samples at the different sliding speeds were analyzed using SEM
microscopy and X-ray energy-dispersive spectroscopy EDS. The friction coefficients of the nitrided and borided W9Mo3Cr4V
steels varied from 0.42 to 0.61, and from 0.33 to 0.52, respectively. As a result of the wear, the friction coefficient was observed
to decrease with an increase in the sliding speeds of the PN and the borided AISI W9Mo3Cr4V steel while an increase was
observed in the wear resistance.
Keywords: plasma nitriding, boriding, W9Mo3Cr4V, sliding speed, wear rate
V tej {tudiji je bil preu~evan vpliv hitrosti drsenja na trenje in vedenje v plazmi nitriranega (PN) in boriranega jekla
(W9Mo3Cr4V). PN-postopek je bil izvr{en v enosmernem plazemskem sistemu 6 h na temperaturi 923 K in v plinski me{anici
80 % N2-20 % H2 pri konstantnem tlaku 5 mbar. Postopek boriranja je bil izvr{en s prahom Ekabor-II pri temperaturi 1223 K in
trajanju 6 h. Rentgenska difrakcijska analiza na povr{ini PN in boriranega jekla je odkrila spojine FeB, Fe2B, CrB, MoB, WB,
FeN, Fe2N, Fe3N in Fe4N. Preizkus obrabe je bil izvr{en na napravi s kroglo na plo{~i pri suhem trenju na sobni temperaturi z
obte`bo 10 N in s hitrostjo drsenja (0,1, 0,3 in 0,5) m/s pri razdalji drsenja 1000 m. Obraba na povr{ini pri vzorcih z razli~no
hitrostjo drsenja je bila analizirana s SEM-mikroskopijo in rentgensko energijsko disperzijsko spektroskopijo. Koeficient trenja
nitriranega in boriranega jekla W9Mo3Cr4V je bil med 0,42 in 0,61 oziroma 0,33 in 0,52. Kot rezultat obrabe PN in boriranega
jekla AISI W9Mo3Cr4V se je koeficient trenja zmanj{eval pri nara{~anju hitrosti drsenja, odpornost proti obrabi pa se je
pove~ala.
Klju~ne besede: nitriranje v plazmi, boriranje, W9Mo3Cr4V, hitrost drsenja, stopnja obrabe
1 INTRODUCTION
Plasma nitriding and boriding are thermo-chemical
surface treatments widely used to improve the tribo-
mechanical properties of engineering components via a
modification of their surface microstructure.1–6 During
the plasma-nitriding process, the nitriding reaction not
only occurs on the surface but also in the sub-surface
owing to the long-distance diffusion of nitrogen atoms
from the surface towards the core. Nitrogen diffused into
a steel surface is combined with alloying elements to
form a fine dispersion of alloy nitrides.7,8 As a result, two
different structures have been identified, the so-called
"white" or "compound" layer and the "diffusion zone".
The first one is the outermost layer, and consists of one
or two iron nitrides (Fe4N–Fe2-3N), depending on the
process parameters. The second layer results from the
reactions produced by the N diffusion, such as precipi-
tation of nitrides, 
-Fe saturation, changes in the residual
tensions and C redistributions.9,10
During the boriding process, boron atoms can diffuse
into ferrous alloys due to their relatively small size and
very mobile nature. They can dissolve in iron inter-
stitially, but can react with it to form FeB and Fe2B
intermetalic compounds. Depending on the potential of
the medium and the chemical compositions of the base
materials, a single or a duplex layer may be formed.
Borided steel components display excellent performance
in several tribological applications in the mechanical
engineering and automotive industries.11–14
Plasma nitriding and boriding are used in numerous
applications in industries such as the manufacture of
machine parts for plastics and food processing,
packaging and tooling, as well as pumps and hydraulic
machine parts, crankshafts, rolls and heavy gears, motor
and car construction, cold- and hot-working dies and
cutting tools. The wear behavior of plasma-nitrided and
borided steels has been evaluated by a number of
researchers.15–20 However, there is no information about
Materiali in tehnologije / Materials and technology 49 (2015) 1, 111–116 111
UDK 621.785.53:620.178.1:539.92 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)111(2015)
the effect of sliding speed on the friction and wear
behaviors of plasma-nitrided and borided W9Mo3Cr4V
steel. The main objective of this study was to investigate
the effect of sliding speed on the friction and wear
behaviors of plasma-nitrided and borided W9Mo3Cr4V
steel. Structural and tribological properties were
investigated using light microscopy, XRD, SEM, EDS,
microhardness tests and a ball-on-disc tribotester.
2 EXPERIMENTAL METHOD
2.1 Plasma Nitriding, Boriding and Characterization
The W9Mo3Cr4V steel contained mass fractions (w)
0.82 % C, 4.10 % Cr, 3.80 % Mo, 7.40 % W and 1.30 % V.
The test samples were cut into Ø 25 mm × 8 mm
dimensions and ground up to 1000 G and polished using
a diamond solution. The PN process was carried out in a
dc plasma system at a temperature of 923 K for 6 h in a
gas mixture of 80 % N2-20 % H2 under a constant
pressure of 5 mbar. The boriding heat treatment was
carried out by using a solid boriding method with
commercial Ekabor-II powders. The samples that were to
be borided were packed in the powder mixture and
sealed in a stainless-steel container. The boriding heat
treatment was performed in an electrical resistance
furnace under atmospheric pressure at 1223 K for 6 h
followed by air cooling.
The microstructures of the polished and etched
cross-sections of the samples were observed under an
Olympus BX-60 light microscope. The presence of
nitrides and borides formed in the coating layer was
detected by means of X-ray diffraction equipment
(Shimadzu XRD 6000) using Cu K
 radiation. The
diffusion zone on the nitrided sample and the thickness
of borides were measured by means of a digital thickness
measuring instrument attached to a light microscope
(Olympus BX60). Micro-hardness measurements were
performed from the surface to the interior along a line in
order to see variations in the hardness of the nitride and
boride layers, the transition zone and the matrix,
respectively. The microhardness of the boride layers was
measured at 8 different locations at the same distance
from the surface by means of a Shimadzu HMV-2
Vickers indenter with a load of 50 g and the average
value was taken as the hardness.
2.2 Friction and Wear
A ball-on-disc test device was used to perform the
friction and wear tests of the borided samples. In the
wear tests, WC-Co balls of 8 mm in diameter supplied
by H. C. Starck Ceramics GmbH were used. Errors
caused by the distortion of the surface were eliminated
by using a separate abrasion element (WC-Co ball) for
each test. The wear experiments were carried out in a
ball-disc arrangement under dry friction conditions at
room temperature with an applied load of 10 N and with
sliding speeds of (0.1, 0.3 and 0.5) m/s at a sliding
distance of 1000 m. Before and after each wear test, each
sample and abrasion element was cleaned with alcohol.
After the test, the wear volumes of the samples were
quantified by multiplying the cross-sectional areas of the
wear by the width of the wear track obtained from the
Taylor-Hobson Rugosimeter Surtronic 25 device. The
wear rate was calculated using the following formula:
Wk = Wv/(M · S) mm
3/(N m) (1)
where Wk is the wear rate, Wv the worn volume, M is the
applied load and S is the sliding distance.
Friction coefficients depending on the sliding dist-
ance were obtained through a friction coefficient pro-
gram. Surface profiles of the wear tracks on the samples
and surface roughness were measured by a Taylor-
Hobson Rugosimeter Surtronic 25. The worn surfaces
were investigated by scanning electron microscopy and
energy-dispersive X-ray spectroscopy (EDS).
3 RESULTS AND DISCUSSION
3.1 Characterization of Plasma Nitriding and Boriding
Coatings
The cross-sections of the optical micrographs of the
nitrided and borided W9Mo3Cr4V steel are given in
I. GUNES: EFFECT OF SLIDING SPEED ON THE FRICTIONAL BEHAVIOR AND WEAR PERFORMANCE ...
112 Materiali in tehnologije / Materials and technology 49 (2015) 1, 111–116
Figure 1: Cross-section of plasma-nitrided and borided W9Mo3Cr4V
steel: a) PN, b) borided
Slika 1: Prerez v plazmi nitriranega in boriranega jekla W9Mo3Cr4V:
a) PN, b) borirano
Figure 1. A typical light photograph of the cross-sec-
tional microstructure of the plasma nitrided
W9Mo3Cr4V sample is given in Figure 1a. This
photograph shows the formation of a diffusion zone on
the sample. However, the formation of a continuous
compound (white) layer on top of the nitrided layer of
the sample is not visible in Figure 1a. Following etching
with a 3 % nital reagent, the nitrided (diffusion) layer
appears as a dark zone on the sample. Around 250 μm
depth of diffusion was obtained from the plasma-nitrided
sample. As can be seen in Figure 1b, the boride layer
was formed on the W9Mo3Cr4V steel borided with a
solid boriding method. Depending on the process time,
the temperature and the chemical composition of
substrates, the thickness of the boride layer was found to
be 55 μm.
The X-ray diffraction patterns of the nitrided and
borided W9Mo3Cr4V steel are given in Figure 2. The
XRD results showed that the nitrided W9Mo3Cr4 steel
contained FeN, Fe2N, Fe3N, Fe4N and M6C phases (Fig-
ure 2a). The XRD results showed that the boride layers
formed on the W9Mo3Cr4V steel contained FeB, Fe2B,
CrB, MoB and WB phases in Figure 2b. The micro-
hardness and the wear resistance of the nitride and
boride layers formed on the steel surface are to a large
extent known with the help of these phases.21–24
Microhardness measurements were carried out on the
cross-sections from the surface to the interior along a
line, as can be seen in Figure 3. The hardness of the
nitrided W9Mo3Cr4V steel varied between 995 HV0.05
and 1286 HV0.05. The hardness of the boride layer on the
W9Mo3Cr4V steel varied between 1378 HV0.05 and 1814
HV0.05. On the other hand, the Vickers hardness values
were 230 HV0.05 for the untreated W9Mo3Cr4V steel.
When the hardnesses of the nitride and boride layers are
compared with the matrix, the nitride layer’s hardness is
approximately five times greater and the boride layer’s
hardness is approximately eight times greater than that of
the matrix.
3.2 Friction and Wear Behavior
The wear mechanism is defined as the physical and
chemical phenomena that take place during wear.
Abrasive wear refers to the mechanism where two bodies
rubbing against each other break off pieces from each
other. The hardness, shape and size or roughness of the
I. GUNES: EFFECT OF SLIDING SPEED ON THE FRICTIONAL BEHAVIOR AND WEAR PERFORMANCE ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 111–116 113
Figure 4: Surface roughness values of the plasma-nitrided and bo-
rided W9Mo3Cr4V steel
Slika 4: Hrapavost povr{ine v plazmi nitriranega in boriranega jekla
W9Mo3Cr4V
Figure 2: X-ray diffraction patterns of plasma-nitrided and borided
W9Mo3Cr4V steel: a) PN, b) borided
Slika 2: Rentgenska difrakcija v plazmi nitriranega in boriranega jekla
W9Mo3Cr4V: a) PN, b) borirano
Figure 3: The variation of hardness depth in the plasma-nitrided and
borided W9Mo3Cr4V steel
Slika 3: Spreminjanje trdote po globini v plazmi nitriranega in
boriranega jekla W9Mo3Cr4V
abrasive material, the angle of contact, the normal load
applied, the sliding speed and the fracture toughness of
the material are all among the important factors in wear
mechanisms.25 Figure 4 shows the surface roughness
values of the nitrided and borided and unborided
W9Mo3Cr4V steel. For the nitrided and borided
W9Mo3Cr4V steel, it was observed that the surface
roughness values increased, as can be seen in Figure 4.
It was observed that the surface roughness of the nitrided
W9Mo3Cr4V steel was higher than that of the borided
W9Mo3Cr4V. These results indicate that both nitriding
and boriding treatments affect the surface roughness of
the specimens. While the surface roughness value of the
borided sample is 0.84 μm, it increased to 0.33 μm as a
result of the boriding process. While the roughness level
of the surface of the non-nitrided sample is 0.93 μm at
the end of the boriding process, it has increased to 0.198
μm.
Figure 5 shows the friction coefficients of the
nitrided and borided W9Mo3Cr4V steels at different
sliding speeds. The friction coefficients of the nitrided
and borided W9Mo3Cr4V steels varied from 0.61 to
0.42 and from 0.52 to 0.33, respectively. The friction
coefficients of both borided and nitrided W9Mo3Cr4V
steels reduced with the increase of the sliding speed up
to 0.5 m/s. The reason behind this may be the oxides on
the surface of the sample. During the wear test, oxides
formed on the steel surface, allowing for low friction
coefficients. The wear tracks on the samples may have
been oxidized due to the frictional heat.
As the sliding speed increased, the steady-state fric-
tion coefficient values slightly decreased. This may be
explained by a decrease in the shear strength at the
sliding interface. Decreasing the shear strength may
occur in two ways. Firstly, the hardness and hence the
shear strength of the coatings decrease with increasing
temperature. It is reported that the hardness of the boride
coatings decreases by a factor of more than two with
increasing temperature up to 1000 °C.12 Secondly, oxi-
dation products (iron, boron, molybdenum, wolfram and
chromium oxides) can reduce the friction. In fact, many
of the oxides have high friction.
Figure 6 shows the effect of sliding speed on the
wear rate of plasma-nitrided and borided W9Mo3Cr4V
steel. It was observed that the borided W9Mo3Cr4V
steel had a lower wear rate than that of the plasma-
nitrided W9Mo3Cr4V steel. For the microhardness of
the FeB, CrB and WB phases, the borided steel showed
more resistance to wear due to the nitrided steel. The
wear test results indicated that the wear resistance of the
nitrided and borided steels was higher than that of the
untreated steel. It is well known that the hardness of the
nitride and boride layers plays an important role in the
improvement of the wear resistance.
As shown in Figures 3 and 6, the relationship bet-
ween the microhardness and the wear resistance of the
nitrided and borided samples also confirms that the wear
resistance was improved by the increase in the hardness.
This is in agreement with reports of previous studies.26–33
In addition, reductions in wear rates were observed with
an increasing rate of sliding speed. The oxides formed on
the samples borided at high sliding speeds may have
caused a reduction in the friction coefficient and an
increase in the wear rate because oxides on the sample
surface act as a solid lubricant. Molinari et al.34 investi-
gated the tribological properties of a plasma-nitrided
Ti-6A14V alloy at different loads and sliding speeds. It
was observed that the increase in the sliding speed at low
loads did not significantly affect the wear volume, while
the wear volume decreased with an increase in the
sliding speed at high loads. Straffelini et al.35 investi-
gated the impact of sliding speed and contact pressure on
the oxidative wear of austempered ductile iron and
reported that the friction coefficient and wear rate
decreased with an increasing rate of sliding speed.
I. GUNES: EFFECT OF SLIDING SPEED ON THE FRICTIONAL BEHAVIOR AND WEAR PERFORMANCE ...
114 Materiali in tehnologije / Materials and technology 49 (2015) 1, 111–116
Figure 6: The effect of the sliding speed on the wear rate of the
plasma-nitrided and borided W9Mo3Cr4V steel
Slika 6: Vpliv hitrosti drsenja na stopnjo obrabe v plazmi nitriranega
ali boriranega jekla W9Mo3Cr4V
Figure 5: The effect of the sliding speed on the friction coefficient of
the plasma-nitrided and borided W9Mo3Cr4V steel
Slika 5: Vpliv hitrosti drsenja na koeficient trenja jekla W9Mo3Cr4V,
nitriranega v plazmi ali boriranega
I. GUNES: EFFECT OF SLIDING SPEED ON THE FRICTIONAL BEHAVIOR AND WEAR PERFORMANCE ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 111–116 115
Figure 8: Sliding-speed-dependent SEM micrographs of the wear surfaces of the borided W9Mo3Cr4V steel at 1223 K for 6 h: a) 0.1 m/s, b) 0.3
m/s, c) 0.5 m/s, d) EDS analysis
Slika 8: SEM-posnetki obrabljene povr{ine v odvisnosti od hitrosti drsenja za 6 h borirano jeklo W9Mo3Cr4V pri 1223 K: a) 0,1 m/s, b) 0,3 m/s,
c) 0,5 m/s, d) EDS-analiza
Figure 7: Sliding-speed-dependent SEM micrographs of the wear surfaces of the nitrided W9Mo3Cr4V steel at 923 K for 6 h: a) 0.1 m/s, b) 0.3
m/s, c) 0.5 m/s, d) EDS analysis
Slika 7: SEM-posnetki obrabljene povr{ine v odvisnosti od hitrosti drsenja za 6 h nitrirano jeklo W9Mo3Cr4V pri 923 K: a) 0,1 m/s, b) 0,3 m/s,
c) 0,5 m/s, d) EDS-analiza
The SEM micrographs and X-ray energy-dispersive
spectroscopy EDS of the worn surfaces of the nitrided
and borided samples are illustrated in Figures 7 and 8.
Figure 7 shows the SEM micrographs of the wear sur-
faces of nitrided W9Mo3Cr4V steel worn at different
sliding speeds. The wear region of the nitrided samples,
debris, surface grooves and cracks on the layer are
shown in Figures 7a to 7c. It was observed that delami-
nation wears, which form as a result of the progress of
microcracks, occurred on the wear tracks of the nitrided
samples. The wear intensity on the surfaces of the ni-
trided samples was observed to decrease with an increase
in the sliding speed (Figures 7a to 7c). Oxidation
products (iron, boron, wolfram and chromium oxides)
formed as a result of the wear test in Figure 7d.
Figure 8 shows the SEM micrographs of the wear
surfaces of borided W9Mo3Cr4V steel worn at different
sliding speeds. There were microcracks, abrasive parti-
cles and small holes on the worn surface of the boride
coatings (Figures 8a to 8c). In the wear region of the
borided W9Mo3Cr4V steel, there were cavities probably
formed as a result of layer fatigue (Figure 8b) and
cracks concluded in the delaminating wear (Figure 8c).
It was observed that delamination wear, which forms as a
result of the progress of microcracks, occurred on the
wear tracks of the borided samples. Oxide layers formed
as a result of the wear test in Figure 8d. The spallation
of the oxide layers in the sliding direction and their
orientation extending along the wear track were iden-
tified (Figures 7d and 8d).
4 CONCLUSIONS
The following conclusions may be drawn from the
present study.
• The multiphase nitride coating that thermo-chemi-
cally occurred on the W9Mo3Cr4V steel consisted of
FeN, Fe2N, Fe3N, Fe4N and M6C phases, while the
boride coating consisted of FeB, Fe2B, CrB, MoB
and WB phases.
• The surface hardness of the nitrided W9Mo3Cr4V
steel was in the range of 995–1286 HV0.05, the surface
hardness of the borided W9Mo3Cr4V steel was in the
range 1378–1814 HV0.05 and the untreated
W9Mo3Cr4V steel substrate was 230 HV0.05.
• The surface roughness values of both the plasma-
nitrided and borided W9Mo3Cr4V steel increased
compared to the untreated W9Mo3Cr4V steel.
• The friction coefficients of the plasma-nitrided and
borided W9Mo3Cr4V steel varied from 0.42 to 0.61,
and from 0.33 to 0.52, respectively.
• Reductions were observed in the wear rates of both
the nitrided and borided W9Mo3Cr4V steel.
• As a result of an increase in the sliding speed, it was
observed that the friction coefficient decreased while
the wear resistance of the plasma nitrided and
borided W9Mo3Cr4V steel increased.
5 REFERENCES
1 A. G. von Matuschka, Boronizing, Heyden and Son Inc., Philadel-
phia, USA 1980, 11
2 C. Bindal, A. H. Ucisik, Surface and Coatings Technology, 122
(1999), 208
3 A. Akbari, R. Mohammadzadeh, C. Templier, J. P. Riviere, Surface
and Coatings Technology, 204 (2010), 4114
4 A. Alsaran, A. Celik, C. Celik, Surface and Coatings Technology,
160 (2002), 219
5 M. Hudakova, M. Kusy, V. Sedlicka, P. Grgac, Mater. Tehnol., 41
(2007) 2, 81–84
6 N. Ucar, O. B. Aytar, A. Calik, Mater. Tehnol., 46 (2012) 6, 621–625
7 S. Y. Sirin, E. Kaluc, Materials and Design, 36 (2012), 741
8 S. Y. Sirin, K. Sirin, E. Kaluc, Materials Characterization, 59 (2008),
351
9 P. Corengia, G. Ybarra, C. Moina, A. Cabo, E. Broitman, Surface
and Coatings Technology, 200 (2005), 2391
10 M. B. Karamis, Wear, 150 (1991), 331
11 S. Taktak, Materials and Design, 28 (2007), 1836
12 S. Taktak, Surface and Coatings Technology, 201 (2006), 2230
13 C. Bindal, A. H. Ucisik, Surface and Coatings Technology, 94–95
(1997), 561
14 C. Martini, G. Palombarini, G. Poli, D. Prandstraller, Wear, 256
(2004), 608
15 A. Alsaran, Materials Characterization, 49 (2003), 171
16 B. Podgornik, J. Vizintin, V. Leskovsek, Surface and Coatings Tech-
nology, 108–109 (1998), 454
17 M. Ulutan, M. M. Yildirim, O. N. Celik, S. Buytoz, Tribology
Letters, 38 (2010), 231
18 T. Eyre, Wear, 34 (1975), 383
19 A. Çelik, Ö. Bayrak, A. Alsaran, I. Kaymaz, A. F. Yetim, Surface
and Coatings Technology, 202 (2008), 2433
20 F. Yýldýz, A. F. Yetim, A. Alsaran, A. Çelik, I. Kaymaz, Tribology
International, 44 (2011), 1979
21 I. Gunes, S. Ulker, S. Taktak, Materials and Design, 32 (2011), 2380
22 L. G. Yu, X. J. Chen, K. A. Khor, G. Sundararajan, Acta Materialia,
53 (2005), 2361
23 I. Gunes, Transactions of the Indian Institute of Metals, 67 (2014),
359
24 Y. T. Xi, D. X. Liu, D. Han, Surface and Coatings Technology, 202
(2008), 2577
25 M. Tabur, M. Izciler, F. Gül, I. Karacan, Wear, 266 (2009), 1106
26 E. Atik, U. Yunker, C. Meric, Tribology International, 36 (2003), 155
27 D. C. Wen, Surface and Coatings Technology, 204 (2009), 511
28 I. Gunes, A. Dalar, Journal of Balkan Tribology Association, 19
(2013), 325
29 B. Selcuk, R. Ipek, M. B. Karamýs, Journal of Materials Processing
Technology, 141 (2003), 189
30 C. Li, B. Shen, G. Li, C. Yang, Surface and Coatings Technology,
202 (2008), 5882
31 S. Ulker, I. Gunes, S. Taktak, Indian Journal of Engineering Mate-
rials Science, 18 (2011), 370
32 M. B. Karamiº, E. Gerçekcioðlu, Wear, 243 (2000), 76
33 I. Gunes, Journal of Materials Science and Technology, 29 (2013),
662
34 A. Molinari, G. Straffelini, B. Tesi, T. Bacci, G. Pradelli, Wear,
203–204 (1997), 447
35 G. Straffelini, M. Pellizzari, L. Maines, Wear, 270 (2011), 714
I. GUNES: EFFECT OF SLIDING SPEED ON THE FRICTIONAL BEHAVIOR AND WEAR PERFORMANCE ...
116 Materiali in tehnologije / Materials and technology 49 (2015) 1, 111–116
B. STOJANOVI] et al.: TRIBOLOGICAL BEHAVIOUR OF A356/10SiC/3Gr HYBRID COMPOSITE ...
TRIBOLOGICAL BEHAVIOUR OF A356/10SiC/3Gr HYBRID
COMPOSITE IN DRY-SLIDING CONDITIONS
TRIBOLO[KO VEDENJE HIBRIDNEGA KOMPOZITA
A356/10SiC/3Gr PRI SUHEM DRSENJU
Bla`a Stojanovi}, Miroslav Babi}, Nenad Miloradovi}, Slobodan Mitrovi}
Faculty of Engineering, University of Kragujevac, Sestre Janjic 6, 34000 Kragujevac, Serbia
blaza@kg.ac.rs
Prejem rokopisa – received: 2014-01-21; sprejem za objavo – accepted for publication: 2014-03-12
The paper presents tribological behaviour of hybrid aluminium composite A356/10SiC/3Gr in dry-sliding conditions. Hybrid
composites were made with a modified compocasting procedure and tribological tests were conducted on a tribometer with a
block-on-disc contact geometry. An analysis of the effects of three sliding speeds ((0.25, 0.5 and 1) m/s) and two normal loads
(10 N and 20 N) on the wear volume and wear rate was conducted for different sliding distances. The total sliding distance was
300 m. SEM and an EDS analysis of the wear surfaces of the tested material were also performed. The analysis of the wear
surfaces shows an intensive adhesive wear of hybrid composites. Also, the appearance of mechanically mixed layers (MMLs)
was confirmed, i.e., an increase in the concentration of iron in the hybrid composite material.
Keywords: hybrid composites, wear, aluminium, SiC, graphite, testing
^lanek opisuje tribolo{ko vedenje hibridnega aluminijevega kompozitnega materiala A356/10SiC/3Gr pri suhem drsenju.
Hibridni kompoziti so bili izdelani s prilagojenim postopkom "kompokasting", tribolo{ki preizkusi pa so bili opravljeni na
tribometru s kontaktno geometrijo "blok na disku". Opravljena je analiza vpliva treh drsnih hitrosti ((0,25, 0,5 in 1) m/s) in dveh
normalnih obremenitev (10 N in 20 N) na intenzivnost in velikost obrabe pri razli~nih dol`inah drsenja. Skupna dol`ina drsenja
je bila 300 m. Hkrati sta bili izvr{eni tudi SEM-mikroskopija in EDS-analiza obrabljene povr{ine preizku{anega materiala.
Analiza odrgnjenih povr{in poka`e intenzivno adhezivno obrabo hibridnega kompozita. Potrjen je bil pojav mehansko me{anih
plasti (MML) oziroma pove~ane koncentracije `eleza v hibridnem kompozitnem materialu.
Klju~ne besede: hibridni kompoziti, obraba, aluminij, SiC, grafit, preizku{anje
1 INTRODUCTION
Composite materials with aluminium bases have an
increasing application in the automotive, airline and
aerospace industries. In the automotive industry, these
materials are used for building engines, cylinder linings,
valve tappets, brakes, cardan shafts, etc. An improve-
ment in the mechanical and tribological characteristics is
achieved by adding appropriate reinforcements to these
materials. Graphite (Gr), silicon carbide (SiC) and
alumina (Al2O3) are commonly used as reinforcements.
By combining two types of reinforcements, hybrid
composites with improved characteristics with regard to
the basic material are obtained. The composites rein-
forced with SiC and graphite are called Al/SiC/Gr hybrid
composites1–5.
Tribological behaviour of the hybrid composites with
Al2219 aluminium bases and reinforced with (5, 10 and
15) % SiC and 3 % Gr was investigated by Basavara-
jappa et al.6 The composite materials were made with the
liquid-metallurgy procedure. Tribological tests were
realized on a tribometer with a pin-on-disc contact
geometry in accordance with the ASTM G99-95
standard and they showed that an increase in the SiC
fraction increases the wear resistance of hybrid compo-
sites and the wear is decreased. The wear rate increases
with the increasing sliding speed and normal load.
Mahdavi and Akhlaghi7 tested the tribological cha-
racteristics of Al/SiC/Gr hybrid composites obtained
with an in-situ powder-metallurgy technique. A hybrid
composite had a matrix made of the Al6061 aluminium
alloy reinforced with 9 % Gr and (0–40) % SiC. The
wear rate of the tested composites decreased with an
increase in the SiC amount up to 20 %. A further
increase in the SiC amount increased the wear. Other
investigations of the same authors relate to the hybrid
composites with 30 % SiC and a variable amount of gra-
phite (0–13 %). The results of the tribological investi-
gations showed that the wear rate also increases with the
increasing sliding distance and the increasing graphite
amount.
Suresha and Sridhara8 investigated the tribological
behaviour of the hybrid composites obtained using a
stir-casting procedure, with variable SiC and Gr amounts
(0–5 %) and the total ratio of 10 %. Ravindran et al.9
investigated the tribological behaviour of the hybrid
composites with aluminium alloy A2024 bases. The
hybrid composites were obtained with powder metal-
lurgy using 5 % of SiC and (0. 5 or 10) % of Gr. Hybrid
composite Al/5SiC/5Gr (5 % SiC and 5 % Gr) showed
the best tribological characteristics, while a further
increase in the graphite amount increased the wear.
The literature gives a small amount of information
about the effect of graphite and SiC particles on wear
Materiali in tehnologije / Materials and technology 49 (2015) 1, 117–121 117
UDK 539.92:669.018.9:66.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)117(2015)
properties. Considering the previous statements, the aims
of this paper are to investigate the tribological behaviour
of the hybrid composites with an A356-aluminium-alloy
base reinforced with SiC and Gr and to provide new
information and knowledge. The compocasting proce-
dure was selected to obtain hybrid composites because it
is economical and it enables a good distribution of the
reinforcements and a favourable composite structure.
The quality/price ratio of the obtained composites is the
best for this procedure, which is also very important. An
investigation of tribological characteristics was con-
ducted by varying normal loads and sliding speeds for
different sliding distances in dry-sliding conditions. An
analysis of the wear was conducted on the basis of the
obtained values of the wear rates and SEM and EDS
analyses of the test samples. Within the framework of
this investigation, the tribological behaviour of the
Al/SiC/Gr hybrid composites obtained with the com-
pocasting procedure and an A356-aluminium-alloy base
reinforced with 10 % SiC and 3 % Gr is presented.
2 EXPERIMENT
2.1 Procedure for obtaining the composites
Hybrid composites were obtained with the compo-
casting procedure or with an infiltration of the particles
into the half-hardened melt of the A356 alloy. The
preparation of the materials consisted of the chemical
cleaning of the base (the A356 alloy) and its infiltration
into a previously preheated electro-resistance furnace
crucible, the melting and overheating up to 650 °C (the
liquid-phase domain) to clean the slag. In order to obtain
a hybrid composite (A356 alloy + 10 w(SiC) + 3 w(Gr)),
the measured quantities of SiC and Gr were previously
homogeneously mixed in the solid state, then preheated
at 150 °C and later used in the infiltration process. The
commercial procedure T6 was used for the heat treat-
ment, consisting of solution annealing at 540 °C (for
6 h), water quenching and artificial aging at a tempera-
ture of 160 °C for 6 h. The mean value of the SiC-par-
ticle diameter was 39 μm and for the graphite-particle
diameter it was 35 μm. The newly obtained composite
had w(SiC) = 10 % and w = 3 % of graphite.
Figure 1a presents the structure of the basic material
and Figure 1b shows the structure of the Al/10SiC/3Gr
composite. In the hybrid composite, it is noticed that the
base is well filled with the reinforcement particles, so
that the base surface without any particles is decreased,
which indicates a good distribution of the particles in the
base. The soft particles of graphite did not sustain their
mean value (35 μm) during the procedure of obtaining
the composite. In fact, their erosion and size reduction
occurred in the process of preparation (mixing them with
the SiC particles).
2.2 Description of the experiment and equipment
The wear behaviour of a block was monitored in
terms of the wear-scar width (Figure 2). Using the
wear-scar width and the geometry of the contact pair, the
wear volume (in accordance with ASTM G77-83) and
the wear rate (expressed in mm3/m) were calculated.
The tested blocks were made of the basic A356
material and the A356/10SiC/3Gr hybrid composite. The
block dimensions were 6.35 mm × 15.75 mm × 10.16
mm, while the roughness of the ground contact surface
was Ra = 0.2 μm. The counter body (disc) was made of
90MnCrV8 steel with a hardness of 62–64 HRC, a dia-
meter of 35 mm and a width of 6.35 mm. The roughness
of the contact surface of the disc was Ra = 0.3 μm. Tribo-
B. STOJANOVI] et al.: TRIBOLOGICAL BEHAVIOUR OF A356/10SiC/3Gr HYBRID COMPOSITE ...
118 Materiali in tehnologije / Materials and technology 49 (2015) 1, 117–121
Figure 2: Scheme of the contact-pair geometry
Slika 2: Shematski prikaz geometrije kontakta
Figure 1: Metallographic structure of: a) basic material A356, b)
A356/10SiC/3Gr hybrid composite
Slika 1: Mikrostruktura: a) osnovni material A356, b) hibridni
kompozit A356/10SiC/3Gr
logical tests were conducted in dry-sliding conditions for
different sliding speeds ((0.25, 0.5 and 1.0) m/s) and
loads (10 N and 20 N). All the experiments were re-
peated 5 times.
The tests were performed for a sliding distance of
300 m. The first tests were used to form wear-rate cur-
ves, while the other tests were conducted without
stopping.
3 EXPERIMENT TEST RESULTS
The results of tribological tests were obtained for
different sliding speeds and different normal loads in
dry-sliding conditions and they are presented in the
following diagrams. The diagrams showing the depen-
dence between the material wear volume and the sliding
distance were obtained on the basis of the wear-scar
widths. The wear-scar widths were measured after the
(30, 60, 90, 150 and 300) m of the sliding distance.
Figure 3 shows the variations in the wear volume in
relation to the sliding distance for all three sliding speeds
((0.25, 0.5 and 1.0) m/s). Full lines show the results for
the basic A356 material, while dashed lines present the
results for the A356/10SiC/3Gr hybrid composite.
The run-in distance for the tested materials depends
on the sliding speeds and normal loads and its range was
between 50 m and 100 m. This is best seen on the
wear-rate curves or the curves showing the dependence
between the wear-scar widths and the sliding distance10.
After this sliding distance, the tested materials entered
the period of normal wear. When analysing the obtained
diagrams, it became clear that the curves obtained for the
base material were steeper than the curves for the hybrid
composite. It was obvious that an addition of SiC and Gr
prolonged the period of normal wear of the tested hybrid
materials.
The effects of the sliding speed and normal loads on
the wear rate of the tested materials for a sliding distance
of 300 m are shown in Figures 4 and 5. With an increase
in the sliding speed, the intensity of the wear increases
for both the base A356 material and for the hybrid
A356/10SiC/3Gr material. The dependence between the
wear and the sliding speed is almost linear (Figure 4).
Figure 5 shows that an increase in normal load
induces an increase in the wear rate for all the test regi-
mes, i.e., for all the sliding speeds. According to the
positions of the curves, it may be concluded that the
effect of normal load on the wear rate is quite pro-
nounced. Also, according to the positions of the curves,
it is obvious that the wear rate of the hybrid composite
with SiC and graphite is considerably smaller than the
wear rate of the base material, especially at the normal
load of 10 N. The wear rate of the basic A356 material
for the normal load of 20 N is (on average) three times
higher than the wear rate for the normal load of 10 N.
B. STOJANOVI] et al.: TRIBOLOGICAL BEHAVIOUR OF A356/10SiC/3Gr HYBRID COMPOSITE ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 117–121 119
Figure 4: Dependence between the wear rate and the sliding speed
Slika 4: Odvisnost med obrabo in hitrostjo drsenja
Figure 3: Dependence between the wear volume and the sliding
distance for all three sliding speeds
Slika 3: Odvisnost med volumnom obrabe in dol`ino poti drsenja za
vse tri hitrosti drsenja
The wear rate of the A356/10SiC/3Gr hybrid composite
for the normal load of 20 N is (on average) seven times
higher than the wear rate for the normal load of 10 N.
4 DISCUSSION
For a more complete analysis, Figures 6a and 6b
show the SEM microphotographs of the worn surfaces of
the tested materials. These microphotographs were ob-
tained during the tribological tests with the sliding speed
of 0.25 m/s and the normal load of 10 N.
When observing the samples, it is clear that the
adhesive wear is the basic wear mechanism. Adhesive
wear occurred as a result of an alternate formation and
destruction of the frictional connections caused by an
atomic and intermolecular interaction of the boundary
layers of the contact bodies. Due to the adhesive wear,
there is a pulling of the material out of the contact sur-
face of the composite material (block). As a consequence
of the adhesive wear, pits of irregular shapes and depths
are visible on the surface of the block6 and in Figure 7.
A tribolayer was formed on the surface of the
A356/10SiC/3Gr composite during sliding and it formed
a protective lubricating film preventing a direct contact
between the aluminium matrix and the steel counter
body. High-volume-fraction ceramic particles act as
load-bearing elements, preventing the subsurface damage
during sliding and suppressing the spalling of tribolayers
from the surface.
In order to rationalize the wear behaviour of the
observed composite, it is important to analyze the com-
position of the microstructure of the tribolayer formed
during the wear process. The EDS analysis was per-
formed together with the SEM of the tested material. The
EDS analysis was employed to determine the elements
on the worn surface, as shown in Figure 8. An occur-
rence of iron (Fe) and oxygen (O) may be spotted on the
analyzed samples. The presence of iron and its oxide
confirms that the wear of the steel disc occurred. The
disc wear occurs due to the effects of SiC from the com-
posite material as well as of the Si phase from the A356
aluminium matrix. Worn particles of iron enter the
surface layer of the composite causing a formation of a
mechanically mixed layer (MML) which is characteristic
for a MMC with an aluminium matrix7,11–13. The pre-
sence of iron and its oxides is confirmed by EDS Spec-
trum 2 in Figure 8. Spectrum 1 in Figure 8 shows the
presence of SiC in the hybrid composite. The lubrication
B. STOJANOVI] et al.: TRIBOLOGICAL BEHAVIOUR OF A356/10SiC/3Gr HYBRID COMPOSITE ...
120 Materiali in tehnologije / Materials and technology 49 (2015) 1, 117–121
Figure 7: Forming of the adhesive-wear pits
Slika 7: Nastanek jamic zaradi adhezivne obrabe
Figure 5: Dependence between the wear rate and normal load
Slika 5: Odvisnost med obrabo in obremenitvijo
Figure 6: SEM microphotographs of worn surfaces of: a) A356, b)
A356/10SiC/3Gr
Slika 6: SEM-posnetka obrabe povr{ine: a) A356, b) A356/10SiC/3Gr
layer is composed of the Al, Si, Fe, O and C elements.
The occurrence of oxygen implies that the surface ele-
ments, such as Fe, were at least partially oxidized during
the sliding. Thus, it may be concluded that the tribolayer
is composed of a mixture of iron oxides, graphite and
fractured SiC particles and some fine particles contain-
ing aluminium.
5 CONCLUSIONS
The testing of the tribological behaviour of the hybrid
A356/10SiC/3Gr composite in dry-sliding conditions
shows superior characteristics of this material compared
to those of the base A356 material. With an increase in
the sliding distance, the amount of the wear or wear rate
of the hybrid composite is always smaller than that of the
base material. The run-in sliding distance of the tested
materials ranges between 50 m to 100 m, after which the
materials enter the period of normal wear. The wear rate
of the tested materials increases with the increasing
sliding speed and normal load for all the test regimes.
SEM confirms a homogenous distribution of SiC parti-
cles within the hybrid composite. Additionally, SEM
microphotographs and an EDS analysis show the pre-
sence of Fe and its oxides as well as the formation of a
mechanically mixed layer (MML).
Acknowledgement
This paper presents the results obtained during the
research within the framework of project TR35021
supported by the Ministry of Education, Science and
Technological Development of the Republic of Serbia.
6 REFERENCES
1 A. Vencl, Tribology of the Al-Si alloy based MMCs and their appli-
cation in automotive industry, In: L. Magagnin (ed.), Engineered
Metal Matrix Composites: Forming Methods, Material Properties
and Industrial Applications, Nova Science Publishers, Inc., New
York, USA 2012, 127–166
2 B. Stojanovi}, M. Babi}, S. Mitrovi}, A. Vencl, N. Miloradovi}, M.
Panti}, Tribological characteristics of aluminium hybrid composites
reinforced with silicon carbide and graphite, A review, Journal of the
Balkan Tribological Association, 19 (2013) 1, 83–96
3 M. Babi}, B. Stojanovi}, S. Mitrovi}, I. Bobi}, N. Miloradovi}, M.
Panti}, D. D`uni}, Wear Properties of A356/10SiC/1Gr Hybrid
Composites in Lubricated Sliding Conditions, Tribology in Industry,
35 (2013) 2, 148–154
4 S. Mitrovi}, M. Babi}, B. Stojanovi}, N. Miloradovi}, M. Panti}, D.
D`uni}, Tribological Potential of Hybrid Composites Based on Zinc
and Aluminium Alloys Reinforced with SiC and Graphite Particles,
Tribology in Industry, 34 (2012) 4, 177–185
5 A. Vencl, I. Bobic, S. Arostegui, B. Bobic, A. Marinkovic, M. Babic,
Structural, mechanical and tribological properties of A356 alumi-
nium alloy reinforced with Al2O3, SiC and SiC + graphite particles,
Journal of Alloys and Compounds, 506 (2010) 2, 631–639
6 S. Basavarajappa, G. Chandramohan, A. Mahadevan, Influence of
sliding speed on the dry sliding wear behaviour and the subsurface
deformation on hybrid metal matrix composite, Wear, 262 (2007)
7–8, 1007–1012
7 S. Mahdavi, F. Akhlaghi, Effect of the Graphite Content on the
Tribological Behavior of Al/Gr and Al/30SiC/Gr Composites
Processed by In Situ Powder Metallurgy (IPM) Method, Tribology
Letters, 44 (2011) 1, 1–12
8 S. Suresha, B. K. Sridhara, Friction characteristics of aluminium
silicon carbide graphite hybrid composites, Materials & Design, 34
(2012), 576–583
9 K. Ravindran, R. Manisekar, P. Narayanasamy, N. Narayanasamy,
Tribological behaviour of powder metallurgy-processed aluminium
hybrid composites with the addition of graphite solid lubricant,
Ceramics International, 39 (2013) 2, 1169–1182
10 S. Mitrovi}, N. Miloradovi}, M. Babi}, I. Bobi}, B. Stojanovi}, D.
D`uni}, Wear behaviour of hybrid ZA27/SiC/graphite composites
under dry sliding conditions, Tribological Journal Bultrib, 3 (2013)
1, 142–147
11 J. Rodriguez, P. Poza, M. A. Garrido, A. Rico, Dry sliding wear
behaviour of aluminium–lithium alloys reinforced with SiC particles,
Wear, 262 (2007) 3–4, 292–300
12 R. N. Rao, S. Das, D. P. Mondal, G. Dixit, Effect of heat treatment
on the sliding wear behaviour of aluminium alloy (Al–Zn–Mg) hard
particle composite, Tribology International, 43 (2010) 1–2, 330–339
13 A. K. Mondal, S. Kumar, Dry sliding wear behaviour of magnesium
alloy based hybrid composites in the longitudinal direction, Wear,
267 (2009) 1–4, 458–466
B. STOJANOVI] et al.: TRIBOLOGICAL BEHAVIOUR OF A356/10SiC/3Gr HYBRID COMPOSITE ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 117–121 121
Figure 8: EDS analysis of hybrid A356/10SiC/3Gr composite
Slika 8: EDS-analiza hibridnega kompozita A356/10SiC/3Gr

V. K. SINGH et al.: NEW SOLID-POLYMER-ELECTROLYTE MATERIAL FOR DYE-SENSITIZED SOLAR CELLS
NEW SOLID-POLYMER-ELECTROLYTE MATERIAL FOR
DYE-SENSITIZED SOLAR CELLS
NOVI ELEKTROLITNI MATERIAL NA OSNOVI TRDNEGA
POLIMERA ZA SON^NE CELICE, OB^UTLJIVE ZA SVETLOBO
Vivek K. Singh, Bhaskar Bhattacharya, Shashank Shukla, Pramod K. Singh
Material research laboratory, School of Basic Sciences & Research, Sharda University, G. Noida 201310, India
vivekv445@gmail.com; singhpk71@gmail.com
Prejem rokopisa – received: 2014-02-08; sprejem za objavo – accepted for publication: 2014-02-26
A solid-polymer electrolyte consisting of polyvinylpyrrolidone (PVP) doped with ammonium iodide (NH4I) was developed and
characterized for a possible application in a dye-sensitized solar cell. Complex impedance spectroscopy revealed an increase in
the conductivity and the maximum conductivity was obtained at the w = 50 % NH4I mass concentration. Light photographs
confirmed an enhancement in the amorphous nature of the host which was affirmed by XRD measurement. The composite
nature of the polymer-electrolyte film was also confirmed with the FTIR spectrum. A dye-sensitized solar cell (DSSC) was
fabricated using the most conductive film that showed an efficiency of 0.025 % at the 1 sun condition.
Keywords: polymer electrolyte, conductivity, FTIR, XRD, dye-sensitized solar cell
Razvit in karakteriziran je bil trdni polimerni elektrolit, ki ga sestavlja polivinil pirolidon (PVP), dopiran z amonijevim iodidom
(NH4I), za morebitno uporabo za son~ne celice, ob~utljive za svetlobo. Kompleksna impedan~na spektroskopija je odkrila
pove~ano prevodnost z maksimumom pri masni koncentraciji w = 50 % NH4I. Posnetki s svetlobno mikroskopijo so odkrili
pove~anje dele`a amorfne osnove, kar so potrdile tudi XRD-meritve. Kompozitno naravo polimerne elektrolitne plasti je potrdil
tudi FTIR-spekter. Son~ne celice, ob~utljive za svetlobo (DSSC), so bile izdelane z uporabo najbolj prevodne plasti, ki je
pokazala u~inkovitost 0,025 % v razmerah 1 sun.
Klju~ne besede: polimerni elektrolit, prevodnost, FTIR, XRD, za svetlobo ob~utljive son~ne celice
1 INTRODUCTION
Polymer electrolytes are promising candidates for
electromechanical-device applications chiefly because
they mechanically behave like solids but their internal
structure and, consequently, the conductivity behavior
closely resemble the liquid state.1 The main advantages
of polymeric electrolytes are satisfactory mechanical
properties, easy fabrication of thin films and an ability to
form a good electrode/electrolyte contact.2–5
Polyvinylpyrrolidone (PVP), also commonly called
polyvidone or povidone, is a polymer made of N-vinyl-
pyrrolidone monomer. PVP was first synthesized by
Prof. Walter Reppe and a patent was filed in 1939. When
dry it is a light flaky powder, which readily absorbs up to
40 % of its weight in atmospheric water. PVP is soluble
in water and other polar solvents. Since it has excellent
wetting properties and readily forms films, it makes a
good coating or an additive to coatings. PVP is used in a
wide variety of applications in medicine, pharmacy, cos-
metics and industrial production. When added to iodine,
PVP forms a complex called povidone-iodine exhibiting
disinfectant properties and being beneficial for dye-sen-
sitized solar-cell applications where iodide/polyiodide
redox couple is frequently added to the electrolyte.
Dye-sensitized solar cells (DSSCs) were first reported by
O’Regan and Grätzel in 1991.6
Over the past decade, DSSCs have been intensely
investigated as potential alternatives to the conventional
inorganic photovoltaic devices due to their low produc-
tion cost and good efficiency for a conversion of solar
energy into electricity. A typical cell consists of a nano-
crystalline mesoporous titanium dioxide film sensitized
with a monolayer dye, an electrolyte containing
iodide/triiodide as the redox couple and a platinum
counter electrode. Liquid electrolytes were replaced with
solid-polymer electrolytes because the former lead to
corrosion, evaporation and leakage. Thus, the solid-poly-
mer electrolytes improved the long-term stability of
DSSCs.
In the present paper, we report on new solid-polymer
electrolyte films of a polyvinylpyrrolidone (PVP) com-
plex with ammonium iodide (NH4I) and a DSSC that was
fabricated using the film with the maximum electrical
conductivity.
2 MATERIALS AND METHOD
Polyvinylpyrrolidone (PVP, Mw = 130,000), ammo-
nium iodide (NH4I) and iodine (I2) were purchased from
Sigma-Aldrich, USA while methanol was purchased
from Qualikems Fine Chem. Pvt. Ltd., Vadodara, India.
The following approach was taken to prepare the
electrolytes. PVP (500 mg) was dissolved in about 4 mL
of methanol under continuous magnetic stirring ( 30
min) or until complete dissolution at room temperature.
Then an appropriate quantity of NH4I salts was added to
Materiali in tehnologije / Materials and technology 49 (2015) 1, 123–127 123
UDK 678.7:544.6.018.47-036.5 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)123(2015)
the PVP methanolic solution and stirred continuously.
After the solvent evaporation the polymer-salt complex
was poured into polypropylene Petri dishes. Free-stand-
ing films of different PVP compositions – w(NH4I)/%
(where w = (10, 20, 30, 40, 50, 60, 70) %) – were ob-
tained and further characterized using various characteri-
zation tools.
A dye-sensitized solar cell with an active area of 0.72
cm2 was fabricated with the procedure reported else-
where.7 With the common procedure the TiO2 paste was
applied on the fluorine-doped SnO2 substrate (FTO)
using the doctor-blade method. The adhesive scotch tape
was used to control the thickness of the as-coated TiO2
film with a thickness of  50 μm, followed by a heat
treatment at 500 °C for 30 min. The porous TiO2 film
formed on the FTO substrate was 10 μm thick and a pore
had a diameter of 10–15 nm.8–13 The porous TiO2 film on
the FTO substrate was then immersed in a ruthenium
sensitizer dye solution (0.5 mmol N-719, Solaronix, in
ethanol) and left overnight to allow a sufficient dye
adsorption. This TiO2 working electrode with the dye
was then rinsed off with distilled water and ethanol solu-
tion. A Pt-thin-film-coated counter electrode was pre-
pared separately by spin-coating the H2PtCl6 solution
onto the FTO substrate. The viscous polymer-electrolyte
solution ( 400 μL) containing PVP:NH4I + I2 (the
maximum ) was finally cast on the working electrode (a
two-step casting) and sandwiched between the platinized
counter electrode and the TiO2 working electrode.
3 RESULTS AND DISCUSSIONS
3.1 Conductivity measurement
Ionic conductivity of the polyvinylpyrrolidone-based
polymer-electrolyte film was measured using a CH
Instruments workstation (model 604D, USA) over a fre-
quency range of 100–105 Hz.
We used steel electrodes as contacts to measure the
ionic conductivity () and we calculated ionic-conduc-
tivity values using the following formula:
 = G · l/A (1)
where  is the ionic conductivity, G is the conductance
(in the case of 1/Rb, Rb is the bulk resistance where the
Nyquist plot intercepts with the real axis), l is the
thickness of the sample and A is the area of the given
sample.
The cole-cole plot (complex impedance plot) of a
typical sample of the PVP + w(NH4I) 50 % polymer
electrolyte is shown in Figure 1. The room-temperature
ionic conductivity (deduced from different cole-cole
plots) of polymer electrolytes as a function of the NH4I
concentration is shown in Figure 2 and its values are
listed in Table 1.
Table 1: Room-temperature ionic conductivity of the PVP:NH4I
polymer electrolyte system
Tabela 1: Ionska prevodnost PVP:NH4I polimernega elektrolitskega
sistema pri sobni temperaturi
Composition (w(NH4I)/%) Conductivity (S cm–1)
10
20
30
40
50
60
70
2.24 × 10–5
2.57 × 10–5
2.63 × 10–5
1.09 × 10–4
7.55 × 10–4
4.85 × 10–4
7.45 × 10–5
As observed in Figure 2 and Table 1, the ionic
conductivity () increases with the increase in the NH4I
concentration and reaches its maximum at the w(NH4I) =
50 % ( = 7.55 × 10–4 S/cm) concentration and then it
decreases. The increase in the ionic conductivity with the
increasing NH4I concentration can be related to the
increase in the number of mobile charge carriers, while
the possible decrease in the ionic conductivity at a NH4I
mass concentration greater than 50 % can be attributed to
the formation of ion multiples.
V. K. SINGH et al.: NEW SOLID-POLYMER-ELECTROLYTE MATERIAL FOR DYE-SENSITIZED SOLAR CELLS
124 Materiali in tehnologije / Materials and technology 49 (2015) 1, 123–127
Figure 2: Effect of the NH4I amount on the conductivity of the
polymer electrolyte (PVP:NH4I) measured at room temperature
Slika 2: Vpliv vsebnosti NH4I na prevodnost polimernega elektrolita
(PVP:NH4I), izmerjeno pri sobni temperaturi
Figure 1: Cole-cole plot of the PVP + w(NH4I) 50 % polymer-elec-
trolyte system
Slika 1: Cole-cole-diagram polimernega elektrolitnega sistema PVP +
w(NH4I) 50 %
The ionic conductivity () in the case of an electro-
lyte system is given as:
 = n q μ (2)
where n is the charge carrier density, q is the charge of
the carrier and μ is the mobility of the carriers. There-
fore, any increase in either n or q will certainly affect
the value of ionic conductivity.
3.2 X-ray diffraction
The crystallinity of the polymer electrolyte was
further affirmed by X-ray diffraction patterns (XRD)
using a Rigaku D/max-2500 XRD diffractometer at a
scan rate of 5° min. The recorded X-ray diffraction
patterns of pure PVP and NH4I doped PVP polymer
electrolytes are shown in Figure 3. It is clear that pure
PVP shows well-known amorphous peaks around 2
values of 23°. The incorporation of NH4I into the PVP
matrix decreases the intensity of the peaks (the
suppression in crystallinity). It also appears that the
XRD data relating to the NH4I doped PVP polymer
electrolyte shows only the peaks related to either PVP or
NH4I, which clearly affirms the composite nature of the
polymer-electrolyte system. Additionally, the PVP-NH4I
data does not contain any other peaks related to the NH4I
sample, affirming a complete dissolution of NH4I in the
Sago Palm matrix.
3.3 FTIR spectroscopy
The FTIR spectra of pristine PVP, NH4I and the PVP
doped with NH4I were recorded between 4000 cm–1 and
400 cm–1 on a PerkinElmer Spectrophotometer 883 as
shown in Figure 4. Pure NH4I shows well defined peaks
at (3131, 1622 and 1398) cm–1, where the first peak
corresponds to the N-H stretch, while the other two
correspond to the N-H bending. In the spectrum of pure
PVP, the peaks at (847, 895 and 934) cm–1 correspond to
para-, di-substituted and mono-substituted C-H bend-
ings. The bands between 1450 cm–1 and 1600 cm–1
correspond to the C=C stretching. The peaks at 1075
cm–1 and 2135 cm–1 correspond to the C-N stretching,
while the ones at 1018 cm–1 and 1172 cm–1 correspond to
the C-C stretching. The CH3 bending is shown at 1375
cm–1. The peak at 1835 cm–1 corresponds to the presence
of C=O bonds. In the spectrum of PVP + NH4I the peaks
at 843 cm–1 and 934 cm–1 correspond to the C-H bending.
The peaks at 1100–1300 cm–1 correspond to the C-C
stretching. The C-N stretching is shown at 1074 cm–1.
The CH2 bending is given by the peak at 1439 cm–1. The
bands between 1550 cm–1 and 1640 cm–1 correspond to
the N-H bending. The peak at 1848 cm–1 corresponds to
the C=O stretching, while the C=C peaks are indicated
by the peaks between 1450 cm–1 and 1600 cm–1.
It is also clear from the figure that almost all the
peaks related to the host materials (PVP and NH4I) are
present in the NH4I doped PVP polymer-electrolyte
sample. The disappearance of any new peaks other than
those of the host materials clearly affirms the composite
nature of the samples, also supported by our XRD data.
3.4 Light microscopy
Light microscopy (LM) of a polymer-electrolyte
sample with the dimensions of 1 cm × 1 cm was carried
out using a Leica Leitz DMRX light microscope. The
obtained photographs are shown in Figure 5. It is noted
that the pure PVP film (Figure 5a) shows well-ordered
patches, confirming its semicrystalline nature. This
pattern is a bit different from the micrographs of the
PEO matrix. Due to an addition of NH4I to the PVP
matrix (Figure 5b) the patch size becomes random and
the crystallinity seems to be disturbed. The decrease in
the crystallinity (an ordered pattern) showed a further
increase in the amorphicity (a non-ordered pattern)
where different sizes of rough patches are distributed
randomly within the polymer matrix. It is believed that
the amorphous regions (the non-ordered pattern) are
conductivity-rich regions and, hence, our light micro-
V. K. SINGH et al.: NEW SOLID-POLYMER-ELECTROLYTE MATERIAL FOR DYE-SENSITIZED SOLAR CELLS
Materiali in tehnologije / Materials and technology 49 (2015) 1, 123–127 125
Figure 4: FTIR spectra of pure PVP, NH4I and PVP + NH4I polymer
electrolyte
Slika 4: FTIR-spektri ~istega PVP, NH4I in PVP + NH4I polimernega
elektrolita
Figure 3: XRD pattern of pure PVP and PVP + NH4I polymer elec-
trolyte
Slika 3: XRD-posnetek ~istega PVP in PVP + NH4I polimernega elek-
trolita
graphs showed good agreement with the ionic-conducti-
vity data.
To further specify the nature of the charge carriers
(ionic or electronic) we carried out the ionic-transfe-
rence-number measurement. Figure 6 shows the
current-versus-time measurement for a typical sample of
the arrowroot-60 % NH4I polymer-electrolyte matrix. In
this study we applied a fixed DC voltage and the current
was recorded with respect to time following a well-
established formula:
t
I I
Iion
initial final
initial
=
−
(3)
where Iinitial is the initial current and Ifinal is the final
residual current.
The observed ionic-transference number is 0.93
showing that our biopolymer electrolyte is essentially an
ion-conducting system.10
3.5 Ion-transport mechanism
The ion-transport mechanism in the PVP polymer-
electrolyte matrix can be easily understood using Figure
7. According to the literature, in most of the polyethers
incorporated with alkali halides the anion contribution is
more dominant.14 However, in the case of the polyethers
doped with the other salts like NH4ClO4 the cationic part
is more dominating15,16. The NH4+ ions of NH4I are
coordinated with the ether oxygen of PVP and I– anions
hang outside. The weakly bonded H in NH4+ can be
easily dissociated under the influence of a DC electric
field forming H+ ions. These H+ or NH4+ ions can jump
via each coordinating site as shown in Figure 7.
3.6 DSSC performance
A dye-sensitized solar cell (DSSC) was prepared
with the PVP:NH4I/I2 polymer electrolyte with the
maximum ionic conductivity (w(NH4I) = 50 %). Iodine
was also added to prepare the redox couple (10 % with
respect to the iodide salt). The recorded J-V characte-
ristic is shown in Figure 8. The fabricated DSSC shows
V. K. SINGH et al.: NEW SOLID-POLYMER-ELECTROLYTE MATERIAL FOR DYE-SENSITIZED SOLAR CELLS
126 Materiali in tehnologije / Materials and technology 49 (2015) 1, 123–127
Figure 7: Mechanism of ion transport in the PVP:NH4I polymer-elec-
trolyte matrix
Slika 7: Mehanizem potovanja ionov v osnovi polimernega elektrolita
PVP:NH4I
Figure 5: Light microscope photographs of: a) pure PVP, b) PVP + 40
% NH4I polymer-electrolyte matrix
Slika 5: Posnetka s svetlobnim mikroskopom: a) ~isti PVP, b) PVP +
40 % NH4I osnova polimernega elektrolita
Figure 8: Current density versus voltage characteristic of the PVP:50
% NH4I/I2 polymer-electrolyte film at the 1 sun condition
Slika 8: Gostota toka v odvisnosti od zna~ilnosti napetosti v plasti
polimernega elektrolita PVP:50 % NH4I/I2 v razmerah 1 sun
Figure 6: Current-versus-time plot (tion measurement) of a typical
PVP:NH4I polymer-electrolyte matrix measured at room temperature
Slika 6: Odvisnost toka od ~asa (tion meritev) zna~ilne PVP:NH4I
osnove polimernega elektrolita, izmerjena pri sobni temperaturi
an open-circuit voltage (Voc) of 0.35 V and a short-cir-
cuit-current density (Jsc) of 0.1 mA/cm2 with the overall
efficiency of 0.025 %.
The observed efficiency was much lower when
compared to the other polymeric systems reported in11–13.
This was expected since the observed conductivity in the
present case was much lower (by  1–2 order of mag-
nitude).
4 CONCLUSION
A solid polymer-electrolyte film consisting of PVP
doped with the NH4I salt was successfully prepared and
characterized. It was observed that NH4I doping
enhances the ionic conductivity and the conductivity
maximum was obtained at the 50 % NH4I salt mass
concentration with the conductivity value of 7.55 × 10–4
S/cm. XRD and IR affirmed the composite nature of the
polymer electrolyte, while optical micrograph and XRD
affirmed the reduction in the crystallinity due to the NH4I
doping. Using a solid polymer electrolyte with the
maximum ionic conductivity, we fabricated a DSSC that
shows a reasonable photo response at the 1 sun con-
dition.
Acknowledgments
This work was supported by a DST project (DST/
TSG/PT/2012/51) of the Government of India.
5 REFERENCES
1 D. E. Fenton, J. M. Parker, P. V. Wright, Polymer, 14 (1973), 589
2 K. N. Kumar, T. Sreekanth, M. J. Reddy, U. V. S. Rao, J. Power
Sources, 101 (2001), 130–133
3 T. M. W. J. Winjendra, P. Ekanayake, M. A. K. L. Dissanayake, I.
Albinsson, B. E. Mellander, J. Solid State Electrochem., 14 (2010),
1221–1226
4 R. Chandrasekaran, S. Selladurai, J. Solid State Electrochem., 5
(2001), 355–361
5 X. Li, D. Zhang, X. J. Yin, S. Chen, J. Shi, Z. Sun, S. Huang, J. Solid
State, Electrochem., 15 (2011), 1271–1277
6 B. O’Regan, M. Grätzel, Nature, 353 (1991), 737–740
7 M. Singh, V. K. Singh, K. Surana, N. A. Jadhav, P. K. Singh, B.
Bhattacharya, H. W. Rhee, J. Ind. Eng. Chem., 19 (2013), 819–822
8 A. F. Nogueira, J. R. Durrant, M. A. De Paoli, Advanced Materials,
13 (2001), 826–830
9 M. Saito, S. Fujihara, Energy & Environmental Science, 1 (2008),
280–283
10 A. Saxena, P. K. Singh, B. Bhattacharya, Mater. Tehnol., 47 (2013) 6,
799–802
11 P. Kumar, P. K. Singh, B. Bhattacharya, Ionics, 17 (2011), 721–725
12 P. K. Singh, B. Bhattacharya, R. K. Nagarale, J. Appl. Polym. Sci.,
118 (2010), 2976–2980
13 P. K. Singh, B. Bhattacharya, R. M. Mehra, H. W. Rhee, Current
Applied Physics, 11 (2011), 616–619
14 B. Bhattacharya, R. K. Nagarale, P. K. Singh, High Performance
Polymers, 22 (2010), 498–512
15 S. A. Hashmi, A. Kumar, K. K. Maurya, S. Chandra, J. Phys. D:
Appl. Phys., 23 (1990), 1307–1314
16 A. Sachdeva, R. Singh, P. K. Singh, B. Bhattacharya, Mater. Tehnol.,
47 (2013) 4, 467–471
V. K. SINGH et al.: NEW SOLID-POLYMER-ELECTROLYTE MATERIAL FOR DYE-SENSITIZED SOLAR CELLS
Materiali in tehnologije / Materials and technology 49 (2015) 1, 123–127 127

T. F. KUBATÍK et al.: COMPACTING THE POWDER OF Al-Cr-Mn ALLOY WITH SPS
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Tomá{ Franti{ek Kubatík1, Zdenìk Pala1, Pavel Novák2
1Institute of Plasma Physics AS CR, Za Slovankou 1782/3, 182 00 Prague 8, Czech Republic
2Department of Metals and Corrosion Engineering, Institute of Chemical Technology, Technická 5, 166 28 Prague 6, Czech Republic
kubatik@ipp.cas.cz
Prejem rokopisa – received: 2014-02-11; sprejem za objavo – accepted for publication: 2014-02-26
Rapidly solidified powder AlCr6Fe2(Mn, Si, Ti) obtained with the melt atomization contained a solid solution of the alloying
elements in aluminium and Al57Mn12, Al42Cr7(Fe) and Al2Cr3(Fe) intermetallic phases. The addition of chromium to aluminium
alloys is acknowledged to efficiently inhibit the growth of refined Al grains. Spark-plasma sintering (SPS) is a modern method
where the diffusion and decomposition of metastable intermetallics are limited due to the high speed of the process and lower
sintering temperatures. This work investigates a compaction of Al-Cr-Fe-Si alloy powders using this method.
Keywords: aluminium alloy, intermetallics, powder metalurgy, spark-plasma sintering
Hitro strjen prah AlCr6Fe2(Mn, Si, Ti), dobljen z atomizacijo taline, je vseboval trdno raztopino legirnih elementov v aluminiju
in intermetalne faze Al57Mn12, Al42Cr7(Fe), Al2Cr3(Fe). Dodatek kroma aluminiju u~inkovito zadr`uje rast drobnih zrn Al.
Sintranje z iskrasto plazmo (SPS) je sodobna metoda, kjer je omejena difuzija in razgradnja metastabilnih intermetalnih zlitin
zaradi velike hitrosti procesa in ni`jih temperatur sintranja. V tem delu se raziskuje kompaktiranje prahov zlitine Al-Cr-Fe-Si po
metodi SPS.
Klju~ne besede: aluminijeva zlitina, intermetalne faze, pra{na metalurgija, sintranje z iskrasto plazmo
1 INTRODUCTION
Aluminium alloys with transition metals (e.g., Fe, Cr)
have been known as attractive materials with an excellent
combination of the tensile strength and thermal stability,1
which makes them useful for the applications at elevated
temperatures in the automotive and aerospace industries.1
Their thermal stability is attributed to both the fact that
these alloying elements have a low diffusion coefficient
and their solubility in aluminium. However, these feat-
ures make these alloys difficult to prepare using conven-
tional metallurgy with a large amount of alloying
elements. Therefore, these materials are made with pow-
der metallurgy (PM) using rapidly solidified powders. A
general feature of rapidly solidified Al-Cr and Al-Mn
based alloys is the presence of many intermetallic pha-
ses, including quasi crystals (Q-phases),2,3 in the intensi-
vely cooled alloys.
In the literature, the presence of multiple species of
icosahedral quasicrystal phases Al95Fe4Cr (Ti, Si),
Al84.6Cr15.4(Ti, Si) and Al74Cr20Si6(Fe, Ti) was described.
The authors stated that in a powder fraction of 25–45
μm, phases Al13Cr2, Al82Fe18, Al95Fe4Cr and Al74Cr20Si6
were present, while in a fraction of 100–125 μm only
phases Al13Cr2, Al32Fe4 and Al82Fe18 were found.4–6 A
quasi-crystalline phase is decomposed at the tempera-
tures around 550 °C.4,5 It was reported6 that the compact
samples prepared with the methods of cold extrusion,
ultra-high pressure (6 GPa) and hot extrusion (480 °C/30
min preheating) had almost no porosity.
The recently developed field-assisted sintering
(FAST),7,8 also known as spark-plasma sintering
(SPS),9–11 or pulsed-electric-current sintering (PECS),12,13
is a non-conventional powder-consolidation technique
used for a very fast densification of ceramic and metal
powders. In this method, a powder is placed inside a tool
between two pressing punches acting also as electrodes.
High-amperage electric-current pulses are passed
through the uniaxially pressed powder compact loaded in
the tool. The powder is rapidly densified due to a simul-
taneous application of pressure and heat in accordance
with the Joule effect, whereas the grain coarsening is
suppressed and the diffusion and decomposition of
metastable intermetallics are limited due to the high
speed of the process and lower sintering temperatures.
The purpose of the present work is to study the compact-
ing of an aluminium-alloy powder using SPS, monitor
the changes taking place during the compacting of
rapidly solidified powders and compare the structure and
phase composition of the obtained samples with the
conventional method of hot extrusion.
2 EXPERIMENTAL WORK
The feedstock material for SPS was an atomized (N2)
powder with a composition of AlCr6Fe2(Mn, Si, Ti).
The powder was sieved into size fractions of  32 μm
and 160–80 μm. The compacting was carried out in a
spark-plasma sintering (SPS) system – model SPS 10-4
(Thermal Technology, USA). The tool system placed
inside the vacuum chamber consisted of a graphite
punch-and-die unit. The powder was manually poured
into the die cavity. The temperature was measured with a
NiCr-Ni thermocouple.
Materiali in tehnologije / Materials and technology 49 (2015) 1, 129–132 129
UDK 669.715:621.762:621.762.5 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)129(2015)
Figure 1 is a standard record of a single cycle. The
program begins with the evacuation of the chamber.
Afterwards, the sample is pre-loaded at 20 MPa,
followed by heating with the speed of 100 °C/min, using
a pulsed DC source. After reaching the desired process
temperature (450, 500 and 550) °C, the sample is loaded
to the final compression pressure of 80 MPa. The dwell
time at this temperature and pressure is 5 min, as shown
in Figure 1.
After the dwell time, the DC pulse source is turned
off and the sample is cooled down to room temperature
with free cooling. After the aeration of the chamber, the
graphite tool is opened and the sample is removed. The
prepared samples are cylindrical with a diameter of 19
mm and height of 4–6 mm. The porosity of the samples
was measured with the Archimedes method. The micro-
structure was observed in the metallographic cut using a
scanning electron microscope EVO MA 15 (Carl Zeiss
SMT, Germany, SEM). Both the starting powders, i.e.,
with the particle sizes smaller than 32 μm and within the
range from 160 μm to 180 μm, and the obtained spark-
plasma-sintered samples were subjected to X-ray diffrac-
tion (XRD) analyses.
The powders were inserted into standard sample hol-
ders with the so-called side-loading procedure that
minimizes the rise of the preferred orientation. These
two samples were measured in the standard Bragg-Bren-
tano geometry with a divergent beam and with a beam
knife placed above the samples in order to minimize the
effects of air scattering. Since our aim was to further
measure the centres of the sintered samples, the irra-
diated volumes were in the middle of the samples’
cross-sections. Therefore, a poly-capillary system and a
collimator of 1 mm in diameter were inserted into the
primary X-ray beam path changing the originally diver-
gent character into a quasi-parallel one; diffraction CuK

radiation was detected with a 1D LynxEye detector. The
used D8 Discover diffractometer was equipped with a
laser system and a compact x, y, z stage facilitating
precise positioning of the measured surface.
3 RESULTS AND DISCUSSION
3.1 Densification behaviour
The diagram in Figure 1 shows the process parame-
ters during the SPS process at 550 °C for the coarse
fraction (160–180 μm). The plots for all the temperatures
and fractions are very similar differing only at low
temperatures, where a small change in the position can
be seen even during the dwell time. The curve of the
position, which shows a change in the position of the
pistons, indicates that the increasing temperature slightly
increases the value of the positions related to thermal
heating and auxiliary piston-guide graphite parts. At the
temperatures slightly above 300 °C, the softening and
deformation of the alloy powder occur (sintering). This
temperature is the same for any sintering temperature
(450, 500 and 550) °C and powder fraction (32 μm and
160–180 μm).
After reaching the temperature, the pressure increases
for one minute, reaching the resulting value of 80 MPa.
During the pressure increase, a very sharp reduction in
the value of the position is detected, indicating an inten-
sive compaction. After reaching the final pressure, the
current values of the temperature and pressure are main-
tained constant for the duration of 5 min. The sintering
temperature of 450 °C leads to a gradual compaction,
during the dwell time, of the fraction of the powder
smaller than 32 μm but to a smaller extent than for the
coarse fraction. This may be related to the fact that the
spherical powder alloy has a hard thin oxide layer, whose
proportion is smaller in the coarse fraction of the alu-
minium alloy. This hypothesis is confirmed with the
results of the open-porosity measurements, whose values
are listed in Table 1. The results also show that the
sintering temperature of 550 °C results in a very low
open porosity, which is almost at the limit of the
quantification obtained with the Archimedes method.
Table 1: Open porosity depending on the sintering temperature and
powder size
Tabela 1: Odprta poroznost je odvisna od temperature sintranja in od
velikosti zrn prahu
Powder
fraction
(μm)
Temperature
of sintering
(°C)
Open
porosity
(%)
Aluminium alloy 32 450 3.05
Aluminium alloy 32 500 2.63
Aluminium alloy 32 550 0.14
Aluminium alloy 160–180 450 2.88
Aluminium alloy 160–180 500 2.88
Aluminium alloy 160–180 550 0.068
3.2 Microstructure
Figure 2 is a cross-sectional microstructure of the
sintered powder fraction of 160–180 μm at 550 °C. The
figure shows that the sintered alloy does not contain any
visible pores and has a nearly homogeneous structure. In
T. F. KUBATÍK et al.: COMPACTING THE POWDER OF Al-Cr-Mn ALLOY WITH SPS
130 Materiali in tehnologije / Materials and technology 49 (2015) 1, 129–132
Figure 1: Parameter progress during the sintering of aluminium alloy
powder (160–180 μm fraction) at 550 °C (SPS)
Slika 1: Spreminjanje parametrov med sintranjem prahu aluminijeve
zlitine (zrnatost 160–180 μm) pri 550 °C (SPS)
the microstructure, only some spherical particles can be
seen. The structure consists of an aluminium matrix with
a small amount of silicon (titanium) in the solid solution,
in which Al45Cr7 intermetallic phases consisting of
rounded particles with an irregular shape are clearly
visible. The observed microstructures are fairly similar at
all the compacting temperatures and for both powder
fractions.
3.3 XRD – phase analysis
The obtained diffraction patterns are juxtaposed in
Figure 3. The phase identification revealed, apart from
the expected fcc (Fm3m) aluminium solid solution, also
aluminium manganese (cubic Pm-3 Al57Mn12) and alumi-
nium chromium (monoclinic C12/m1 Al45Cr7) phases. It
is clearly seen that the differences between all the SPS-
sample diffraction patterns are virtually non-existent and
that the peak positions match those of the coarser
powder. However, several differences between the
diffraction patterns of both powders can be observed
(Figure 4) and, hence, the sieving led to minor changes
in the phase composition.
The phase ID of the powder with smaller particles
gave hints that the Al-Cr-Mn phase could be present.
There are only two such phases in the ICSD database,
namely, cubic Im-3 Al12Cr0.5Mn0.5 and tetragonal
I4/mmm Al2Cr3Mn. The latter proved to be the one
present according to the Rietveld refinement, which gave
the quantity of almost w = 1 % of this tetragonal phase.
The results of the quantitative Rietveld refinement of
both powders are summarized in Table 2. Since the
diffraction patterns of the sintered samples do not exhibit
significant differences, similar information about the
phase presence is not relevant and, thus, not provided.
Both the Al-Cr (ICSD code of 57652) and Al-Mn (ICSD
code of 14374) phases were used for the Rietveld
refinement performed in TOPAS 4.2 with sufficiently
good matches between the structure model and measured
data. There is, however, about the mass fraction 3 %
more of the Al-Cr phase in the sintered samples due to
various reasons; the difference in the irradiated volume
can be one of them.
Table 2: Results of a quantitative Rietveld refinement of both powders
Tabela 2: Rezultati kvantitativnega Rietveldovega udrobljenja obeh
prahov
Powder parti-
cle fraction
w(Al)
/%
w(Al57Mn12)
/%
w(Al42Cr7)
/%
w(Al2Cr3Mn)
/%
< 32 μm 63.9 19.0 16.3 0.8
160–180 μm 68.0 9.4 22.6
mass fraction, w/%
4 CONCLUSIONS
Using the SPS method, a compact alloy with a very
low open porosity was prepared. The porosity values are
dependent on the powder fraction and sintering tempe-
rature. The microstructure of the compacted samples is
formed using an aluminium-based solid solution with
homogeneously dispersed intermetallic particles Al45Cr7
at all the sintering temperatures. The performed eva-
T. F. KUBATÍK et al.: COMPACTING THE POWDER OF Al-Cr-Mn ALLOY WITH SPS
Materiali in tehnologije / Materials and technology 49 (2015) 1, 129–132 131
Figure 4: Differences between diffraction-peak positions for the two
fractions of the used powder
Slika 4: Razlika med pozicijo difrakcijskih vrhov za dve frakciji
uporabljenih prahov
Figure 2: Microstructure of the alloy sintered at 550 °C for 5 min
(fraction of 160–180 μm, SEM)
Slika 2: SEM-posnetek mikrostrukture zlitine, sintrane 5 min na 550
°C (zrnatost 160–180 μm)
Figure 3: Plain comparison of the obtained diffraction patterns. The
substantial difference between the intensities is due to the differences
in the irradiated volume of the powders and the centres of the SPS
sample cross-sections.
Slika 3: Primerjava dobljenih difrakcijskih vzorcev. Ob~utna razlika
v intenzitetah je zaradi razlik v obsevanem volumnu prahov in sredine
prereza SPS-vzorcev.
luation of diffraction data indicates that all the sintered
samples are very similar with respect to the phase
composition. Moreover, the shapes of diffraction profiles
are also the same which means that the real structures,
namely, the microstrains and coherent domain sizes,
remain virtually the same for all the investigated sam-
ples.
It should be noted that the irradiated volumes were in
the centre of the cross-sections and, hence, certain diffe-
rences could occur on the sample surfaces. Nevertheless,
we unfolded comparatively significant differences bet-
ween the phase compositions of both powder fractions.
The powder with smaller particles includes almost mass
fractions 1 % of the tetragonal Al-Cr-Mn phase, 10 %
more of the Al-Mn phase and about 6 % less of the
Al-Cr phase. The probable reason for the differences lies
in the fact that the powder fractions differ in the cooling
rate. Finer powders solidify more rapidly than the
coarser ones. This difference can cause changes to the
types of intermetallics that are formed. Moreover, the
identified phase compositions include more manganese
than expected on the basis of the initial chemical com-
position of the atomized powder. This discrepancy can
be explained with, e.g., the substitution of manganese
atoms by iron.
The SPS method is a very effective method for com-
pacting alumina-based powders. The sample size is
limited by the size of the die. In this case, the whole
sample body is heated in accordance with the Joule
effect, which enables a production of the samples with
almost no porosity at very short dwell times, unlike in
the case of hot extrusion. The results of the XRD anal-
ysis show that there are almost no observable changes
during the SPS compacting process with respect to the
phase composition and real structure.
5 REFERENCES
1 M. J. F. Gándara, Aluminium: the metal of choice, Mater. Tehnol., 47
(2013) 3, 261–265
2 M. Yamasaki, Y. Nagaishi, Y. Kawamure, Inhibition of Al grain
coarsening by quasicrystalline icosahedral phase in the rapidly
solidified powder metallurgy Al-Fe-Cr-Ti alloy, Scripta Materialia,
56 (2007), 785–788
3 G. Lojen, T. Bon~ina, F. Zupani~, Thermal stability of Al-Mn-Be
Melt-spun ribbons, Mater. Tehnol., 46 (2012) 4, 329–333
4 B. Bártová, D. Vojtìch, J. Verner, A. Gemperle, V. Studni~ka, Struc-
ture and properties of rapidly solidified Al-Cr-Fe-Ti-Si powder
alloys, Journal of Alloys and Compounds, 387 (2005), 193–200
5 D. Vojtìch, A. Michalcová, J. Pilch, P. [ittner, J. [erák, P. Novák,
Structure and thermal stability of Al-5.7Cr-2.5Fe-1.3Ti alloy
produced by powder metallurgy, Journal of Alloys and Compounds,
475 (2009), 151–156
6 D. Vojtìch, A. Michalcová, F. Prù{a, K. Dám, P. [edá, Properties of
the thermally stable Al95Cr3.1Fe0.8 alloy prepared by cold-com-
pression at ultra-high pressure and by hot-extrusion, Materials
characterization, 66 (2012), 83–92
7 Y. Minamino, Y. Koizumi, M. Tsuji, N. Hirohata, K. Mizuuchi, Y.
Ohkanda, Microstructures and mechanical properties of bulk
nanocrystalline Fe-Al-C alloys made by mechanical alloying with
subsequent spark plasma sintering, Science and Technology of
Advanced Materials, 5 (2004), 133–143
8 T. Masao, Handbook of Advanced Ceramics, Chapter 11.2.3 – Spark
Plasma Sintering (SPS) Method, System and Applications, Second
Edition, 2013, 1149–1177
9 T. Skiba, P. Hau{ild, M. Karlík, K. Vanmeensel, J. T. Vleugels,
Mechanical properties of spark plasma sintered FeAl intermetallics,
Intermetallics, 18 (2010), 1410–1414
10 K. Sastry, Field assisted sintering and characterization of ultrafine
and nanostructured aluminium alloys, PhD thesis, Ku Leuven, 2006,
p. 255
11 K. Vanmeensel, A. Laptev, J. Hennicke, J. Vleugels, O. Van Der
Biest, Modelling of the temperature distribution during field assisted
sintering, Acta Materialia, 53 (2005), 4379–4388
12 J. Zhang, A. Zavaliangos, Jr. J. Groza, Field activated sintering
techniques: a comparison and contrast, Powder Metallurgy Science
Technology and Application, 5 (2003), 17–21
13 R. Orrú, R. Licheri, A. Locci, A. Cincotti, G. Cao, Consolidation
synthesis of materials by electric current activated/assisted sintering,
Materials Science and Engineering R, 63 (2009), 127–287
T. F. KUBATÍK et al.: COMPACTING THE POWDER OF Al-Cr-Mn ALLOY WITH SPS
132 Materiali in tehnologije / Materials and technology 49 (2015) 1, 129–132
F. NIKOOSOHBAT et al.: EFFECT OF TEMPERING ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
EFFECT OF TEMPERING ON THE MICROSTRUCTURE AND
MECHANICAL PROPERTIES OF RESISTANCE-SPOT-WELDED
DP980 DUAL-PHASE STEEL
VPLIV POPU[^ANJA NA MIKROSTRUKTURO IN MEHANSKE
LASTNOSTI TO^KASTO VARJENEGA DVOFAZNEGA JEKLA
DP980
Fateme Nikoosohbat1, Shahram Kheirandish2, Masoud Goodarzi2,
Majid Pouranvari3
1West Coast Stainless, California, USA
2Materials Science and Metallurgical Engineering Department, Iran University of Science and Technology (IUST), Tehran, Iran
3Materials and Metallurgical Engineering Department, Dezful Branch, Islamic Azad University, Dezful, Iran
mpouranvari@yahoo.com,
Prejem rokopisa – received: 2014-02-12; sprejem za objavo – accepted for publication: 2014-03-12
This paper aims at investigating the effect of an in-process tempering cycle on the microstructure and mechanical properties of
the DP980 dual-phase-steel resistance spot welds. A microstructural characterization, a microhardness test and a quasi-static
tensile-shear test were conducted. The results showed that the fusion-zone microstructure of as-welded spot welds mainly
consisted of lath martensite. A tempered martensite region was observed in the sub-critical heat-affected zone, producing a
relatively soft region compared to the base metal. It was found that applying an in-process tempering for a sufficient time
promoted the heat-affected-zone softening and produced a softer fusion zone due to martensite tempering which, in turn,
improved the energy-absorption capability of spot welds.
Keywords: resistance spot welding, dual-phase steel, temper, microstructure
^lanek opisuje vpliv cikla popu{~anja med postopkom to~kastega varjenja dvofaznega jekla DP980 na mikrostrukturo in
mehanske lastnosti to~kastega zvara. Izvr{ena je bila karakterizacija mikrostrukture, izmerjena mikrotrdota in opravljen kvazi
stati~ni natezno-stri`ni preizkus. Rezultati so pokazali, da je mikrostruktura podro~ja zlivanja to~kastega zvara prete`no iz
latastega martenzita. V obmo~ju podkriti~ne toplotno vplivane cone je bilo opa`eno podro~je popu{~enega martenzita, ki je
relativno mehko podro~je v primerjavi z osnovnim materialom. Ugotovljeno je bilo, da uporaba dovolj dolgega popu{~anja med
postopkom pospe{uje meh~anje toplotno vplivane cone in nastanek mehkej{e cone zlivanja zaradi popu{~nega martenzita, kar
hkrati izbolj{a zmo`nost absorpcije energije to~kastega zvara.
Klju~ne besede: uporovno to~kasto varjenje, dvofazno jeklo, popu{~anje, mikrostruktura
1 INTRODUCTION
Dual-phase (DP) steels exhibit a unique microstruc-
ture consisting of soft ferrite and hard martensite that
offers a favorable combination of strength, high work-
hardening rate, ductility and formability. Due to these
features, automotive companies are finding that the use
of these steels can enable them not only to reduce the
overall weight of an automobile, but also providing an
improved crash protection to the vehicle occupants.1–3
Resistance spot welding (RSW) is the main joining
process of sheet metal, particularly in the automotive
industry. Vehicle crashworthiness, which is defined as
the capability of a car structure to provide an adequate
protection to its passengers against injuries in the event
of a crash, largely depends on the spot-weld structural-
mechanical behavior.4–7 Considering the development
and commercialization of new DP steels for the applica-
tion in automotive bodies, there is an increasing need to
study the spot-welding behavior of these materials.
According to the literature, the problems associated with
the resistance spot welding of DP steels can be summa-
rized as follows:8–15
(i) High susceptibility to failure in the interfacial failure
mode
(ii) High susceptibility to expulsion
(iii) Sensitivity to the formation of shrinkage voids
(iv) High hardness of the fusion zone due to martensite
formation. This can have adverse effects on the fail-
ure mode under certain loading conditions (e. g., a
peel test).
Very high cooling rates in RSW cause a formation of
martensite even in low-carbon steels. Tempering can
then be performed by adding a post-weld current pulse
(the in-process tempering step) to the welding sequence.
This tempering step is a relatively simple way of reduc-
ing the weld hardness which, in turn, lowers the brittle-
ness of the weld in C-Mn steels.16 After the weld nugget
has been formed, it is held between the electrodes long
enough to sufficiently quench to martensite. A subse-
quent temper pulse is then applied to soften the micro-
structure of the weld nugget. Steels of varying compo-
sition and processing react differently to tempering and it
is unknown how some of the newer transformation-
hardened steels would react during such tempering.16
Materiali in tehnologije / Materials and technology 49 (2015) 1, 133–138 133
UDK 621.791.76/.79:539.4.016:621.785.72 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)133(2015)
The objective of the present research is to investigate
the effect of a tempering cycle on the microstructure,
strength and failure energy of DP980 resistance spot
welds.
2 EXPERIMENTAL PROCEDURE
A DP980-steel sheet 2 mm thick was used as the base
metal. The chemical composition and mechanical pro-
perties of DP980 steel are given in Tables 1 and 2,
respectively.
Table 1: Chemical composition of the investigated steel in mass frac-
tions (w/%)
Tabela 1: Kemijska sestava preiskovanega jekla v masnih dele`ih
(w/%)
P S Al Cu Cr Si Mn C
0.004 0.014 0.036 0.196 0.2 0.13 1.326 0.13
Table 2: Tensile properties of the investigated DP980 steel
Tabela 2: Natezna trdnost preiskovanega jekla DP980
Yield stress
(MPa)
Ultimate tensile
stress (MPa)
Total elongation
(%)
532 990 18
Spot welding was performed using a 120 kVA
AC-pedestal-type resistance-spot-welding machine,
controlled by a PLC. The welding was conducted using a
45° truncated-cone RWMA Class-2 electrode with an
8 mm face diameter. The welding and tempering were
performed in accordance with the schedule given in
Table 3.
Table 3: Welding and tempering schedule
Tabela 3: Potek varjenja in popu{~anja
Squeeze
time
(cycle)
Welding
current
(kA)
Welding
time
(cycle)
Elec-
trode
force
(kN)
Holding/
Cooling
time
(cycle)
Tem-
pering
current
(kA)
Temper-
ing time
(cycle)
40 9 35 5 40 5 20, 40
The quasi-static tensile-shear test samples were pre-
pared according to ANSI/AWS/SAE/D8.9-97 standard.17
Figure 1 shows the sample dimensions used for the ten-
sile-shear tests. The tensile-shear tests were performed at
a cross head of 2 mm/min with an Instron universal
testing machine, recording the load-displacement curves.
The peak load (measured as the peak point in the
load-displacement curve) and the failure energy
(measured as the area under the load-displacement curve
up to the peak load) were extracted from the load-dis-
placement curve.
Light microscopy was used to examine the macro-
structures and to measure the weld-fusion-zone (the
weld-nugget) size. The microstructures of the samples
were also examined with a scanning electron microscope
(SEM). The microhardness test, a technique that has
proven to be useful in quantifying microstructure/mecha-
nical property relationships, was used to determine the
hardness profile parallel to the sheet interface (20 μm
away from the weld centerline), using a 100 g load on a
Bohler microhardness tester.
3 RESULTS AND DISCUSSION
3.1 Microstructure and hardness profile of as-welded
DP980 RSWs
The microstructure and hardness profile of resistance
spot welds play important roles in their failure behavior.
The rapid heating and cooling induced by resistance-
spot-welding thermal cycles significantly alter the micro-
structure in the joint zone. A typical macrostructure of a
F. NIKOOSOHBAT et al.: EFFECT OF TEMPERING ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
134 Materiali in tehnologije / Materials and technology 49 (2015) 1, 133–138
Figure 2: a) Typical macrostructure of an as-welded DP980 joint and
b) hardness profile of the DP980 spot weld in the as-welded condition
(i.e., tempering time = 0) and after the in-process tempering for the
tempering times of 20 and 40 cycles
Slika 2: a) Zna~ilna makrostruktura zvarjenega spoja DP980 in b)
profil trdote to~kastega zvara DP980 v zvarjenem stanju (~as po-
pu{~anja = 0) in po popu{~anju med postopkom za ~ase popu{~anja od
20 do 40 ciklov
Figure 1: Tensile-shear specimen dimensions
Slika 1: Dimenzije preizku{anca za natezno-stri`ni preizkus
DP980 spot weld is shown in Figure 2a indicating three
distinct zones, namely, the fusion zone (FZ), the heat-
affected zone (HAZ) and the base metal (BM).
The hardness profile of the DP980 RSW is shown in
Figure 2b. The hardness variation across the joint can be
analyzed in terms of the microstructure of the joint. The
DP980 base-metal microstructure, as shown in Figure
3a, consists of dispersed martensite islands embedded in
a ferrite matrix; the corresponding hardness is 300 HV
(Figure 2b). As can be seen in Figure 3b, the FZ
predominantly exhibits lath martensite with the average
hardness of 390 HV (Figure 2b). The formation of a
high-volume fraction of martensite in the FZ explains the
higher hardness of the FZ compared to the BM hardness.
The martensite formation in the FZ is attributed to the
inherently high cooling rate of the resistance-spot-weld-
ing process due to the presence of water-cooled copper
electrodes and their quenching effect as well as the short
welding cycle. Gould et al.18 developed a simple analyti-
cal model predicting the cooling rates of resistance spot
welds. According to this model, the cooling rate for a 2
mm thickness is about 2000 K s–1. For steels, the
required critical cooling rate to achieve martensite in the
microstructure can be estimated using the following
equation:19
lg v = 7.42 – 3.13 w(C) – 0.71 w(Mn) – 0.37 w(Ni) –
– 0.34 w(Cr) – 0.45 w(Mo) (1)
where v is the critical cooling rate (K h–1).
The calculated critical cooling rate for the investi-
gated steel is 280 K s–1. Since the recorded cooling rates
in the FZ are significantly higher than the critical cooling
rates needed for a martensite formation, it is not
surprising that a martensite structure is present in the FZ.
The hardness of the HAZ is higher than the BM hard-
ness due to the formation of non-equilibrium phases. The
microstructure of the HAZ is more heterogeneous than
that of the FZ, as verified by the hardness profile. The
material in the HAZ experiences the peak temperature
and cooling rate, inversely proportional to its distance
from the fusion line. Thus, the high-thermal cycle gra-
dient coupled with the resulted austenite-grain structure
can explain the observed microstructure gradient in the
HAZ. One of the interesting features of the DP980 RSW
hardness profile is the HAZ softening (i.e., the HAZ
hardness reduction with respect to the BM). As can be
seen, the HAZ of DP980 welds experiences an approxi-
mately 50 HV softening. The HAZ softening during the
welding of DP steels was reported in20–22. The location of
the HAZ softening corresponds to the sub-critical HAZ
(i.e., the region in the HAZ that experiences the tem-
peratures below the Ac1 critical temperature). It is well
documented that this phenomenon is due to the temper-
ing of the pre-existing martensite in the sub-critical areas
of the HAZ.22 Figure 3c shows a microstructure of this
region indicating a presence of the tempered martensite.
3.2 Criteria for the selection of tempering parameters
The FZ size of the spot welds made with the welding
parameters from Table 3 is about 8 mm, which is well
above the conventional recommendation of 4 t0.5 and
5 t0.5 (t is the sheet thickness (mm)). However, the spot
welds produced under the presented welding condition
failed in the interfacial-failure mode. As mentioned
above, it is well documented that DP steels are more
prone to fail in the interfacial-failure mode. The
in-process quenching and tempering of welds were used
to improve the mechanical properties of the spot welds
with interfacial failure.
To perform a proper quenching and tempering treat-
ment on a weldment, the following points should be
considered:
(i) The selected quenching time should be such that the
weldment cools to a temperature slightly below the
martensite-finish temperature (Mf) before being
heated for tempering. If this temperature is above Mf,
there can be untransformed austenite left in the FZ
and it can re-decompose into the untempered marten-
F. NIKOOSOHBAT et al.: EFFECT OF TEMPERING ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 133–138 135
Figure 3: SEM micrographs showing the microstructures of: a) base
metal, b) fusion zone, c) HAZ softening region in the as-welded joint
Slika 3: SEM-posnetki mikrostrukture: a) osnovna kovina, b) pod-
ro~je zlivanja, c) HAZ-podro~je zmeh~anja v varjenem spoju
site upon cooling to room temperature after the tem-
pering.
(ii) In this study, to estimate the temperature distribution
during the cooling cycle of spot welding, a simple
analytical model predicting the cooling rates of resi-
stance spot welds developed by Gould et al.16,18 was
used. Based on some simplification assumptions, the
temperature distribution in the through-thickness
direction of a steel sheet after the welding current is
turned off, is expressed as follows:
T T
K
K
t
t t
x
=
+ ⎛⎝
⎜ ⎞
⎠
⎟⎛
⎝
⎜
⎞
⎠
⎟
⎛
⎝
⎜
⎞
⎠
⎟
⎛
⎝
⎜
⎞
⎠
⎟
+
P
E
S
S
E S
1
2
2
1
π
π
cos
2
2
π
⎛
⎝
⎜ ⎞
⎠
⎟⎛
⎝
⎜
⎞
⎠
⎟
⎛
⎝
⎜
⎞
⎠
⎟
K
K
t
t
e t
t
E
S
S
E
S
2

#
$
(2)
where TP is the peak temperature in the spot weld, while
tS and tE are the steel-sheet thickness and the electrode
face diameter, respectively. KS and KE are the thermal
conductivity coefficients of steel and copper, respec-
tively. X is the distance from the weld faying surface to
the electrode face, and 
 is the thermal diffusivity of
steel.
Figure 4 shows the calculated temperature distribu-
tion for different levels of penetration in the workpiece,
ranging from 0 % (the faying surface) to 100 % (the
electrode-sheet surface). The temperatures in the weld at
the beginning of the cooling range (0-cycle cooling in
Figure 3a) suggest a full austenization of the weld
through thickness. However, the temperatures in the
extremities of the weld, approaching the electrode-sheet
interfaces, are substantially cooler, being as low as 600
°C. This, of course, is well below the A3 temperature for
the material studied here, suggesting that the peripheral
areas of the weld do not achieve even a partial re-auste-
nization and, therefore, are not subjected to local marten-
site decomposition. Based on Figure 4, the quenching
time between the welding cycle and the tempering cycle
was selected as 40 cycles. Using a quenching time of 40
cycles allows the weldment to cool roughly below the Mf
temperature, facilitating a transformation into martensite.
3.3 Effect of the tempering cycle on the microstructure
and hardness
Figure 2b shows the effect of the tempering cycle on
the weld hardness. According to the hardness values of
the FZ, the tempering time of 20 cycles is insufficient for
the martensite to decompose. According to Figure 5, the
FZ microstructure of the welds tempered for the temper-
ing time of 20 cycles is still martensitic.
After a tempering time of 40 cycles a weld exhibits a
relatively complex distribution of the hardness (Figure
2b). The macrograph of a weld after the tempering time
of 40 cycles (Figure 6a) shows that the periphery of the
weld nugget is surrounded by a narrow shell. The hard-
ness of the weld-nugget centre is higher than that of the
shell. Indeed, the weld nugget consists of an untempered
fusion zone (UTFZ) and a tempered fusion zone (TFZ).
This is caused by a non-uniform temperature distribution
inherent during the in-situ tempering. The temperature
distribution in the spot weld during the tempering is
roughly parabolic in shape, with the peak temperatures
always on the weld centerline.16 Accordingly, the centre
of the weld nugget can be overheated and re-austenized
during the tempering cycle and, subsequently, retrans-
formed into hard martensite during cooling, as confirmed
by the high hardness of the weld center. In the shell of
the TFZ, the microstructure shows evidence of tempering
as shown in Figure 6b. Moreover, according to Figure
2b, the tempering cycle resulted in a softer HAZ due to a
longer time of tempering the martensite present in the
initial microstructure of the BM.
F. NIKOOSOHBAT et al.: EFFECT OF TEMPERING ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
136 Materiali in tehnologije / Materials and technology 49 (2015) 1, 133–138
Figure 4: Calculated cooling profiles of a spot weld for different le-
vels of penetration based on the Gould’s model16,18
Slika 4: Profili ohlajanja to~kastega zvara za razli~ne nivoje penetra-
cije, izra~unani na osnovi Gouldovega modela16,18
Figure 5: FZ microstructure after 20 cycles of the in-process temper-
ing
Slika 5: Mikrostruktura podro~ja zlivanja (FZ) po 20 ciklih popu{~a-
nja med procesom
3.4 Effect of the tempering cycle on the mechanical
properties
The peak load of the spot welds depends on several
factors including the physical weld attributes (mainly the
FZ size and the indentation depth), the failure mode and
the strength of the failure location. The failure energy of
spot welds, measured as the area under the load-dis-
placement curve up to the peak point, can be expressed
as follows:
Energyabsorption Fdl P l
l
= ∝ ×∫
0
max
max max (3)
where Pmax is the peak load and lmax is the maximum
displacement, corresponding to the peak load. The
maximum displacement (lmax) that represents the duc-
tility of the spot welds depends on the ductility of the
failure location. Therefore, the energy absorption de-
pends on the factors governing the peak load and the
ductility of the failure location.
Figure 7a shows the tensile-shear peak loads of the
welds indicating that the as-welded joint (the untem-
pered condition) exhibits the highest peak load. This can
be related to the tensile-shear test configuration, which
tends to load the welds in shear. Under the shear loading,
any tendency for brittle fracture is minimized.23 There-
fore, the shear strength of the spot welds failed in the IF
mode during the tensile-shear test depends on the
strength/hardness correlation of the FZ. In the as-welded
joint, the fully martensitic microstructure provides the
highest peak load. The tempering which softens the FZ
microstructure leads to a reduction in the strength of the
joints.
Figure 7b shows the failure energy of the welds
indicating a positive influence of the tempering cycle on
the failure energy of the welds. According to Figure 7b,
applying a tempering time of 40 cycles improved the
weld failure energy. This can be related to the improved
ductility of the FZ due to the tempering of the as-welded
martensite which, in turn, enhances the energy-absorp-
tion capability of the weld. It was shown that there is a
direct relationship between the failure energy in the sta-
tic tensile-shear test and the impact tensile-shear test.24
Therefore, it can be concluded that the tempering of spot
welds that improves the weld energy-absorption capa-
bility increases the weld performance reliability against
the impact loads such as crash accidents of cars.
F. NIKOOSOHBAT et al.: EFFECT OF TEMPERING ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 133–138 137
Figure 7: Effect of the tempering cycle on the weld strength and
energy
Slika 7: Vpliv cikla popu{~anja na trdnost in energijo zvara
Figure 6: a) Weld-nugget macrograph after 40-cycle in-process tem-
pering showing tempered fusion zone (TFZ) and untempered fusion
zone (UTFZ), b) tempered martensite in the TFZ
Slika 6: a) Posnetek makrostrukture jedra zvara po 40 ciklih po-
pu{~anja ka`e popu{~eno podro~je zlivanja (TFZ) in nepopu{~eno
podro~je zlivanja (UTFZ), b) popu{~eni martenzit v TFZ
4 CONCLUSIONS
1. The FZ of the as-welded spot weld exhibited a lath
martensite microstructure. A reduction in the hard-
ness (softening) with respect to the BM was observed
in the HAZ due to the tempering of the martensite
present in the initial microstructure of the base metal.
2. A short tempering cycle (i.e., a tempering time of 20
cycles) was insufficient for the lamellar martensite
microstructure in the FZ to decompose. With the
increasing tempering time, the martensite tempering
in the FZ increased, resulting in a reduction of the
average hardness of the FZ compared with that of the
as-welded condition. Moreover, applying a tempering
cycle led to an enlargement of the HAZ softening
region.
3. It was shown that applying the in-process tempering
for a sufficient time generates a pronounced HAZ
softening and a softer fusion zone due to martensite
tempering that, in turn, improves the energy-absorp-
tion capability of the welds.
5 REFERENCES
1 M. Pouranvari, Influence of welding parameters on peak load and
energy absorption of dissimilar resistance spot welds of DP600 and
AISI1008 steels, Can Metall Q, 50 (2001), 381–88
2 P. Movahed, S. Kolahgar, S. P. H. Marashi, M. Pouranvari, N. Parvin,
The effect of intercritical heat treatment temperature on the tensile
properties and work hardening behavior of ferrite–martensite dual
phase steel sheets, Mater Sci Eng A, 518 (2009), 1–6
3 M. Delincea, Y. Brechet, J. D. Emburyc, M. G. D. Geersd, P. J.
Jacquesa, T. Pardoena, Separation of size-dependent strengthening
contributions in fine-grained dual phase steels by nanoindentation,
Acta Mater, 55 (2007), 2337–50
4 M. Pouranvari, E. Ranjbarnoodeh, Dependence of the fracture mode
on the welding variables in the resistance spot welding of ferrite-
martensite DP980 advanced high-strength steel, Mater. Tehnol., 46
(2012) 6, 665–671
5 P. Podrzaj, S. Simoncic, Resistance spot welding control based on
fuzzy logic, Int J Adv Manuf Technol., 52 (2011), 959–967
6 P. Podrzaj, S. Simoncic, Resistance spot welding control based on
the temperature measurement, Sci Technol Weld Join, 18 (2013),
551–557
7 M. Pouranvari, S. M. Mousavizadeh, Failure Mode of M130 Marten-
sitic Steel Resistance-Spot Welds, Mater. Tehnol., 47 (2013) 6,
771–776
8 M. Pouranvari, S. P. H. Marashi, Critical review of automotive steels
spot welding: process, structure and properties, Sci Technol Weld
Join, 18 (2013), 361–403
9 M. M. H. Abadi, M. Pouranvari, Failure-mode transition in resistance
spot welded DP780 advanced high-strength steel: effect of loading
conditions, Mater. Tehnol., 48 (2014) 1, 67–71
10 M. Pouranvari, S. P. H. Marashi, D. S. Safanama, Failure mode
transition in AHSS resistance spot welds, Part II: Experimental
investigation and model validation, Mater Sci Eng A, 528 (2011),
8344–52
11 M. Pouranvari, H. R. Asgari, S. M. Mosavizadeh, P. H. Marashi, M.
Goodarzi, Effect of weld nugget size on overload failure mode of
resistance spot welds, Sci Technol Weld Join, 12 (2007), 217–25
12 M. Pouranvari, S. P. H. Marashi, Failure mode transition in AHSS
resistance spot welds, Part I, Controlling factors, Mater Sci Eng A,
528 (2011), 8337–43
13 M. Pouranvari, Failure mode transition in similar and dissimilar
resistance spot welds of HSLA and low carbon steels, Can Metall Q,
51 (2012), 67–74
14 S. Simoncic, P. Podrzaj, Image-based electrode tip displacement in
resistance spot welding, Meas Sci Technol, 23 (2012) 6, 065401
15 X. Sun, E. V. Stephens, M. A. Khaleel, Effects of fusion zone size
and failure mode on peak load and energy absorption of advanced
high strength steel spot welds under lap shear loading conditions,
Eng Fail Anal, 15 (2008), 356–67
16 W. Chuko, J. Gould, Development of Appropriate Resistance Spot
Welding Practice for Transformation-Hardened Steels, Report to the
American Iron and Steel Institute, 2002
17 Recommended Practices for Test Methods and Evaluation the
Resistance Spot Welding Behavior of Automotive Sheet Steels,
ANSI/AWS/SAE D8.9–97, 1997
18 J. E. Gould, S. P. Khurana, T. Li, Predictions of microstructures
when welding automotive advanced high-strength steels, Weld J, 86
(2006), 111–6
19 K. E. Easterling, Modelling the weld thermal cycle and transfor-
mation behavior in the heat-affected zone, In: H. Cerjak, K. E.
Easterling (Eds.), Mathematical Modelling of Weld Phenomena, 1st
ed., Maney Publishing, 1993
20 F. Nikoosohbat, S. Kheirandish, M. Goodarzi, M. Pouranvari, S. P.
H. Marashi, Microstructure and failure behaviour of resistance spot
welded DP980 dual phase steel, Mater Sci Technol, 26 (2010),
738–44
21 M. Pouranvari, S. P. H. Marashi, Key factors influencing mechanical
performance of dual phase steel resistance spot welds, Sci Technol
Weld Join, 15 (2010), 149–155
22 V. H. B. Hernandez, S. K. Panda, M. L. Kuntz, Y. Zhou, Nanoinden-
tation and microstructure analysis of resistance spot welded dual
phase steel, Mater Lett, 64 (2010), 207–210
23 R. W. Hertzberg, Deformation and fracture mechanics of engineering
materials, 2nd ed, John Wiley & Sons, 1996
24 M. Zhou, H. Zhang, S. J. Hu, Critical Specimen Sizes for Ten-
sile-Shear Testing of Steel Sheets, Weld J, 78 (1999), 305–313
F. NIKOOSOHBAT et al.: EFFECT OF TEMPERING ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
138 Materiali in tehnologije / Materials and technology 49 (2015) 1, 133–138
M. SARIKAYA et al.: OPTIMIZATION OF THE PROCESS PARAMETERS FOR SURFACE ROUGHNESS ...
OPTIMIZATION OF THE PROCESS PARAMETERS FOR SURFACE
ROUGHNESS AND TOOL LIFE IN FACE MILLING USING THE
TAGUCHI ANALYSIS
OPTIMIZACIJA PROCESNIH PARAMETROV GLEDE NA
HRAPAVOST POVR[INE IN TRAJNOSTNO DOBO ORODJA PRI
^ELNEM REZKANJU Z UPORABO TAGUCHIJEVE ANALIZE
Murat Sarýkaya1, Hakan Dilipak2, Akýn Gezgin3
1Department of Mechanical Engineering, Sinop University, 57030 Sinop, Turkey
2Manufacturing Department, Technology Faculty, Gazi University, 06500 Ankara, Turkey
3Mechanical Education, Institute of Science and Technology, Gazi University, 06500 Ankara, Turkey
msarikaya@sinop.edu.tr
Prejem rokopisa – received: 2014-02-13; sprejem za objavo – accepted for publication: 2014-03-12
In this study, the Taguchi method, which is a powerful tool to design quality optimization, is used to find the optimum surface
roughness and tool life in milling operations. An orthogonal array, a signal-to-noise (S/N) ratio, and an analysis of variance
(ANOVA) are employed to investigate the tool life and the surface-roughness characteristics of AISI D3 steel. Accordingly, the
lowest surface roughness and the highest tool life were estimated to be 0.436 μm and 434.1 s, respectively and, finally, the
Taguchi method allowed the optimization of the system for the verification of the tests.
Further, ANOVA analysis was revealed that the number of cutter insert was the most important parameter influencing
the surface roughness with a 75.27 %, and cutting speed was the most important parameter influencing the tool
life with a 95 %.
Keywords: material machinability, face milling, Taguchi method, optimization, experimental design
V tej {tudiji je bila uporabljena Taguchijeva metoda kot mo~no orodje za ugotavljanje razmer za doseganje optimalne hrapavosti
in trajnostne dobe orodja pri rezkanju. Ortogonalna matrika, signal hrupa (S/N) in analiza variance (ANOVA) so bili uporabljeni
za preiskavo zdr`ljivosti orodja in hrapavosti povr{ine jekla AISI D3. Na podlagi rezultatov je bila ocenjena najmanj{a hra-
pavost povr{ine 0,436 μm in najve~ja zdr`ljivost orodja 434,1 s, Taguchijeva metoda pa je omogo~ila optimizacijo sistema pri
verifikaciji preizkusov.
Nadalje je ANOVA analiza odkrila, da je {tevilo rezalnih vlo`kov najbolj pomemben parameter, ki s 75,27 % vpliva na hrapa-
vost povr{ine, hitrost rezanja pa je najpomembnej{i parameter, ki s 95 % vpliva na `ivljenjsko dobo orodja.
Klju~ne besede: obdelovalnost materiala, ~elno rezkanje, Taguchijeva metoda, optimizacija, na~rtovanje preizkusov
1 INTRODUCTION
Basically, milling is one of the most commonly used
chip removal operations in manufacturing processes and
machined parts are usually utilized to assembly with
other parts in aerospace, die, medical, automotive, de-
fense industry and machine design as well as in manu-
facturing industries.1 In addition to the cutting insert, the
tool holder, workpiece material, cutting speed (V), feed
rate (f), depth of cut (a) and number of milling cutting
inserts are the most important cutting parameters that
highly affect the performance characteristics such as tool
life and surface roughness.
Generally, researchers have focused on the tool
deformation, the effects of cutting-tool coatings and
environmental conditions using a single insert. The costs
of cutting tools are important to manufacturers. The cost
factor causes lowering to the minimum level of an
implementation with the least number of inserts. In
contrast, the increase in the duration of the process is
also a well-known factor.2 The aim of this study is to find
out the effects of the cutting parameters (the cutting
speed, the feed rate and the number of cutting inserts) on
the surface roughness and the tool life at high cutting
speeds. The roughness of machined surface is an import-
ant quality indicator in machining processes and the
various properties of machined parts such as corrosion,
wear, friction, and heat transmission are also influenced
by surface roughness.3,4 Most of the process parameters
like spindle speed, feed rate, number of insert, depth of
cut, tool holder geometry, cutting insert geometry, tool
material, cooling condition affect the tool life and
surface roughness. Thus, it is difficult to define a general
model for tool life and surface roughness.5 Some stati-
stical methods like Taguchi, Response Surface Methodo-
logy (RSM), desirability functional analysis, ANOVA
and Grey Relational Analysis (GRA) have been applied
for optimization and analysis of process parameters. The
optimization using Taguchi method has revealed a
unique and powerful optimization tool that differs from
traditional applications.6 For instance, Kivak et al.7
studied the optimization of drilling parameters based on
the Taguchi method to minimize the surface roughness
(Ra) and thrust force (Ff). Their study showed that the
Materiali in tehnologije / Materials and technology 49 (2015) 1, 139–147 139
UDK 621.7.01:621.937:519.61/.64 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(1)139(2015)
cutting tool was the most significant parameter for the
Ra. Moreover, the results of verification test demon-
strated the Taguchi method for drilling operations was
successful to obtain the better surface quality of the
machined parts. Aslan et al.8 evaluated the Ra and
cutting-tool wear during the machining of AISI 4140 (63
HRC) steel with an experiment according to the
Taguchi’s L27 orthogonal array. From the ANOVA table,
it was found out that tool wear is affected by cutting
speed with 30 %. They suggested a 250 m/min cutting
speed, a 0.25 mm depth of cut and a 0.05 mm/r feed rate
to minimize the Ra value. Gunay and Yucel9 investigated
the Ra during the machining of high-alloy white cast iron
with an experiment according to the Taguchi L18
orthogonal array. According to ANOVA data, they
explained that the most important parameter was feed
rate for Ni-Hard with 62 HRC although the cutting speed
was the most important parameter for Ni - Hard with 50
HRC.
Neseli et al.10 studied the influence of the tool geome-
try on the surface finish when turning the AISI 1040
steel with an Al2O3/TiC tool using the response-surface
methodology (RSM). Their results indicated that the
tool-nose radius was the dominant factor for the surface
roughness with a 51.45 % contribution to the total vari-
ability of the model.
Asilturk and Akkus11 applied the Taguchi method to
minimize the surface roughness, which is Ra and Rz, in
turning of hardened AISI 4140 steel (51 HRC) using
coated carbide tools. In addition, their study explored the
effects of the cutting speed, the feed rate and the depth of
cut on the responses. From the ANOVA analysis, it was
determined that the feed rate was the most significant
parameter on results. Further, it was seen that the opti-
mum machining parameters for Ra and Rz were different.
Kacal and Gulesin12 optimized the machining parameters
for the finish turning of austempered cast iron
(GJS-400-15). Their experimental investigation was con-
ducted based on Taguchi’s L18 orthogonal array. Stati-
stical analysis, which was ANOVA, demonstrated that
feed rate is the most important factor with 69.5 % for
surface roughness (Ra). In addition, they identified the
best machining parameters for Ra to be: an austempering
temperature of 290 °C, a ceramic tool, a cutting speed of
800 m/min and a feed rate of 0.05 mm/r.
Sarikaya and Gullu13 studied the Taguchi design and
a response-surface-methodology-based analysis of the
machining parameters for CNC turning under MQL. It
was found that the most effective parameter for the sur-
face roughness is the feed rate. In addition, the cooling
conditions also significantly affect the surface roughness.
Kadirvel and Hariharan14 investigated the optimization of
the die-sinking micro-edm process for multiple perfor-
mance characteristics using the Taguchi-based grey rela-
tional analysis. Their study indicated that, on the basis of
a confirmation test, the improvement in the performance
characteristics was found to be as follows: MRR 3.86 %,
TWR 4.20 % and SR 3.51 %. Kivak15 investigated the
Taguchi-method-based optimization of drilling para-
meters when drilling the AISI 316 steel with a PVD
monolayer- and multilayer-coated HSS drills. The
analysis results revealed that the feed rate was the
dominant factor affecting the surface roughness and the
cutting speed was the dominant factor affecting the flank
wear. Koklu16 investigated the influence of the process
parameters and the mechanical properties of aluminum
alloys on the burr height and the surface roughness in dry
drilling using the Taguchi method. The analysis of
variance and Taguchi techniques were applied in order to
determine the effects of the drilling parameters.
The literature survey demonstrates that traditional
experimental design procedures are too complicated and
not easy to use. A large number of experimental tasks
have to be performed when the number of process para-
meters increases. The Taguchi method uses a special
design of orthogonal arrays to study the entire parameter
space with only a small number of experiments for solv-
ing this problem.11
The purpose of this study is to obtain the optimum
milling parameters (the cutting speed, the feed rate, the
number of cutting inserts) for the minimum surface
roughness and the maximum tool life during the milling
of AISI D3 steel. The Taguchi parameter-design app-
roach was employed to achieve these goals. Moreover, a
statistical analysis (ANOVA) was carried out to see
which machining parameter is statistically significant.
Finally, confirmation tests were conducted using the
optimum cutting conditions determined with the Taguchi
optimization method.
2 EXPERIMENTAL PROCEDURE
2.1 Milling experiments
In this study, AISI D3 cold-work tool steel was used
as the workpiece material. This steel type is used in
many manufacturing industries such as the ones manu-
facturing cold-extrusion and drilling moulds, mould
plates, powder-metallurgy kits, ceramic shaping moulds
and cold punches. The dimensions of the workpiece were
100 mm × 48 mm × 300 mm. The chemical composition
of the work material is as follows: 1.938 % C; 0.37 % Si;
0.22 % Mn; 10.66 % Cr; 0.22 % Ni; 0.135 % V; 0.062 %
W; 0.07 % Cu; 85.82 % Fe. The milling tests were per-
formed using a Johnford VMC 550 model, a three-axis
CNC vertical machine centre, equipped with the maxi-
mum spindle speed of 8000 r/min and a 10 kW drive
motor. The values of process parameters were selected
from the manufacturer’s handbook recommended and the
preliminary experiments for the tested material. Process
parameters and their levels are shown in Table 1. In all
the experiments, the depth of cut was determined as 1
mm. The experiments were carried out under dry cutting
conditions. To protect these experiment conditions for
each test, a new cutting insert was used for each experi-
M. SARIKAYA et al.: OPTIMIZATION OF THE PROCESS PARAMETERS FOR SURFACE ROUGHNESS ...
140 Materiali in tehnologije / Materials and technology 49 (2015) 1, 139–147
ment. The flowchart for the optimization of the milling
parameters is shown in Figure 1. In the milling experi-
ments, the coated carbide inserts manufactured by Mitsu-
bishi Carbide were used. Based on the ISO description,
they create a SEMT 13T3AGSN tool geometry and a
JM-chip-breaker form. In the experiments, the
ASX445-080A06R model of the tool holder was used.
The geometry of the cutting tool is shown in Figure 2.
The geometric features of the tools are listed Table 2.
2.2 Measurements
One of the most important quality indicators of the
machined parts is surface roughness or surface quality.
According to the standard, the average surface roughness
is defined as Ra. In present work, Ra was determined by
using a Mahr Perthometer M1 with a cut-off length of
0.8 mm and a sampling length of 5 mm based on ISO
4287 standard. It was considered by measuring the mean
of the three roughness results performed from different
locations on the workpiece.
Generally, the cutting tool-wear occurs in combina-
tion with the predominant wear mode, dependent upon
the cutting parameters, cooling/lubrication conditions,
workpiece material, cutting tool material and the tool
insert geometry. The forms of the cutting-tool wear,
M. SARIKAYA et al.: OPTIMIZATION OF THE PROCESS PARAMETERS FOR SURFACE ROUGHNESS ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 139–147 141
Figure 3: Conventional features of tool-wear measurements17
Slika 3: Obi~ajne meritve obrabe orodja17
Figure 2: Cutting tool employed in the experiments
Slika 2: Rezalno orodje, uporabljeno pri preizkusih
Table 1: Chosen factors and their levels
Tabela 1: Izbrani faktorji in njihovi nivoji
Symbol Cutting parameter
Levels for surface roughness Levels for tool life
1 2 3 1 2 3
A Cutting speed, V/(m/min) 80 120 180 416 500 600
B Feed rate, f/(mm/r) 0.08 0.12 0.18 0.08 0.1 0.125
C Number of cutting inserts, z 1 3 6 1 3 6
Table 2: Standard tool holder and cutting insert
Tabela 2: Standardno dr`alo orodja in rezalnega vlo`ka
Standard tool holder Standard cutting insert
ASX445-080A06R SEMT13T3AGSN-JM
D/mm d/mm L/mm Z aa/mm D1/mm) S1/mm) F1/mm) Re/mm
80 27 50 6 6 13.4 3.97 1.9 1.5
Figure 1: Flowchart for optimization of milling parameters
Slika 1: Diagram poteka optimizacije parametrov rezkanja
often expressed as the principal types of the tool wear
are nose, flank, notch and crater wear types, and Figure
3 shows how these wear features are usually measured.17
In the present work, tool deformations were measured
and investigated using a Mitutoyo light microscope with
a 0.001 mm sensitivity and a capability of magnifying 5
to 10 fold. The experiments revealed that the notch wear
was observed during the machining at high cutting
speeds. The determination of the machining time was
based on ISO 2688-1. When determining the tool life,
VBmax was taken as 1 mm. The flank wear on the cutting
tools was periodically measured and recorded to deter-
mine the tool life.
2.3 Control factors and orthogonal array
The cutting speed V (m/min), feed rate f (mm/r) and
number of cutting inserts (z) were selected as the control
factors for the surface-roughness and tool-life values,
and their levels were determined as shown in Table 1.
The orthogonal array (OA) enables an effective tool to
conduct the test with the small number of studies in the
Taguchi experimental method. The total degree-of-free-
dom (DFT) is the basis of orthogonal array for experi-
mental design.18 In present work, because there are three
control factors and three levels, the DFT is given as
twenty-six. In order to make the performance com-
parisons between control factors and its different
combinations, an OA having at least nine or twenty-
seven test trials (DFT + 1) shoud be selected. Therefore,
the standard L27 (33) OA is selected for the study and the
surface-roughness and tool-life values are measured via
the experimental design for each combination of the
control factors. Determination of the quality characteri-
stics of the measured control factors is provided by the
S/N ratios.
3 RESULTS AND DISCUSSIONS
3.1 Analysis of the signal-to-noise (S/N) ratio
With the Taguchi method we used the signal-to-noise
(S/N) ratio as the quality characteristic of choice. In
milling operations, a lower surface roughness and larger
tool life are indications of a better performance. There-
fore, in order to obtain the optimum machining perfor-
mance, the smaller-the-better (Equation 1) and larger-
the-better (Equation 2) ratios were selected for the
minimum surface roughness and maximum tool life,
respectively:
M. SARIKAYA et al.: OPTIMIZATION OF THE PROCESS PARAMETERS FOR SURFACE ROUGHNESS ...
142 Materiali in tehnologije / Materials and technology 49 (2015) 1, 139–147
Table 3: Experimental results, means and corresponding S/N ratios
Tabela 3: Eksperimentalni rezultati, povpre~ja in ustrezna razmerja S/N
Exp. No.
Cutting-parameter level Measured
surface
roughness,
Ra/ìm
Calculated
(S/N)/dB
Measured tool
life, T/s
Calculated,
(S/N)/dB
A
Cutting speed,
V/(m/min)
B
Feed rate,
f/(mm/r)
C
Number of
inserts, z
1 1 1 1 0.160 15.9176 435 52.7698
2 1 1 2 0.582 4.7015 426 52.5882
3 1 1 3 1.079 –0.6604 403 52.1061
4 1 2 1 0.383 8.3360 401 52.0629
5 1 2 2 0.784 2.1137 365 51.2459
6 1 2 3 1.194 –1.5401 338 50.5783
7 1 3 1 0.586 4.6420 369 51.3405
8 1 3 2 1.109 –0.8986 354 50.9801
9 1 3 3 1.400 –2.9226 328 50.3175
10 2 1 1 0.206 13.7227 266 48.4976
11 2 1 2 0.723 2.8172 236 47.4582
12 2 1 3 1.084 –0.7006 233 47.3471
13 2 2 1 0.456 6.8207 231 47.2722
14 2 2 2 0.869 1.2196 222 46.9271
15 2 2 3 1.261 –2.0143 193 45.7111
16 2 3 1 0.724 2.8052 216 46.6891
17 2 3 2 1.160 –1.2892 210 46.4444
18 2 3 3 1.409 –2.9782 167 44.4543
19 3 1 1 0.303 10.3711 118 41.4376
20 3 1 2 0.708 2.9993 110 40.8279
21 3 1 3 1.104 –0.8594 100 40.0000
22 3 2 1 0.552 5.1612 102 40.1720
23 3 2 2 0.945 0.4914 97 39.7354
24 3 2 3 1.290 –2.2118 86 38.6900
25 3 3 1 0.761 2.3723 97 39.7354
26 3 3 2 1.211 –1.6629 91 39.1808
27 3 3 3 1.449 –3.2214 82 38.2763
Smaller–the–better:
S
N n
y i
i
n
= −
⎛
⎝
⎜ ⎞
⎠
⎟
=
∑10 110 2
1
log (1)
Larger–the–better:
S
N n y ii
n
= −
⎛
⎝
⎜
⎞
⎠
⎟
=
∑10 1 110 2
1
log (2)
Here, yi is the ith measure of the experimental results
in a run/row and n gives the number of measurements in
each trial/row test.
The S/N ratios for the surface roughness and tool life
were calculated using Equations 1 and 2, as shown in
Table 3. The milling parameters were divided by consi-
dering different levels and possible effects, according to
the selected orthogonal array. According to the experi-
mental results, the mean of the surface roughness was
0.8701 μm and its mean S/N ratio was 2.353 dB. The
mean value of tool life and its mean S/N ratio were also
calculated as 232.4 s, and 46.031 dB, respectively.
Further, the effects of input parameters on the responses
can be analyzed with help of S/N ratios. These effects are
defined and evaluated according to the total mean values
of the experimental-trial results or S/N ratios. The mini-
mum surface-roughness and the maximum tool-life
values can be calculated from the total mean values of
the experimental-trial results for the surface roughness
and tool life. An important requirement when calculating
the optimum points is to identify the optimum levels of
machining parameters. They were defined by assessing
different levels of the input parameters, based on the
results from combinations produced by the OA. The
levels of the control factors were determined for both the
surface roughness and tool life, presented in Table 4, and
S/N graphics of these levels were used for the evaluation
(Figures 4 and 5). As shown in Figure 4, the surface
roughness is found to be minimal at low cutting speed
and feed rate and the minimum number of cutting
inserts. The surface roughness increases with a rise in the
cutting speed for a given value of the feed rate. Since the
temperature increases between the tool material and the
workpiece material through the friction under higher
cutting speed, the cut chips by the cutting tool stick over
the cutting insert with help of the high temperature.19
Therefore, the surface roughness increases during the
experiments at higher cutting speeds. The increased feed
rate and number of cutting inserts lead to vibration,
generating more heat and, thereby, a higher surface
roughness occurs.20 The tool life based on the wear value
of VBmax = 1 mm for all the experiments is drawn.
According to Figure 5, the longest tool life was obtained
at the lowest values of the cutting speed, feed rate and
number of cutters. When Figure 5 is evaluated, it is
possible to observe that an increased number of cutting
inserts leads to a decreased tool life. This can be
explained with the cutting temperature and the vibrations
generated at the tool/chip contact when the number of
cutting inserts is increased. It is known that a high cutt-
ing temperature occurs in the primary deformation zone
M. SARIKAYA et al.: OPTIMIZATION OF THE PROCESS PARAMETERS FOR SURFACE ROUGHNESS ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 139–147 143
Table 4: Response table for S/N ratios (dB) and means
Tabela 4: Tabela odzivov (dB) za razli~na razmerja S/N in za razli~na sredstva
Control
factors
Surface roughness (Ra) Tool life (T)
Level 1 Level 2 Level 3 Delta Level 1 Level 2 Level 3 Delta
S/N ratios
A 3.30 2.26 1.49 1.80 51.55 46.76 39.78 11.77
B 5.37 2.04 –0.35 5.72 47.00 45.82 45.27 1.73
C 7.79 1.16 –1.90 9.69 46.66 46.15 45.28 1.39
Means
A 0.81 0.88 0.92 0.12 379.89 219.33 98.11 281.78
B 0.66 0.86 1.09 0.43 258.56 226.11 212.67 45.89
C 0.46 0.90 1.25 0.79 248.33 234.56 214.44 33.89
Figure 4: Main effects plot of the factors and S/N graph for the sur-
face roughness
Slika 4: Prikaz glavnih faktorjev in odvisnost S/N od hrapavosti
povr{ine
with the increasing cutting speed. Thus, the wear mecha-
nisms are accelerated and so the tool life is decreased.
The S/N ratios of the surface-roughness data obtained
from the experimental work, later used to determine the
optimum level of each variable, were calculated in Table
3. Figure 4 illustrates the graphs of the S/N ratios that
were calculated for the surface roughness. As mentioned
above, the maximum value of the S/N ratios gives the
optimum cutting conditions. Thus, the optimum combi-
nation for the Ra was determined as A1B1C1 (A1 = 80
m/min, B1 = 0.08 mm/r, C1 = 1 cutting insert). The S/N
ratios for the tool-life data obtained from the experi-
mental results were calculated in Table 3. Figure 5
illustrates the graphs of the S/N ratios that were
calculated for the tool life. The optimum combination for
the tool life was determined as A1B1C1 (A1 = 416 m/min,
B1 = 0.08 mm/r, C1 = 1 cutting insert).
3.2 Analysis of variance (ANOVA)
The Taguchi method was used for determining the
optimum cutting conditions according to the S/N ratio,
while the control-factor correlation with the experimen-
tal results was determined with the help of an analysis of
variance (ANOVA). The analysis of variance (ANOVA)
was employed through the Minitab 16.0 Program. Table
5 shows the results of ANOVA for the surface roughness
and tool life. In addition to the degree of freedom, the
mean of squares (MS), the sum of squares (SS), the
F-ratio, P-values and the contribution (PCR) associated
with each factor were presented. This analysis was per-
formed for a confidence level of 95 %. The importance
of the input parameters in ANOVA analysis was identi-
fied by comparing the F-values of each input parameters.
The F-value determined in the ANOVA table was
compared with the value according to standard F-tables
for a given statistical level of importance.21 According to
the ANOVA table, the P-value is effective for all three
levels at the reliability level of 95 %, because the results
for the surface roughness and tool life are lower than
0.05. As the results of the evaluation of the surface
roughness, the percentage contributions of input para-
meters for A, B and C were determined as: (1.63, 21.96
and 75.27) %, respectively, and the error was 1.14 %
(Table 5). Thus, it was found that the number of cutting
inserts and the feed rate vary significantly more than the
cutting speed regarding the surface roughness in milling
the AISI D3 steel. The ANOVA table indicates that the
variable most significantly affecting the surface-rough-
ness value is the number of cutting inserts with 75.27 %
of PCR. This result clearly shows the effects of the
M. SARIKAYA et al.: OPTIMIZATION OF THE PROCESS PARAMETERS FOR SURFACE ROUGHNESS ...
144 Materiali in tehnologije / Materials and technology 49 (2015) 1, 139–147
Table 5: Results of ANOVA for responses
Tabela 5: Rezultati ANOVA-odzivov
Variation of
source
Degree of
freedom (DF)
Sum of squares
(SS)
Mean of squares
(MS) F-ratio P-value
Contribution
(%)
Surface roughness (Ra)
A 2 0.06141 0.03071 14.26 0.000 1.63
B 2 0.82931 0.41466 192.53 0.000 21.96
C 2 2.84270 1.42135 659.97 0.000 75.27
Error (e) 20 0.04307 0.00215 1.14
Total 26 3.77650 100
Tool life (T)
A 2 359615 179807 975.62 0.000 95.00
B 2 10018 5009 27.18 0.000 2.65
C 2 5228 2614 14.18 0.000 1.38
Error (e) 20 3686 184 0.97
Total 26 378547 100
Figure 5: Main effects plot of the factors and S/N graph for the tool
life
Slika 5: Prikaz glavnih faktorjev in odvisnost S/N od zdr`ljivosti
orodja
number of cutting inserts on the surface roughness in
milling the AISI D3 steel whose vibration generated a
higher frequency between the tool and the workpiece
with the increasing number of cutting inserts. The other
variable that has an effect on the Ra is the feed rate with
21.96 % of PCR. It is known that an increasing feed rate
increases the chip volume removed per unit time.7
Accordingly, Table 5 shows that the percentage
contributions of factors A, B, and C for the tool life are
(95.00, 2.65 and 1.38) %, respectively, and the error is
0.97 %. Thus, Table 5 indicates that the most effective
variable for the tool life is the cutting speed (95.00 %).
The feed rate and number of cutting inserts do not affect
the tool life. In milling operations, the cutting speed is
the most effective cutting parameter, reducing the tool
life, because an increasing cutting speed increases the
cutting temperature in the primary deformation zone. As
a result of this situation, the wear mechanisms are
accelerated.
3.3 Verification experiments
In the last step of the Taguchi approach, the optimi-
zation was confirmed via verification tests then the deter-
mination of the parameter levels giving the optimum
results. The verification-test results were obtained for the
optimum parameter levels (A1B1C1) of the surface
roughness. Later, the calculation of the predicted mini-
mum surface roughness Rap from Equation 3, taking into
consideration individual effects of factors A, B, C and
their levels, is as follows:
R T A T B T C Tap = + − + − + −Re Re Re Re( ) ( ) ( )1 1 1 (3)
where A1 is the mean (0.81 μm) of the experimental
trials at the first level of factor A. B1 is the mean (0.66
μm) of the experimental trials at the first level of factor
B. C1 is the mean (0.46 μm) of the experimental trials at
the first level of factor C. With Equation 3, the mini-
mum surface roughness was calculated as 0.147 μm. As
determined in Figure 5, A, B, C and their levels were
used for the calculation of the predicted optimum tool
life. The equation for the predicted optimum tool life is
as follows:
T T A T B T C Tp T T T T= + − + − + −( ) ( ) ( )1 1 1 (4)
where A, B and C are the means (379.89, 258.56 and
248.33) s of the experimental trials at the optimum
levels of these factors. With Equation 4, the maximum
tool life was calculated as 441.67 s. The confidence
interval (CI) was conducted to confirm the output para-
meters of the verification test. The CI for the estimated
optimum results was calculated using the following
equation:22
CI F V
n r
= ⋅ ⋅ +
⎛
⎝
⎜
⎞
⎠
⎟
 ; ;1
1 1
Ve ep
eff
(5)
where, with respect to F
;1;Ve, F is the ratio of signifi-
cant level 
, 
 is the significance level, 1 – 
 is the
confidence level and Ve is the degree of freedom of the
pooled error variance. Vep is the pooled error variance, r
is the number of repeated trials (r%), N is the total num-
ber of experimental trials, neff is the number of effective
measured results defined as22:
n
N
beff
=
+1
(6)
where b is total degress of freedom associated with
items used in estimate.
In the present investigation, three verification tests (r
= 3) were made to assess the performance of the experi-
mental study used for the surface roughness at optimum
parameters (A1B1C1). The value of F
;1;Ve = 4.30 which
has a 95 % confidence level for the surface roughness,
was found with respect to the values of 
 = 0.05 and Ve =
20, based on the look-up table. According to Equations 5
and 6, the confidence interval (CI) is calculated as 0.07.
The result value of the confirmation test performed for
the surface roughness is expected to be in the confidence
interval of (0.147 ± 0.07) or (0.077 – 0.217) with a 95 %
confidence level.
In the present work, the values of the surface rough-
ness from the three confirmation tests performed with
regard to the optimum levels (A1B1C1) were measured as
(0.159, 0.162 and 0.168) μm. As shown in Table 6, the
mean of the measurements was 0.163 μm. The mean
result of the confirmation tests in the optimum condi-
tions is within the confidence interval (0.077 < 0.163 <
0.217). As the mean result falls within this limit, the
experiment is considered satisfactory. The optimization
of process parameters was achieved using the Taguchi
method for the surface roughness at the confidence level
of 95 %.
Three confirmation experiments were carried out
under the optimum conditions (A1B1C1) for the tool life.
The value of F
;1;Ve = 4.30 which has a 95 % confidence
level for the tool life, was found with respect to the
values of 
 = 0.05, and Ve = 20, based on the look-up
table. According to Equations 5 and 6, the confidence
interval (CI) is calculated as 21.4 s. The result value of
the confirmation test performed for the tool life is
expected to be in the confidence interval of (441.7 ±
21.4) or (420.3 – 463.1) with a 95 % confidence level. In
this study, the values of the tool life from the three con-
firmation tests performed with regard to the optimum
levels (A1B1C1) were measured as (440.2, 436.4 and
425.7) s. As shown in Table 7, the mean of the measure-
ments was 434.1 s. The mean result of the confirmation
tests in the optimum condition is within the confidence
interval (420.3 < 434.1 < 463.1). As the mean result falls
within this limit, the experiment is considered satisfac-
tory. The optimization of the process parameters was
achieved using the Taguchi method for the tool life at the
confidence level of 95 %. According to the optimum test
and the predicted combination, the comparisons of the
surface roughness and tool life, and the combinations
M. SARIKAYA et al.: OPTIMIZATION OF THE PROCESS PARAMETERS FOR SURFACE ROUGHNESS ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 139–147 145
(A2B1C2) selected in the twenty-seven initial trials are
given in Table 6. According to the comparison table, the
surface roughness and tool life are reduced from 0.723
μm to 0.163 μm, and from 236 s to 434.1 s, respectively.
The improved efficiency due to the optimum combina-
tion was increased for the surface roughness and tool life
by up to 77.45 % ((0.723 – 0.163)/0.723) and by up to
45.63 % ((434.1 – 236)/434.1), respectively. The perfor-
mance comparisons between the initial parameters and
the optimum conditions are given in Table 6. In addition
to, the quality losses are listed in Table 7.
The quality losses between the initial and optimum
combinations for both the surface roughness and tool life
are calculated as follows18:
L y
L y
n
opt
int
( )
( )
≈ ⎛⎝
⎜ ⎞
⎠
⎟1
2
3
Δ
(7)
where Lopt(y) and Lint(y) are the optimum and initial
combinations, respectively. n is the difference between
the S/N ratios for the optimal and initial combinations.
The differences between S/N ratios that can be used to
evaluate the quality loss of the optimum combination
for the surface roughness and tool life, were found to be
13.08 (n = 13.08 (= 15.9 – 2.82)) and 5.29 (n = 5.29
(52.74 – 47.45)), respectively. The quality loss was cal-
culated as 0.049 using Equation 7 for the surface rough-
ness. Thus, the quality loss of the surface roughness for
the optimum combination is only 4.9 % of the initial
combination. Therefore, the quality losses of the surface
roughness were reduced to 95.1 % through the Taguchi
application. The quality loss for the tool life was cal-
culated as 0.295, using Equation 7. The quality loss of
the tool life for the optimum combination is only 29.5 %
of the initial combination. Consequently, the quality
losses of the tool life were reduced to 70.5 % through
the Taguchi method.
4 CONCLUSIONS
This study focuses on the Taguchi method used for
investigating the influence of the cutting parameters on
the surface roughness and tool life when face milling the
AISI D3 steel. In the milling experiments, different
levels of the cutting speed, the feed rate and the number
of cutting inserts as the machining parameters are used
under dry cooling conditions. The experimental results
were evaluated using the analysis of the signal-to-noise
ratio, the main effect graphs of means and ANOVA. The
optimum operating parameters are determined using the
Taguchi method. The results can be drawn as follows:
The optimum levels of the control factors providing a
better surface roughness and tool life were: A1 (the cutt-
ing speed, 80 m/min), B1 (the feed rate, 0.08 mm/r), and
C1 (the number of cutting inserts, 1 insert), and A1 (the
cutting speed, 416 m/min), B1 (the feed rate, 0.08 mm/r),
C1 (the number of cutting inserts, 1 insert), respectively.
The effects of the process parameters on the surface
roughness and tool life were detected by ANOVA anal-
ysis. It was revealed that the number of cutting inserts
the most important parameter influencing the surface
roughness with 75.27 %. Further, cutting speed was the
most important parameter influencing the tool life with
95 %.
Through the confirmation experiment, the surface
roughness was obtained, with one of the initial combina-
tions, as 0.723 μm, and the surface roughness was
improved to 77.45 %. The tool life was obtained, with
one of the initial combinations, as 236 s, and it was im-
proved to 45.63 %. The quality losses for the surface
roughness and the tool life determined from optimum
points were calculated as 4.9 % and 29.5 %, respectively.
The confirmation-experiment results indicated that
the observed values were within the calculated confi-
dence interval (CI) for the confidence level of 95 %.
M. SARIKAYA et al.: OPTIMIZATION OF THE PROCESS PARAMETERS FOR SURFACE ROUGHNESS ...
146 Materiali in tehnologije / Materials and technology 49 (2015) 1, 139–147
Table 6: Comparisons of surface roughness and tool life
Tabela 6: Primerjave hrapavosti povr{ine in zdr`ljivosti orodja
Comparison
Surface roughness Tool life
Level Ra/ìm (S/N)/dB Level T/s (S/N)/dB
Initial combination A2B1C2 0.723 2.82 A2B1C2 236 47.45
Optimum combination
(Experiment) A1B1C1 0.163 15.9 A1B1C1 434.1 52.74
Optimum combination
(Prediction) A1B1C1 0.147 16.63 A1B1C1 441.67 52.92
Table 7: Comparisons of experimental trials
Tabela 7: Primerjave eksperimentalnih poskusov
Comparison
Surface roughness Tool life
Level Ra/ìm Quality loss Level T/s Quality loss
Initial combination A2B1C2 0.723 – A2B1C2 236 –
Optimum combination
(Prediction) A1B1C1 0.147±0.07 – A1B1C1 434.1±21.4
Optimum combination
(Confirmation) A1B1C1 0.163 4.9 % A1B1C1 441.67 29.5 %
The present work demonstrates to industrial users
how to apply the Taguchi parameter design for opti-
mizing the machining performance with a minimum cost
and time in milling the AISI D3 steel. Moreover, this
study considered other factors (the cutting speed, the
feed rate, the number of cutting inserts) to find how the
control factors affect the surface roughness and tool life.
Future works may be extend to analyzing the effects of
some additional variables such as different tool holders,
tool materials, workpiece materials and cooling condi-
tions (wet, HPC, MQL, cryogenic, etc.).
5 REFERENCES
1 T. S. Lee, Y. J. Lin, A 3D predictive cutting force model for end
milling of parts having sculptured surfaces, Int. J. Adv. Manuf. Tech.,
16 (2000), 773–783
2 H. Dilipak, A. Guldas, A. Gezgin, An investigation of the number of
ýnserts effect on the machining time and metal removal rate during
the milling of AISI D3 steel at high cutting speeds, Journal of Me-
chanical Engineering, 55, (2009) 7–8, 438–443
3 E. Kabakli, M. Bayramoðlu, N. Geren, Evaluation of the surface
roughness and geometric accuracies in a drilling process using the
Taguchi analysis, Mater. Tehnol., 48 (2014) 1, 91–98
4 H. Durmuº, Optimization of multi-process parameters according to
the surface quality criteria in the end milling of the AA6013 alumi-
num alloy, Mater. Tehnol., 46 (2012) 4, 383–388
5 S. Tamas, Cutting force modeling using artificial neural networks,
Journal of Materials Processing Technology, 92 (1999), 344–349
6 M. S. Phadke, Quality engineering using robust design, Prentice
Hall, Englewood Cliffs, New Jersey 1989
7 T. Kivak, G. Samtas, A. Cicek, Taguchi method based optimisation
of drilling parameters in drilling of AISI 316 steel with PVD
monolayer and multilayer coated HSS drills, Measurement, 45
(2012), 1547–1557
8 E. Aslan, N. Camuscu, B. Birgoren, Design optimization of cutting
parameters when turning hardened AISI 4140 steel (63 HRC) with
Al2O3 + TiCN mixed ceramic tool, Mater. Des., 28 (2007),
1618–1622
9 M. Gunay, E. Yucel, Application of Taguchi method for determining
optimum surface roughness in turning of high-alloy white cast iron,
Measurement, 46 (2013), 913–919
10 S. Neseli, S. Yaldýz, E. Turkes, Optimization of tool geometry para-
meters for turning operations based on the response surface metho-
dology, Measurement, 44 (2011) 3, 580–587
11 I. Asilturk, H. Akkus, Determining the effect of cutting parameters
on surface roughness in hard turning using the Taguchi method,
Measurement, 44 (2011), 1697–1704
12 A. Kacal, M. Gulesin, Determination of optimal cutting conditions in
finish turning of austempered ductile iron using Taguchi design
method, J. Sci. Ind. Res. India, 70 (2011), 278–283
13 M. Sarikaya, A. Gullu, Taguchi design and response surface metho-
dology based analysis of machining parameters in CNC turning
under MQL, Journal of Cleaner Production, 65 (2014), 604–616
14 A. Kadirvel, P. Hariharan, Optimization of the die-sinking micro-
edm process for multiple performance characteristics using the
Taguchi-based grey relational analysis, Mater. Tehnol., 48 (2014) 1,
27–32
15 T. Kivak, Optimization of surface roughness and flank wear using the
Taguchi method in milling of Hadfield steel with PVD and CVD
coated inserts, Measurement, 50 (2014), 19–28
16 U. Koklu, Influence of the process parameters and the mechanical
properties of aluminum alloys on the burr height and the surface
roughness in dry drilling, Mater. Tehnol., 46 (2012) 2, 103–108
17 D. E. Dimla Sr., P. M. Lister, On-line metal cutting tool condition
monitoring I: force and vibration analyses, International Journal of
Machine Tools & Manufacture, 40 (2000), 739–768
18 Y. T. Liu, W. C. Chang, Y. A. Yamagata, A Study on optimal com-
pensation cutting for an aspheric surface using the Taguchi method,
CIRP Journal of Manufacturing Science and Technology, 3 (2010),
40–48
19 V. Savas, C. Ozay, Analysis of the surface roughness of tangential
turn-milling for machining with end milling cutter, Journal of Mate-
rials Processing Technology, 186 (2007), 279–283
20 I. Korkut, M. Kasap, I. Ciftci, U. Seker, Determination of optimum
cutting parameters during machining of AISI 304 austenitic stainless
steel, Materials and Design, 25 (2004), 303–305
21 J. Antony, D. Preece, Understanding, Managing and Implementing
Quality, Frameworks Techniques and Cases, Routledge, London
2002
22 R. K. Roy, Design of Experiments Using the Taguchi Approach,
John Wiley & Sons, USA 2001
M. SARIKAYA et al.: OPTIMIZATION OF THE PROCESS PARAMETERS FOR SURFACE ROUGHNESS ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 139–147 147

N. REVATHI et al.: THIN TIN MONOSULFIDE FILMS DEPOSITED WITH THE HVE METHOD ...
THIN TIN MONOSULFIDE FILMS DEPOSITED WITH THE HVE
METHOD FOR PHOTOVOLTAIC APPLICATIONS
TANKA PLAST HVE KOSITROVEGA MONOSULFIDA ZA
UPORABO V FOTOVOLTAIKI
Naidu Revathi, Sergei Bereznev, Julia Lehner, Rainer Traksmaa, Maria Safonova,
Enn Mellikov, Olga Volobujeva
Tallinn University of Technology, Department of Materials Science, Ehitajate tee 5, 19086 Tallinn, Estonia
revathi.naidu@ttu.ee
Prejem rokopisa – received: 2013-10-01; sprejem za objavo – accepted for publication: 2014-02-17
In the present study, thin films of SnS photo-absorbers with two different thicknesses of 0.5 μm and 1 μm were deposited onto
molybdenum-covered soda-lime-glass substrates using high-vacuum evaporation (HVE). The changes in the structural,
phase-composition, morphological and photo-electrochemical properties depending on the film thickness were studied. The
films showed a polycrystalline orthorhombic SnS crystal structure with (040) as the preferred orientation. It was observed that
the films showed improved crystallinity with the increase in the film thickness. Raman analysis confirmed the presence of single
phase SnS without any other binary phases. Photoconductivity measurements revealed that the layers have the p-type
conductivity. The SnS layers grown at a thickness of 1 μm showed a high photosensitivity and could be considered as an
absorber layer for solar-cell applications.
Keywords: SnS, thin films, HVE, photo-electrochemical, solar cell
V tej {tudiji je bila z visokovakuumskim naparevanjem (HVE) nanesena tanka plast SnS fotoabsorberja z dvema debelinama 0,5
μm in 1 μm na z molibdenom prekrito podlago iz natrij-kalcijevega stekla. Preu~evane so bile spremembe v strukturi, sestavi
faz, morfolo{ke in fotoelektrokemijske lastnosti v odvisnosti od debeline plasti. Te plasti so pokazale polikristalno ortorombi~no
kristalno strukturo SnS z (040) kot prednostno orientacijo. Opa`eno je bilo, da se s pove~anjem debeline plasti izbolj{a njena
kristalini~nost. Ramanska analiza je potrdila prisotnost samo ene faze SnS brez drugih binarnih faz. Meritve fotoprevodnosti so
odkrile, da imajo plasti prevodnost p-tipa. Plasti SnS, ki so zrasle na debelino 1 μm, imajo veliko fotoob~utljivost in so lahko
uporabne pri son~nih celicah kot absorpcijska plast.
Klju~ne besede: SnS, tanke plasti, HVE, fotoelektrokemi~en, son~na celica
1 INTRODUCTION
Tin monosulfide (SnS) is a layered compound semi-
conductor that crystallizes in an orthorhombic structure
with the lattice parameters of a = 0.4329 nm, b = 1.1193
nm and c = 0.398 nm.1 It has a high absorption coeffi-
cient with a photon energy threshold of 1.3 eV and
exhibits the p-type conductivity.2 Besides, its constituent
elements Sn and S are earth-abundant and nontoxic. SnS
has been used in photocondutors,3 photovoltaic conver-
sion,1 holographic recording media,4 solar control,5
near-infrared detector,6 etc. Owing to these advantages
SnS has to be the promising material in the fabrication of
thin-film solar cells. So far, the best reported conversion
efficiency for SnS solar cells is about 4.4 %,7 which is a
very low value when compared to CIGS (about 21.7 %)8
and CZTS (12.6 %) solar cells.9 The reason for the low
efficiency of the SnS solar cells so far has not been
understood. In the present study, SnS thin films with two
different thicknesses of 0.5 μm and 1 μm were grown on
molybdenum (Mo) covered soda-lime-glass substrates
and the influence of the preparative parameters of depo-
sition on the microstructural parameters and photo-
current response of the films is reported.
2 EXPERIMENTAL WORK
Thin films of SnS were deposited onto Mo-covered
soda-lime-glass substrates with the high-vacuum evapo-
ration technique using BOC EDWARDS Auto 500 sys-
tems. The films with two different thicknesses of 0.5 μm
and 1 μm were deposited at a constant substrate tem-
perature (Ts) of 300 °C, with a deposition rate of
0.5 nm/s. The deposition was carried out in a vacuum
chamber at a pressure of around 1.33 × 10–6 mbar with a
deposition source-to-substrate distance of  25 cm. The
crystalline-phase identification and phase-composition
determination were done with a Bruker D5005 X-ray
diffractometer and a Horiba LabRam HR spectrometer,
respectively. The surface morphology of the as-deposited
layers was observed using high-resolution scanning
electron microscopy (HR-SEM Zeiss ULTRA 55) and an
atomic force microscope (AFM, Bruker Nanoscope V
controller with the application module MultiMode 8.10).
The type of conductivity of the prepared SnS films was
identified by performing the photo-electrochemical
measurements in a background electrolyte solution of
0.1 M H2SO4.
Materiali in tehnologije / Materials and technology 49 (2015) 1, 149–152 149
UDK 532.6:669.058.4:669.058.61 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(1)149(2015)
3 RESULTS AND DISCUSSION
Figure 1 shows X-ray diffraction (XRD) profiles of
the SnS thin films deposited onto Mo-coated glass
substrate at a substrate temperature of 300 °C and a
detailed XRD pattern of the (040) peak is shown in the
inset. XRD profiles of all the layers had diffraction peaks
at 2  (26.12°, 31.96° and 39.04°), corresponding to
the polycrystalline orthorhombic crystal structure of SnS
with (040) as the dominant peak and being in agreement
with the JCPDS data (card no. 39-0354). The strong-
characteristic (040) diffraction peak could be due to the
fact that most of the SnS crystallites grew preferentially
along the (010) crystal plane that has the lowest surface
energy,10 with its orientation parallel to the substrate. At
a thickness of 1 μm, the (040) plane of the as-deposited
SnS film showed an enhanced intensity. The average
crystallite size of the films was calculated using the full
width at half maximum (FWHM) of the (040) plane with
Scherrer’s relation.11 The FWHM of the (040) peak
decreased with the film thickness, which indicates that
the film crystallinity had improved. The evaluated
average crystallite size of the prepared SnS films was
found to be 18.9 nm and 58.9 nm for the thicknesses of
500 nm and 1 μm, respectively. The observed increase in
the crystallite size at the higher thickness may be due to
a decrease in the disorderness and microstrain in the
thicker films. Analogous results for nanocrystalline SnS
thin films prepared with electron-beam evaporation were
also reported by Shaaban et al.12 The microstrains of SnS
thin films were calculated from equation13 and the cal-
culated values are 0.00183 and 0.00061 for the thick-
nesses 500 nm and 1 μm, respectively.
Raman measurements were performed to get detail
information of the phase composition of the films. The
room-temperature Raman spectra of the SnS thin films
were measured in a wavenumber range of 50–550 cm–1
and are shown in Figure 2. The peaks at (96, 163, 192
N. REVATHI et al.: THIN TIN MONOSULFIDE FILMS DEPOSITED WITH THE HVE METHOD ...
150 Materiali in tehnologije / Materials and technology 49 (2015) 1, 149–152
Figure 2: Raman spectra of SnS films at the thicknesses of 500 nm
and 1 μm
Slika 2: Ramanska spektra plasti SnS pri debelinah 500 nm in 1 μm
Figure 1: XRD patterns of SnS thin films (the inset: detailed (040)
peak profiles)
Slika 1: XRD-posnetka tankih plasti SnS (vlo`ek: podrobnej{i profil
vrhov (040))
Figure 3: SEM pictures of SnS thin films
Slika 3: SEM-posnetka povr{ine tanke plasti SnS
and 220) cm–1 can be attributed to the orthorhombic SnS.
The Raman modes observed at (96, 192 and 220) cm–1
belong to the Ag mode whereas the band line at 163 cm–1
can be ascribed to the B2g mode.14 The spectra do not
indicate the presence of traces of SnS2 or Sn2S3 phases in
the films.
The HR-SEM images of the SnS thin films deposited
at the thicknesses of 0.5 μm and 1 μm are presented in
Figure 3. The surfaces of the films are composed of den-
sely packed flake-like particles with the average diame-
ters of about 179 nm and 429 nm for the thicknesses of
500 nm and 1 μm, respectively.
Figure 4 shows the AFM micrographs of the SnS
thin films that were scanned over an area of 1 μm ×
1 μm. It is clear from the images that the thinner films
(500 nm) have smoother surfaces compared to the
thicker films (1 μm). The root-mean-square roughness
(Rq) and the mean-value roughness (Ra) for the film with
the 500 nm thickness are 20.6 nm and 16.7 nm, whereas
for the thicker film (1 μm) the values of Rq and Ra are
about 25.8 nm and 20.3 nm, respectively. The average
grain size of the as-deposited SnS films increased from
24 nm to 60 nm with the increase in the film thickness.
Figure 5 shows the photocurrent response of the SnS
thin films, taken under illumination in the background
electrolyte. The photocurrent was found to be increased
towards the negative region of the applied potential, con-
firming that the films exhibit the p-type
electroconductivity. The results indicate that the best
photosensitivity was obtained with the SnS absorber
layer with a thickness of 1 μm.
The above studies revealed that the SnS thin films
deposited at a thickness of 1 μm showed a larger crystal-
lite size of 58.9 nm with a surface roughness of 25.8 nm
and a high photosensitivity. These results are promising
for the fabrication of complete solar-cell structures based
on a SnS photo-absorber.
4 CONCLUSIONS
Thin films of SnS were successfully deposited onto a
Mo-coated soda-lime-glass substrate using the HVE
technique with two different thicknesses of 500 nm and
1 μm. The XRD and Raman analysis showed that the
films have a single-phase SnS orthorhombic crystal
structure without any other binary traces of SnS2, Sn2S3,
etc. The layers deposited with a thickness of 1 μm com-
pared with the layers with a thickness of 0.5 μm show an
improved crystallinity, a uniform and densely packed
surface morphology and a high surface roughness. They
exhibit the p-type conductivity and high photosensitivity.
Therefore, the prepared SnS photo-absorber layers could
be applied in complete solar-cell structures.
Acknowledgements
Estonian Centre of Excellence in Research Project
TK117T "High-technology Materials for Sustainable
Development", Estonian Energy Technology Program
(project AR 10128), Estonian Ministry of Higher
Education and Science (targeted project T099) and
Estonian Science Foundation (MJD213, G8147) are
acknowledged for the financing of the research.
N. REVATHI et al.: THIN TIN MONOSULFIDE FILMS DEPOSITED WITH THE HVE METHOD ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 149–152 151
Figure 5: PEC measurements of a SnS film deposited with a thickness
of 1 μm
Slika 5: PEC-meritev nanesene plasti SnS z debelino 1 μm
Figure 4: AFM images of SnS films of: a) 0.5 μm and b) 1 μm
Slika 4: AFM-posnetka plasti SnS: a) 0,5 μm in b) 1 μm
5 REFERENCES
1 B. Gosh, M. Das, P. Banerjee, S. Das, Semicond. Sci. Technol., 24
(2009) 2, 025024
2 J. Vidal, S. Lany, M. d’Avezac, A. Zunger, A. Zakutayev, J. Francis,
J. Tate, Appl. Phys. Lett., 100 (2012) 3, 032104
3 G. H. Yuea, W. Wanga, L. S. Wanga, X. Wanga, P. X. Yanb, Y. Che-
na, D. L. Peng, J. Alloy Compd., 474 (2009) 1–2, 445–449
4 G. Valiukonis, D. A. Guseinova, G. Krivaite, A. Sileica, Phys. Stat.
Sol. B, 135 (1986), 299–307
5 A. Ortiz, J. C. Alonso, M. Garcia, J. Toriz, Semicond. Sci. Technol.,
11 (1996) 2, 243–247
6 N. Koteswara Reddy, K. T. Ramakrishna Reddy, G. Fisher, R. Best,
P. K. Dutta, J. Phys. D: Appl. Phys., 32 (1999) 9, 988–990
7 P. Sinsermsuksakul, L. Sun, S. W. Lee, H. H. Park, S. B. Kim, C.
Yang, R. G. Gordon, Adv. Energy Mater., 4 (2014) 15, 1400496
8 ZSW sets 21.7 % thin film efficiency record: http://www.pv-
magazine.com/news/details/beitrag/zsw-sets-217-thin-film-efficiency-
record_100016505/#axzz3OhYBRNbe (accessed 22 September
2014)
9 W. Wang, M. T. Winkler, O. Gunawan, T. Gokmen, T. K. Todorov, Y.
Zhu, D. B. Mitzi, Adv. Energy Mater., 4 (2014) 7, 1301465
10 W. G. Pu, Z. Z. Lin, Z. W. Ming, G. X. Hong, C. W. Qun, H. Tana-
mura, M. Yamaguchi, H. Noguchi, T. Nagatomo, O. Omoto, 24th
IEEE PV Specialists Conference, 1 (1994), 365–368
11 B. E. Warren, X-ray Diffraction, Reprint, Dover Publications, Inc.,
New York 1990, 253
12 E. R. Shaaban, M. S. Abd El-Sadek, M. El. Hagary, I. S. Yahia, Phys.
Scr., 86 (2012) 1, 015702
13 K. Santhosh Kumar, C. Manoharan, L. Amalraj, S. Dhanapandian, G.
Kiruthigaa, K. Vijayakumar, Cryst. Res. Technol., 47 (2012) 7,
771–779
14 P. M. Nikolic, Lj. Miljkovic, P. Mihajlovic, B. Lavrencic, J. Phys. C:
Solid State Phys., 10 (1977) 11, L289–L292
N. REVATHI et al.: THIN TIN MONOSULFIDE FILMS DEPOSITED WITH THE HVE METHOD ...
152 Materiali in tehnologije / Materials and technology 49 (2015) 1, 149–152
A. DELI] et al.: MECHANICAL PROPERTIES OF THE AUSTENITIC STAINLESS STEEL X15CrNiSi20-12 ...
MECHANICAL PROPERTIES OF THE AUSTENITIC STAINLESS
STEEL X15CrNiSi20-12 AFTER RECYCLING
MEHANSKE LASTNOSTI AVSTENITNEGA NERJAVNEGA JEKLA
X15CrNiSi20-12 PO RECIKLIRANJU
Alen Deli}1, Omer Kablar2, Almir Zuli}1, D`emal Kova~evi}1, Nusmir Hod`i}1,
Enis Bar~i}1
1CIMOS TMD Automobilska industrija, d. o. o., Sarajevska 62, Grada~ac, Bosnia and Herzegovina
2ArcelorMittal Zenica, Steel Plant Department, Bulevar Kralja Tvrtka I, no. 17, Zenica, Bosnia and Herzegovina
alen.delic@cimos.eu
Prejem rokopisa – received: 2013-11-21; sprejem za objavo – accepted for publication: 2014-03-28
X15CrNiSi20-12 is a steel that has excellent resistance to corrosion and heat as well as good strength at room and elevated
temperatures. According to EN 10095, the forged steel is used for products resistant to hot gases and combustion products at
temperatures above 550 °C. As annealed, the steel is nonmagnetic, but becomes slightly magnetic after cold working. Typical
uses include furnace parts, heating elements, aircraft and jet-engine parts, heat exchangers, carburizing and annealing boxes,
sulfite liquor-handling equipment, kiln liners, boiler baffles, refinery and chemical processing equipment, and auto-exhaust
parts.
The influence of re-melted back (waste) materials on the quality of X15CrNiSi20-12 steel was determined for re-melted (MM)
and re-alloyed (AM) material with tensile strength tests using two standard test specimens at room temperature and 740 °C. The
tensile properties were used for the analysis and a comparison of the features in order to obtain suitable material for recycling.
Keywords: austenitic stainless steel, mechanical properties, re-alloying, recycling
Jeklo X15CrNiSi20-12 je toplotno obstojno, z odli~no odpornostjo proti koroziji in visoki temperaturi ter z dobro trdnostjo pri
sobni in pri povi{anih temperaturah. Skladno s standardom EN 10095 se jeklo uporablja za odkovke, ki so odporni proti vro~im
plinom in produktom zgorevanja pri temperaturah nad 550 °C. @arjeno jeklo je nemagnetno in pri hladni predelavi postane rahlo
magnetno. Zna~ilna uporaba je za dele pe~i, grelne elemente, letalstvo in dele reaktivnih motorjev, za toplotne izmenjevalce,
{katle za naoglji~enje in `arjenje, za opremo, ki pride v stik z raztopino sulfita, vodila v ogrevnih pe~eh, za {pirale kotlov, v
rafinerijah, v opremi kemijske procesne industrije in za dele avtomobilskih izpu{nih sistemov.
Vpliv pretaljevanja odpadkov materiala na kvaliteto jekla X15CrNiSi20-12 je bil dolo~en na pretaljenem (MM) in na dodatno
legiranem (AM) materialu z nateznim preizkusom materiala z navadnimi vzorci za natezne preizkuse pri sobni temperaturi in pri
740 °C. Natezne lastnosti so bile uporabljene za analizo in za dolo~itev, kak{en material je primeren za recikliranje.
Klju~ne besede: avstenitno nerjavno jeklo, mehanske lastnosti, ponovno legiranje, recikliranje
1 INTRODUCTION
X15CrNiSi20-12 steel is resistant to corrosion and is
generally considered to be a fire-proof alloy, with the
temperature for the destructive loss of weight being
about 1093 °C. The standard EN 10095 lists the require-
ments for the chemical composition of the austenitic
stainless steel X15CrNiSi20-12. Table 1 shows the
chemical composition of the raw material (RM), the
re-melted material without re-alloying (MM) and the
material with re-alloying (AM).1 The examinations were
performed with a SPECTRO LAB LAVM10 metal
analyzer.
1.1 Mechanical and working properties
The basic requirements for these steels involve their
applications at elevated temperatures and at higher loads.
The X15CrNiSi20-12 steel has a stable austenitic
microstructure and it can be used up to 1000 °C. The
occurrence of undesirable 	-ferrite is possible only in
exceptional cases, when the contents of the elements that
stabilize the ferrite are at the upper limit of the allowable
content, with a low content of austenite-stabilizing
elements that stabilize the austenite, especially in the
case of low contents of carbon and nitrogen2.
Materiali in tehnologije / Materials and technology 49 (2015) 1, 153–156 153
UDK 669.14.018.8:658.567.1:620.172 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(1)153(2015)
Table 1: Chemical composition of the X15CrNiSi20-12 steel according to EN 100951
Tabela 1: Kemijska sestava jekla X15CrNiSi20-12 skladno z EN 1000951
Material condition
Chemical composition (mass fractions, w/%)
C Mn Si Cr Ni N P S
EN 10095 <0.20 <2.00 1.5–2.5 19–21 11–13 <0.11 <0.04 <0.030
Raw – RM 0.149 1.89 1.53 19.33 11.39 0.061 0.028 0.002
Remelted – MM 0.130 1.59 1.34 19.30 11.44 0.065 0.029 0.013
Realloyed – AM 0.133 1.76 1.50 19.28 11.34 0.066 0.030 0.020
Austenitic steels show very good ductility and large
elongations at low temperatures. Furthermore, the tensile
strength is significantly increased at low temperatures.
These properties make austenitic steels suitable for use at
low temperatures. However, when increasing the tempe-
rature the strength properties are decreased. The influ-
ence of delta ferrite can also not be ignored, as it
increases the tensile strength and also the dispersion-
hardening mechanism that increases the yield strength3.
Table 2 lists the basic mechanical properties accord-
ing to the EN 10095 standard.
Table 2: Mechanical properties of the austenitic stainless steel
X15CrNiSi20-121
Tabela 2: Mehanske lastnosti avstenitnega nerjavnega jekla
X15CrNiSi20-121
Yield strength Rp0.2 /MPa > 230
Yield strength Rp1.0 /MPa > 270
Tensile strength Rm/MPa 500–750
Elongation A/%) > 28
Hardness, HB < 223
2 EXPERIMENTAL
2.1 Tensile properties at 20 °C and 740 °C
The tensile strength of the re-melted (MM) and
re-alloyed (AM) steel after annealing was determined at
room temperature and at 740 °C. The results showed that
the tensile properties of both melts are uniform and
within standard limits4. A metallographic analysis
showed an austenitic microstructure with carbide preci-
pitates at the grain boundaries. In the re-melted steel
(MM) the grains are coarser due to the longer holding
time at the quenching temperature5. The hardness testing
was carried at room temperature and the results are given
in Table 3.
Figure 1 provides a graphical representation of the
tensile strength at room temperature and at 740 °C and
the values prescribed in the standard EN 10095 for room
temperature. The tensile properties at 740 °C are not
prescribed by this standard.
The test results show that all the tensile strengths at
20 °C are within the limit values of the EN 10095 stan-
dard. The tensile strengths of the raw and re-melted steel
are similar and close to the maximum prescribed in the
standard6. The highest strength was obtained for the raw
steel (RM) 730 MPa, followed by the re-melted steel
(MM) 703 MPa, and the lowest value of 689 MPa for the
re-alloyed steel (AM) (Figure 2).
The tensile strength at the test temperature of 740 °C
shows a reverse trend, from 281 MPa for the raw steel,
followed by the re-melted steel without re-alloying at
297 MPa, and finally for the re-alloyed steel a value of
311 MPa. The elongation shows the same behavior as the
tensile strength for all three steels. The highest value of
44 % was for the raw steel, for the re-melted steel with-
A. DELI] et al.: MECHANICAL PROPERTIES OF THE AUSTENITIC STAINLESS STEEL X15CrNiSi20-12 ...
154 Materiali in tehnologije / Materials and technology 49 (2015) 1, 153–156
Table 3: Results of the tensile tests at temperatures of 20 °C and 740 °C
Tabela 3: Rezultati nateznih preizkusov pri temperaturah 20 °C in 740 °C
Steel condition
X15CrNiSi20-12
20 °C 740 °C
Rp0.2/MPa Rm/MPa A/% HB Rp0.2/MPa Rm/MPa A/%
EN 10095 > 230 500–750 > 28 < 223 – – –
Raw – RM 400 730 44 194 177 281 41
Remelted – MM 376 703 43 190 168 297 36
Realloyed – AM 345 689 41 180 150 311 39
Table 4: Tensile properties at 20 °C for the melted and re-alloyed steel, cast and annealed
Tabela 4: Lastnosti pri nateznem preizkusu pri 20 °C za staljeno in ponovno legirano jeklo, ulito in `arjeno
Steel condition
X15CrNiSi20-12
Remelted steel - MM Realloyed steel - AM
Rp0.2/MPa Rm/MPa A/% Rp0.2/MPa Rm/MPa A/%
Casted 628 813 28 367 696 42
Annealed 403 717 41 346 691 38
Figure 1: Tensile tests at temperatures of 20 °C and 740 °C: a) Rp0.2,
Rm, b) A
Slika 1: Natezni preizkusi pri temperaturah 20 °C in 740 °C: a) Rp0,2,
Rm, b) A
out re-alloying the value was 43 % and for the re-melted
and re-alloyed steel it is 41 %. The elongation at 740 °C
is above the prescribed value for room temperature and it
is not prescribed in the standard.
Figure 2 shows the results of tensile tests for the
re-melted and re-alloyed steel, cast and annealed, and in
Table 4 are the results of the tensile tests for the cast and
annealed, re-melted and re-alloyed steel.
A comparison of the results in Table 4 indicates that
the properties of the re-melted (MM) and the re-alloyed
(AM) steel are reduced after annealing and are much
closer and more balanced than the re-alloyed steel. The
thermal annealing treatment was performed to obtain the
same conditions as were present for the raw steel.
The hardness values are the values prescribed by the
standard for all steels and the maximum hardness is 223
HB. The hardness of the steel is the highest for raw steel,
HB 194, for the re-melted steel it is 190 HB, and it is the
lowest value, 180 HB, for the re-melted and re-alloyed
steel.
2.2 Metallographic analysis
The microstructure was examined with a light
microscope, OLYMPUS, BX60M, according to ASTM E
4077. Figure 3 shows the microstructure of the raw steel
(RM), the re-melted (MM) and the re-alloyed steel (AM)
after quenching with a magnification of 100-times and
500-times is shown after etching with the Kalling
reagent8.
For all three steels the microstructure consists of an
austenitic base and carbide precipitates in the interior
and boundaries of the austenite grains. Figure 3 reveals
the difference in the amount of carbides. In the re-
alloyed material a certain amount of carbides are inside
the grains. This may be regarded as a consequence of the
cooling rate. Considering the location and the size of the
carbide it is assumed that these are complex chromium
carbides9. The composition of complex chromium car-
bides can be found in a number of references.
The examinations were performed with a Philips
XL-30 Scanning Electron Microscope. Figure 4 shows
the analysis of the carbides in the re-alloyed steel.
Figure 4 shows the results of the EDS analysis of the
carbides in the re-alloyed steel that confirms the
presence of chromium carbides.
3 CONCLUSIONS
Based on an analysis of the literature data and the
conducted experimental studies, as well as the results of
A. DELI] et al.: MECHANICAL PROPERTIES OF THE AUSTENITIC STAINLESS STEEL X15CrNiSi20-12 ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 153–156 155
Figure 4: EDS analysis of carbides in re-alloyed steel (AM)
Slika 4: EDS-analiza karbidov v dolegiranem jeklu (AM)
Figure 3: Microstructure of: a), b) the raw steel (RM), c), d) re-melted
steel (MM) and e), f) re-alloyed steel (AM)
Slika 3: Mikrostruktura: a), b) jeklo (RM), c), d) pretaljeno jeklo
(MM), e), f) dolegirano jeklo (AM)
Figure 2: Results of tensile tests at 20 °C in the as-cast and annealed
conditions for: a) re-melted and b) re-alloyed steel
Slika 2: Rezultati nateznih preizkusov pri temperaturi 20 °C v litem in
`arjenem stanju za: a) pretaljeno in b) ponovno legirano jeklo
laboratory tests, the possibility of recycling the revert
scrap of the austenitic stainless steel X15CrNiSi20–12
was investigated.
Re-melted (MM) steel and re-melted steel with
re-alloying (AM) were processed into square rods of 36
mm, and quenched to obtain the same condition and
dimensions as the raw steel (RM). The tensile properties
and microstructure of both melts X15CrNiSi20-12 were
examined with forging and rolling deformation and
subsequent heat treatment of the experimental melts,
approximately equal raw steels were obtained and then
used for comparative laboratory tests. All the technical
parameters of the heating, deformation and heat treat-
ment were fixed.
The mechanical properties were tested at room and
elevated temperatures. The mechanical properties at
room temperature show that the tensile strength of raw
steel (RM), of re-melted steel (MM) and of (AM) are
fairly uniform and are close to the upper prescribed limit
of 750 MPa in the standard EN 10095. The highest
strength is obtained for the raw steel (RM), then the
re-melted steel (MM) and the lowest value of 689 MPa
for the re-melted and re-alloyed steel (AM). The same
trend of results is present for the yield stress Rp0.2 and the
elongation. At elevated temperatures the strength shows
the opposite trend, with the maximum value of 311 MPa
for re-melted and re-alloyed steel (AM) and for the raw
steel (RM) a value of 281 MPa. For the yield stress Rp0.2
the trend is reversed with the largest value of 177 MPa
for the (RM) to a minimum of 150 MPa for the AM. The
high tensile properties are due to the austenitic structure
and the carbide precipitates at the austenite grain boun-
daries.
The experimental results show that with the recycling
process for X15CrNiSi20-12 steel it is possible to obtain
the steel properties required by the standard EN 10095.
The obtained results can be used with recycling techno-
logy in industrial conditions. The proper recycling pro-
cess can ensure competitive products by the use of an
accurate regime of re-melting, processing and heat treat-
ment.
4 REFERENCES
1 EN 10095 Standard, Heat resisting steels and nickel alloys, 1999
2 F. Rakoski, Nichtmetallische Einschlüsse in Stählen, Stahl und Eisen,
114 (1994) 7, 71–76
3 A. Gigovi} - Geki}, Kvantifikacija uticaja alfa-gama obrazuju}ih ele-
menata na mehani~ke osobine i pojavu delta ferita kod nehr|aju}eg
austenitnog ~elika NITRONIC 60, Doktorska disertacija, Fakultet za
metalurgiju i ~elike, Univerzitet u Zenici, 2010
4 M. Oru~, M. Rimac, O. Beganovi}, A. Deli}, New Materials as Base
for Development of Modern Industrial Technologies, 13th Inerna-
tional Research/Expert Conference Trends in the Development of
Machinery and Associated Technology, TMT 2009, Hammamet,
Tunisia, 2009
5 J. D. Verhoeven, Metallurgy of Steel for Bladesmiths & Others who
Heat Treat and Forge Steel, Iowa State University, 2005, 66–68
6 M. Torkar, M. Lamut, A. Millaku, Recycling of steel chips, Mater.
Tehnol., 44 (2010) 5, 289–292
7 ASTM E 407 Standard, Standard Practice for Microetching Metals
and Alloys, 2007
8 H. [uman, Metalografija, Tehnolo{ko metalur{ki fakultet, Univerzitet
u Beogradu, Beograd, 1989
9 H. K. D. H. Bhadeshia, R. W. K. Honeycombe, Steels Microstructure
and Properties, Third Edition, Elsevier Ltd, 2006, 264–266
A. DELI] et al.: MECHANICAL PROPERTIES OF THE AUSTENITIC STAINLESS STEEL X15CrNiSi20-12 ...
156 Materiali in tehnologije / Materials and technology 49 (2015) 1, 153–156
Z. A]IMOVI]-PAVLOVI] et al.: COMPARISON OF REFRACTORY COATINGS BASED ON TALC ...
COMPARISON OF REFRACTORY COATINGS BASED ON TALC,
CORDIERITE, ZIRCON AND MULLITE FILLERS FOR
LOST-FOAM CASTING
PRIMERJAVA OGNJEVZDR@NIH PREMAZOV NA OSNOVI
SMUKCA, KORDIERITA, CIRKONA IN MULITNIH POLNIL ZA
ULIVANJE V FORME Z IZPARLJIVIM MODELOM
Zagorka A}imovi}-Pavlovi}1, Anja Terzi}2, Ljubi{a Andri}3, Marko Pavlovi}1
1University of Belgrade, Faculty of Technology and Metallurgy, Karnegy st. 4, 11000 Belgrade, Serbia
2Institute for Materials Testing, Vojvode Mi{i}a Bl. 43, 11000 Belgrade, Serbia
3Institute for Technology of Nuclear and Other Raw Mineral Materials, Franchet d’Esperey 86, 11000 Belgrade, Serbia
anja.terzic@institutims.rs
Prejem rokopisa – received: 2013-12-23; sprejem za objavo – accepted for publication: 2014-01-15
This study presents the results of an investigation of high-temperature materials – refractory coatings based on different refrac-
tory fillers including talc, cordierite, zircon and mullite applied in the lost-foam casting process. Design and optimization of a
coating composition with controlled rheological properties and synthesis were achieved by applying different coating compo-
nents, suspension agents and fillers and by altering the coating-production procedure. A morphological and microstructural
analysis of the fillers was carried out with scanning electron microscopy. An X-ray diffraction analysis was applied to determine
and monitor the phase-composition changes of the refractory fillers. The particle size and shape were assessed with the PC
software application package OZARIA 2.5. To assess the effects of the applications of individual refractory coatings, a detailed
investigation of the structural and mechanical properties of the mouldings obtained was performed. Highlight was placed on
revealing and analyzing surface and volume defects present on the mouldings. Radiographic moulding tests were carried out
with an X-ray device, SAIFORT type-S200. The attained results are essential for the synthesis of the refractory coatings based
on high-temperature fillers and their applications in the lost-foam casting process for manufacturing the mouldings with
in-advance-set properties.
Keywords: high-temperature materials, refractory coatings, talc, cordierite, zircon, mullite, lost-foam casting
Ta {tudija predstavlja rezultate raziskave visokotemperaturnih materialov – ognjevzdr`nih premazov na osnovi razli~nih
ognjevzdr`nih polnil, smukca, kordierita, cirkona in mulita, uporabljenih pri postopku ulivanja v forme z izparljivim modelom.
Na~rtovanje in optimizacija sestave premazov s kontroliranimi reolo{kimi lastnostmi in sintezo sta bila dose`ena z uporabo
razli~nih komponent premaza, sredstev za suspenzijo in polnil ter s spreminjanjem postopka izdelave premaza. Morfolo{ka in
mikrostrukturna analiza polnil sta bili izvr{eni z vrsti~no elektronsko mikroskopijo. Rentgenska difrakcijska analiza je bila
uporabljena za ugotavljanje in kontrolo sprememb fazne sestave ognjevzdr`nih sestavin. Velikost in oblika zrn sta bili dolo~eni s
PC programskim paketom OZARIA 2.5. Za ugotovitev vpliva posameznega ognjevzdr`nega premaza je bila izvr{ena podrobna
analiza strukturnih in mehanskih lastnosti form. Pozornost je bila namenjena ugotavljanju in analizi povr{ine ter volumenskih
napak v formah. Rentgenski pregledi form so bili opravljeni z rentgensko napravo SAIFORT type-S200. Dobljeni rezultati so
pomembni za sintezo ognjevzdr`nih premazov na osnovi visokotemperaturnih polnil in njihovo uporabo pri postopku ulivanja z
izparljivimi modeli za izdelavo form z vnaprej dolo~enimi lastnostmi.
Klju~ne besede: visokotemperaturni materiali, ognjevzdr`ni premazi, smukec, kordierit, cirkon, mulit, ulivanje z izparljivimi
modeli
1 INTRODUCTION
The most important characteristic of the lost-foam
casting process is the fact that the patterns and the mould
gating made of polymers remain in a cast until the
liquid-metal inflow. In contact with the liquid metal the
polymer is decomposed and evaporated in a relatively
short time. The previously described process is accom-
panied by moulding crystallization1–3. In order to achieve
a high-quality and profitable moulding production using
the lost-foam method, it is necessary to reach a balance
in the system involving evaporable polymer pattern –
liquid metal – refractory coat – sand mould during the
phases of metal inflow, decomposition and evaporation
of the polymer pattern, and formation and solidification
of castings (Figure 1).4 This requires systematic re-
searches of the complex appearances and processes
occurring in the pattern, the metal and the cast as well as
the appearances and processes in the contact zone metal
involving pattern and metal – refractory coating –
sand.5–8
Polymer decomposition is an endothermic process
which starts with a liquid-metal inflow. The pattern-
decomposition kinetics is the function of the liquid-metal
temperature with which the polymer gets in contact.
During the inflow stage, while the metal passes through
the polymer pattern, 70–90 % of the products is liquid.
During the process, the decomposed liquid products are
being pushed towards the casting-cavity upper surface,
afore the liquid-metal front. In the case of a minor
permeability of the refractory coating and the sand for
modelling, these liquid products of the polymer decom-
Materiali in tehnologije / Materials and technology 49 (2015) 1, 157–164 157
UDK 621.74.045:666.76 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(1)157(2015)
position remain in the upper parts of the mouldings
causing surface, subsurface and volume defects. A fur-
ther decomposition of the liquid phase occurs due to
evaporation (the creation of boiling phase layer) forming
a solid residue of the polymer chain, monomer and ben-
zene, small quantities of toluene and ethyl benzene.9–11
Important factors influencing the pattern decompo-
sition and evaporation process, besides the temperature
and the pattern density, are: the type and thickness of the
refractory coating layer covering the evaporable poly-
mer, the type and size of the sand grain for modelling,
the respective permeability of the sand for modelling, the
mouldings and the gating of the mould construction. The
pattern density and permeability of the refractory coating
and the sand mould determine the polymer evaporation
velocity. The velocity of the liquid metal entering the
mould and its contact with the pattern are controlled with
a proper definition of the mould gating12–20.
In order to obtain the mouldings of the required
quality, critical process parameters should be determined
for each particular polymer pattern, as well as the type of
the alloy for casting. This requires long-lasting re-
searches aimed to achieve the optimization of the lost-
foam casting process and to obtain the mouldings with
previously specified properties. In order to correctly
understand the optimization of the lost-foam casting
process, it is necessary to know the various types of
moulds and their properties21–28. Beside this necessity to
obtain the mouldings with the previously specified pro-
perties, the fundamental structure dependence on the
technology should also be determined by applying the
control of the critical process parameters and the control
of the useful moulding properties. Thus, in this work a
special consideration was dedicated to this correlation.
Z. A]IMOVI]-PAVLOVI] et al.: COMPARISON OF REFRACTORY COATINGS BASED ON TALC ...
158 Materiali in tehnologije / Materials and technology 49 (2015) 1, 157–164
Figure 1: Lost-foam-process system balance: liquid metal – refractory
coating – pattern – sand
Slika 1: Postopek z izparljivim modelom – ravnote`je sistema: talina
– ognjevzdr`na obloga – izparljiv model – pesek
Figure 2: Lost-foam-process phases: a) coating-excess removal from the pattern and layer equalizing on the pattern surface, b) coating-layer
drying, c) staged probe prepared for casting, d) easy coating removal from the casts
Slika 2: Faze postopka ulivanja z izparljivim modelom: a) odstranjevanje odve~nega premaza iz modela in izena~enje nanosov na povr{ini
modela, b) su{enje plasti nanosa, c) forma, pripravljena za ulivanje, d) enostavna odstranitev premaza iz ulitka
Table 1: Compositions of the refractory-coating series A, B, C and D
Tabela 1: Sestava ognjevzdr`nih premazov vrste A, B, C in D
Coat type
Refractory filler
Bonding agent Coating maintenancematerial
Solvent for the coating
density, 2 g/cm3Granulation, μm Quantity, %
A
talc bentonite 3.5–5 %
bindal H 8 %
dextrin 0.3–0.5 %
lucel 0.3–0.5 % water30–35 85–88
B
cordierite bentonite 1–3 %
bindal H 5–7 %
dextrin 0.5 %
lucel 0.5 % water35–40 88–90
C
zircon sodium silicate
(water glass) 2–4 %
dextrin 0.5 %
lucel 0.5 % water35–40 90–95
D
mullite bentonite 4 %
bindal H 2–4 %
dextrin 0.5 %
lucel 0.5 % water40–45 88–90
2 MATERIALS AND METHODS
Four series of experiments were carried out to ana-
lyze the possibilities of applying the refractory coatings
based on talc (series mark A), cordierite (series mark B),
zircon (series mark C) and mullite (series mark D)
(Table 1) and the methods for preparing the coating
components were determined. The grinding of the refrac-
tory fillers including talc, cordierite, zircon and mullite
was conducted in a mill with the balls of Cr-Ni steel (the
Retsch-PM4 type), with a capacity of 20 kg/h, a mill
load of 70 % and the grinding time of 45–60 min.
Refractory talc, cordierite, zircon and mullite pow-
ders used for both the coating production and the appli-
cation in the lost-foam casting process depend crucially
on the rheological coating quality and the sedimentary
coating stability. Both, the coating compositions and
their preparation procedures were altered to achieve the
specified coating properties. While applying a refractory
coating on the polymer pattern with the techniques of
immersion into the tank containing the coating, over-
flowing and coating with a brush, a special attention was
given to the coating quality control. The basic criteria for
the quality evaluation of this type of refractory products
were: the pertinence of the application, the drying beha-
viour, the resistance to attrition, the sedimentation and
the penetration.
The experimental parameters of the lost-foam casting
process (the selection of the composition and preparation
of the refractory coats) were determined as presented in
Table 2.
The lost-foam-process phases are showed in Figure
2.
Table 3 shows the process parameters for all the
series of the refractory-coating production, the methods
of the coating application on the patterns and drying.
3 RESULTS AND DISCUSSION
In Figures 3 to 6 the results of the XRD phase ana-
lysis of samples A to D are presented. The results of
XRD show that talc is dominant in sample A (Figure
3)29, cordierite is dominant in sample B where spinel and
Z. A]IMOVI]-PAVLOVI] et al.: COMPARISON OF REFRACTORY COATINGS BASED ON TALC ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 157–164 159
Table 2: Experimental parameters of the lost-foam casting process
Tabela 2: Eksperimentalni parametri postopka ulivanja z izparljivim modelom
Parameter Parameter description
Tested alloy AlSi10Mg
Preparation methods
for the liquid die
-refinement with the compounds of NaCl and KCl with a quantity of 0.1 % of the die mass;
-degasification with briquette C2Cl6 with a quantity of 0.3 % of the die mass;
-modification with sodium with a quantity of 0.05 %.
Casting temperature
(°C)
735; 760; 795
Evaporable
polystyrene pattern
-density (kg/m3): 20; 25
-pattern construction: a plate of 200 mm × 50 mm × 20 mm and a staged probe with different
wall-thickness values (mm): 10; 20; 30; 40; 50
-polystyrene grain size of 1–1.5 mm
Mounting pattern for
the casting
a cluster of four pattern plates set on the central runner gate and a cluster of two staged probes set
on the central runner gate
Gating of the moulds -central runner gate of 40 mm × 40 mm × 400 mm
-ingates of 20 mm × 20 mm × 10 mm, 2 pieces
Dry quartz sand for
the cast production
Grain size (mm): 0.17; 0.26; 0.35
Table 3: Optimum process parameters of the production and preparation of the coating, the lining and coat drying on polystyrene patterns
Tabela 3: Optimalni procesni parametri pri izdelavi premaza, podlage in su{enju premaza na polistirenskem modelu
Parameter Parameter description
Coat density 2 g/cm3
Coat temperature 25 °C
Removing the coat excess from the pattern after pulling it out
of the tank for lining
the coat is being seeped for 5–10 s while the pattern is in the
vertical position and then kept for 5 s at an angle of 45° in
order for the coat layers on the pattern surface to become even
Slow mixing of the coat in the tank while applying the coat on
the pattern - velocity of 1 r/min
Methods of applying the coat on the pattern
-cluster immersion into the tank with the coating
- overflowing
-coating with a brush
Drying
- first layer for 1.5 h
- final layer for 24 h
Thickness of the coat layer on the pattern after drying (mm) 0.5; 1; 1.5
quartz are also present (Figure 4), zircon is dominant in
sample C (Figure 5) and mullite is dominant in sample
D (Figure 6)30.
Figures 7 and 8 show the diagrams of the grain-size
and grain-shape factors of samples A to D. The mean
grain size of the refractory filler is between 35–40 μm
and the mean grain-shape factor is 0.63–0.75, which
means that the grains are elongated rounds (the shape
factor of 0 corresponds to the needle shape and the factor
of 1 corresponds to the circle). SEM microphotographs
show that grains are in fact irregularly shaped with a
large number of angles that strengthen the locking of the
particles in the coating mixture and are preferable for
obtaining a homogeneous coating mixture. This conclu-
sion is in accordance with the previous investigations22,23
emphasizing the fact that the filler particle size and shape
are crucial for the quality of casting coatings. The con-
ditions of the filler particle size, shape and surface
depend on the type of the mill. The roughness of the sur-
faces of filler particles may cause an increased absorp-
Z. A]IMOVI]-PAVLOVI] et al.: COMPARISON OF REFRACTORY COATINGS BASED ON TALC ...
160 Materiali in tehnologije / Materials and technology 49 (2015) 1, 157–164
Figure 5: Diffractogram of the filler based on zircon
Slika 5: Difraktogram polnila na osnovi cirkona
Figure 6: Diffractogram of the filler based on mullite30
Slika 6: Difraktogram polnila na osnovi mulita30
Figure 4: Diffractogram of the filler based on cordierite
Slika 4: Difraktogram polnila na osnovi kordierita
Figure 3: Diffractogram of the filler based on talc
Slika 3: Difraktogram polnila na osnovi smukca
Figure 8: Histogram of the grain-shape factors of samples A to D
Slika 8: Histogram faktorja oblike zrn vzorcev od A do D
Figure 7: Histogram of the grain sizes of samples A to D
Slika 7: Razporeditev velikosti zrn vzorcev od A do D
tion of the liquid metal onto the coating surface, which
deteriorates the coating quality.
Figures 9 to 12 show the shapes of fillers A to D.
The ceramic-powder particles have similar morphologies
but different particle sizes. This is favourable since the
particles of diverse granulations contribute to the forma-
tion of an even and continuous coating layer on the poly-
mer pattern, due to a better harmony among the particles.
For the mullite-sample preparation, a longer grinding
time was required (approximately 90 min) to achieve a
particle granulation of less than 40 μm. Figure 9 shows
that the talc-based filler consists of proper foliate aggre-
gates. The morphology of the cordierite sample (Figure
10) shows that cordierite particles have improper forms
and different sizes. The zircon-powder particles also
have improper forms with a typical shell-like break and
different sizes (Figure 11). The mullite sample (Figure
12) shows improper forms and different sizes of the
particles. Pursuant to the previous work31, it was showed
that an application of the talc-based filler with elongated
and smooth particles contributed to the production of the
coatings of a higher quality, easily adherent to the
patterns and not forming sintered sand faults on the
mouldings. Applying the coating of zircon flour plus
colloidal silica13,17, thin coating layers with a high gas
permeability were obtained, while the silica-amount
increase contributed to the reduction of the metal pene-
tration into the mould. Surveying the references, no
application of cordierite as a refractory filler was
recorded. Referencing earlier works22,23, this research
used different filler granulations that were proved to be
appropriate for obtaining homogenous coating layers and
achieving a better coating adherence to the pattern
surface. Further, the bentonite quantity was reduced, i.e.,
the carboxymethyilcellulose (CMC) quantity increased
causing a reduction in the cracks in dried coating layers,
in accordance with18,21,25.
The control of the critical process parameters for the
refractory-coating production and the control of the
coating properties determine that the coatings of all the
series comply with the conditions for the application in
the lost-foam process. It was determined that the coat-
ings were easily applied on the polymer patterns, being
evenly lined during the overflowing and immersion, were
easy to be coated with a brush, without any marks of
brushing, leakage, drops or clots. After drying, the coat-
ing surface was smooth; the coating layers were of equal
Z. A]IMOVI]-PAVLOVI] et al.: COMPARISON OF REFRACTORY COATINGS BASED ON TALC ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 157–164 161
Figure 12: Microphotograph of the refractory filler based on mullite
Slika 12: Posnetek ognjevzdr`nega polnila na osnovi mulita
Figure 10: Microphotograph of the refractory filler based on cor-
dierite
Slika 10: Posnetek ognjevzdr`nega polnila na osnovi kordierita
Figure 11: Microphotograph of the refractory filler based on zircon
Slika 11: Posnetek ognjevzdr`nega polnila na osnovi cirkona
Figure 9: Microphotograph of the refractory filler based on talc
Slika 9: Posnetek ognjevzdr`nega polnila na osnovi smukca
thickness everywhere on the pattern surface, without
bubbles, crazing, peelings or attrition. The coating quali-
ty and the refractory-filler homogeneity of the coating
depend on the coating preparation. To achieve an even
coat-layer thickness on the pattern surfaces, it is neces-
sary to slowly and constantly mix the coating during its
application on the patterns, maintaining the defined
density (2 g/cm3) and temperature (25 °C) of the coating,
otherwise the coating composition becomes non-homo-
geneous.
In the case of applying a coating layer of a higher
thickness that dried out fast, crazing and tweaking
emerged on dried coating layers, having negative effects
on the mould-surface quality (uneven surface, roughness,
sintered sand, cracks, folds, as seen on Figures 13a and
13c) and there were the risks of a liquid-metal pene-
tration into the cast and the sinking of the cast (Figure
13b). The coating-layer perforation and metal
penetration into the cast were also noticed on the series
of the casts with the coating layers of a lower thickness
(0.5 mm) on the moulding part, with a wall thickness
larger than 40 mm, and applied at higher casting
velocities. When the coating layers of a higher thickness
were applied (above 1 mm) this fault did not emerge on
the mouldings.
To observe the effects of the casting process, evaluate
certain operation phases and analyse the refractory-coat-
ing influence, a visual control of the obtained mouldings
was carried out, testing their structural and mechanical
properties. After pulling the formed "clusters" from the
mould, their surfaces are covered with a coating layer
that is easy to remove so the cleaning is not necessary,
which significantly reduces the production costs. The
refractory coatings of all the series demonstrated positive
effects on the surface quality – shiny and smooth mould-
ing surfaces were obtained. The mouldings were true
copies of the patterns (dimensionally accurate) indicating
a total decomposition and evaporation of the polystyrene
pattern and a satisfactory gating of the moulds. It was
noticed that the lower moulding parts of all the series
have flat and sharp edges, clean and shiny surfaces. In
the cases of some mouldings from the series with the
coating layer of a higher thickness (above 1.5 mm) the
upper moulding surfaces are slightly uneven and folded,
and on certain moulding parts surface roughness also
appears, more often on the mouldings from the series
with a pattern density greater than 20 kg/m3.
The results of the investigations of the moulding
structural and mechanical characteristics were within the
limits predicted by the standards for this type of alloys.
These included the mouldings from the series with the
used polystyrene patterns of up to 20 kg/m3, refractory
coatings of a lower layer thickness, below 1 mm, quartz
sand for modelling with its grain size larger than 0.26
mm, the casting temperature within the limits of
760–780 °C and the casting velocity which enabled even
decomposition and evaporation of polystyrene with a
complete elimination of the gassy products from the
patterns, without the cast falling in or the liquid metal
penetrating into the sand (Figure 14).
On the other hand, the mouldings from the series
with the applied patterns with the densities greater than
20 kg/m3 and the coat thickness greater than 1.5 mm
exhibited subsurface and also volumetric porosity
(Figure 15). This indicates that the reasons for these
types of defects are, primarily, the polystyrene pattern,
followed by the refractory coating and the high casting
velocity. This is in accordance with6,17,24,26,27 showing that
a success of the lost-foam casting process may be
achieved with the polymer patterns with a density of less
than 20 kg/m3 and with an application of thinner layers
of refractory coatings of different compositions based on
silicondioxide31, zircon17, talc31, mica26, with various
components within the coating composition, an applica-
tion of polysaccharide and an increased amount of CMC
(1–5 %). Contrary to our research where pure compo-
nents for the refractory fillers were used, the works32,33
showed that a combination of different refractory fillers,
such as zircon flour, alumina, silica, magnesite, clays,
talc, chromite, mica, together with various suspension
agents for the control of the system rheology, may be
used to attain the significant effects of an increase in
both the coating refractoriness and gas permeability. All
Z. A]IMOVI]-PAVLOVI] et al.: COMPARISON OF REFRACTORY COATINGS BASED ON TALC ...
162 Materiali in tehnologije / Materials and technology 49 (2015) 1, 157–164
Figure 14: Characteristic sample radiogram: a casting without poro-
sity
Slika 14: Zna~ilen radiogram vzorca: ulitek brez poroznosti
Figure 13: Defects on castings: a) casting lumpiness of the surface
reproduced from the pattern, surface porosity, fold, b) metal penetra-
tion into the sand, c) sintered sand on the casting surface
Slika 13: Napake na ulitkih: a) odtis povr{ine iz izparljivega modela,
poroznost povr{ine, guba, b) prodor taline v pesek, c) sintrani pesek na
povr{ini ulitka
of these contributed to achieving an increased quality of
the surfaces of the lost-foam mouldings. To apply this
type of coating on large and complex forms, a flow
coating method was successively developed.
4 CONCLUSION
Utilization of refractory talc, cordierite, zircon and
mullite powders for a coating production and application
in the lost-foam casting process depends, crucially, on
the rheological properties of the coatings, and, respecti-
vely, on the sedimentary coating stability. The presented
research results show the optimum coating compositions
and their preparation procedures with the goal of
achieving positive effects on the moulding quality. Shiny
and smooth moulding surfaces are obtained without
detectable defects and volumetric porosity, folds, sin-
tered sand in the series of casting with polymer patterns
of a lower density, application of refractory coating
layers of a lower thickness and quartz sand for the casts
of a higher permeability. The achieved results of
applying refractory coatings on the mullite base are the
initial ones. Further studies should be done with the goal
to determine the correlations between the coating com-
position and the layer thickness, and the moulding
structural and mechanical characteristics. By developing
refractory coatings on talc, cordierite, zircon and mica
bases and optimizing the technological parameters of the
lost-foam casting process, the mouldings with the
previously defined quality, i.e., desired properties could
be obtained with significantly lower costs with respect to
the mouldings based on sand.
Acknowledgements
This investigation was supported and funded by the
Ministry of Education, Science and Technological
Development of the Republic of Serbia and Projects
33007, 172057 and 45008.
5 REFERENCES
1 R. Monroe, Expandable Pattern Casting, AFS Inc., Des Plaines,
Illinois, USA 1992, 71
2 H. Mae, X. Teng, Y. Bai, T. Wierzbicki, Comparison of ductile frac-
ture properties of aluminum castings: Sand mold vs. metal mold,
International Journal of Solids and Structures, 45 (2008), 1430–1444
3 G. W. Powell, The fractography of casting alloys, Materials Charac-
terization, 33 (1994), 275–293
4 A. Prsti}, Z. A}imovi}-Pavlovi}, M. \uri}, A. Terzi}, Lj. Pavlovi},
Synthesis, Characterization and Application of the Filler Based on
Corundum for Obtaining the Ceramic Coats in Foundry, Metalurgia
International, 17 (2012), 83–87
5 R. E. Moore, Refractories, Structure and Properties, Encyclopedia of
Materials: Science and Technology, Pergamon Press, Oxford 2001,
8079–8099
6 Y. Motoyama, H. Takahashi, Y. Inoue, K. Shinji, M. Yoshida, Deve-
lopment of a device for dynamical measurement of the load on cast-
ing and the contraction of the casting in a sand mold during cooling,
Journal of Materials Processing Technology, 212 (2012), 1399–1405
7 M. Daniel, K. Jarrett New, J. Dantzig, A sand surface element for
efficient modeling of residual stress in castings, Applied Mathe-
matical Modeling, 25 (2001), 825–842
8 D. G. Eskin, L. Katgerman, Contraction behavior of aluminum alloys
during solidification, International Foundry Research, 59 (2007),
8–13
9 L. Y. Pio, S. Sulaiman, A. M. Hamouda, M. Mohamad, H. M.
Ahmad, Grain refinement of LM6 Al–Si alloy sand castings to
enhance mechanical properties, Journal of Materials Processing
Technology, 162–163 (2005), 435–441
10 D. Apelian, G. K. Sigworth, K. R. Whaler, Assessment of Grain
Refinement and Modification of Al–Si Foundry Alloys by Thermal
Analysis, American Foundrymen’s Society, Inc., 1984, 99297–99307
11 U. C. Nwaogu, T. Poulsen, R. K. Stage, C. Bischoff, N. S. Tiedje,
New sol–gel refractory coatings on chemically-bonded sand cores for
foundry applications to improve casting surface quality, Surface and
Coatings Technology, 205 (2011), 4035–4044
12 J. R. Brown, The Foseco Foundryman’s Handbook, Pergamon Press,
Oxford 2000
13 F. L. Pirkle, D. A. Podmeyer, Zircon: Origin and Uses, Transactions,
vol. 292, Society for Mining, Metallurgy and Exploration, Inc., USA
1988, 62
14 B. Schillinger, E. Calzada, C. Eulenkamp, G. Jordan, W. W.
Schmahl, Dehydration of moulding sand in simulated casting process
examined with neutron radiography, Nuclear Instruments and
Methods in Physics Research A, 651 (2011), 312–314
15 J. L. Li, R. S. Chen, W. Ke, Microstructure and mechanical proper-
ties of Mg-Gd-Y-Zr alloy cast by metal mould and lost foam casting,
Transactions of Nonferrous Metals Society of China, 21 (2011),
761–766
16 A. Wróbel, E. Kulej, B. Kucharska, The Effect of Corrosion Wear on
the Recorded Texture of a Clad Al-Si Coating, Archives of Metal-
lurgy and Materials, 4 (2011), 897–902
17 M. Karimian, A. Ourdjini, M. H. Idris, T. Chuan, H. Jafari, Process
Control of Lost Foam Casting using Slurry Viscosity and Dipping
Time, Journal of Applied Sciences, 11 (2011), 3655–3658
18 W. D. Griffiths, P. J. Davies, The permeability of Lost Foam pattern
coatings for Al alloy castings, Journal of Materials Science, 43
(2008), 5441–5447
19 A. Charchi, M. Rezaei, S. Hossainpour, J. Shayegh, S. Falak, Nume-
rical simulation of heat transfer and fluid flow of molten metal in
MMA–St copolymer lost foam casting process, Journal of Materials
Processing Technology, 210 (2010), 2071–2080
Z. A]IMOVI]-PAVLOVI] et al.: COMPARISON OF REFRACTORY COATINGS BASED ON TALC ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 157–164 163
Figure 15: Characteristic sample radiogram: a casting with expressed
porosity, a higher pattern density and a coating layer with a higher
thickness
Slika 15: Zna~ilen radiogram vzorca: ulitek s poroznostjo, ve~jo go-
stoto modela in ve~jo debelino premaza
20 M. Khodai, N. Parvin, Pressure measurement and some observation
in lost foam casting, Journal of Materials Processing Technology,
206 (2008), 1–6
21 F. W. Pursall, Coatings for Moulds and Cores, In: K. Straus (Ed.),
Applied Science in the Casting of Metals, Pergamon Press, Oxford
1970
22 Lj. Trumbulovic, Z. Acimovic-Pavlovic, Z. Gulisija, Lj. Andric,
Correlation of technological parameters and quality of castings
obtained by the EPC method, Materials Letters, 56 (2004),
1726–1731
23 Z. A}imovi}, Lj. Pavlovi}, Lj. Trumbulovi}, Lj. Andri}, M. Sta-
matovi}, Synthesis and characterization of the cordierite ceramics
from nonstandard raw materials for application in foundry, Materials
Letters, 57 (2003), 2651–2656
24 T. Hübert, S. Svoboda, B. Oertel, Wear Resistant Alumina Coatings
Produced by a Sol-Gel Process, Surface and Coatings Technology,
201 (2006), 487–491
25 W. Gasior, P. Fina, Z. Moser, Modeling of the Thermodynamic
Properties of Liquid Fe-Ni and Fe-Co Alloys From the Surface
Tension Data, Archives of Metallurgy and Materials, 1 (2011), 13–23
26 D. Gan, S. Lu, C. Song, Z. Wang, Physical Properties of Poly(ether
ketone ketone)/Mica Composites: Effect of Filler Content, Materials
Letters, 48 (2001) 5, 299–302
27 P. L. Jain, Principle of Foundry Technology, 4th Edition, McGraw-
Hill, New Delhi 2006
28 D. A. Caulk, A foam melting model for lost foam casting of alumi-
num, International Journal of Heat and Mass Transfer, 49 (2006),
2124–2136
29 Lj. Andri}, A. Terzi}, Z. A}imovi}-Pavlovi}, Lj. Pavlovi}, M. Petrov,
Comparative Analysis of Process Parameters of Talc Mechanical
Activation in Centrifugal and Attrition Mill, Physicochemical Prob-
lems of Mineral Processing, 50 (2013), 433–452
30 Z. A}imovi}, A. Terzi}, Lj. Andri}, Lj. Pavlovi}, M. Pavlovi}, Syn-
thesizing a new type of mullite lining, Mater. Tehnol., 47 (2013) 6,
777–780
31 M. Yekeler, U. Ulusoy, C. Hicyilmaz, Effect of particle shape and
roughness of talc mineral ground by different mills on the wettability
and floatability, Powder Technology, 140 (2004), 68–78
32 U. C. Nwaogu, N. S. Tiedje, Foundry Coating Technology: A Re-
view, Materials Science and Applications, 2 (2011), 1143–1160
33 D. Kascheev, N. Yu. Novozhilov, E. V. Tsarevskii, V. A. Perepelitsyn,
V. A. Ryabin, N. F. Seliverstov, Refractory Coatings for Foundry
Moulds and Cores, Journal of Refractories and Ceramics, 23 (1982),
36–139
Z. A]IMOVI]-PAVLOVI] et al.: COMPARISON OF REFRACTORY COATINGS BASED ON TALC ...
164 Materiali in tehnologije / Materials and technology 49 (2015) 1, 157–164
V. LAZI] et al.: SELECTION OF THE MOST APPROPRIATE WELDING TECHNOLOGY FOR HARDFACING ...
SELECTION OF THE MOST APPROPRIATE WELDING
TECHNOLOGY FOR HARDFACING OF BUCKET TEETH
IZBIRA NAJBOLJ PRIMERNE TEHNOLOGIJE TRDEGA
NAVARJANJA ZOBA ZAJEMALKE
Vuki} Lazi}1, Aleksandar Sedmak2, Ru`ica R. Nikoli}1, Milan Mutavd`i}3,
Srbislav Aleksandrovi}1, Bo`idar Krsti}1, Dragan Milosavljevi}1
1University of Kragujevac, Faculty of Engineering, Sestre Janji} 6, 34000 Kragujevac, Serbia
2University of Belgrade, Faculty of Mechanical Engineering, Kraljice Marije 16, 11000 Belgrade, Serbia
3Company for Road Construction Kragujevac, Tanaska Raji}a 16, 34000 Kragujevac, Serbia
vlazic@kg.ac.rs
Prejem rokopisa – received: 2014-01-02; sprejem za objavo – accepted for publication: 2014-01-15
A possibility of extending the service life of the working parts of construction machinery with particular attention to hardfacing
of loader bucket teeth was investigated. In the first part of this paper the tribological processes typical for this machinery is
analysed. Worn excavator parts are made of conditionally weldable cast steel that requires a special hardfacing technology, so
numerous investigations were performed to obtain the most appropriate technology. In the experimental part of the paper, the
selection of the optimum hardfacing technology for bucket teeth and the procedure of the manual arc hardfacing are presented.
The samples were first hardfaced using different techniques and technologies and then the microstructure and microhardness of
characteristic hardfaced layers were studied. Specially prepared samples were used for tribological investigations. The results of
experimental investigations enabled the selection of the most suitable hardfacing technology and its application to real parts.
The bucket teeth, with their hardfaced layers applied vertically, horizontally or in a honeycomb pattern were mounted onto a
loader bucket, alternated with the new non-hardfaced teeth and their performance during the operation was regularly monitored.
After a certain period, the degrees of the wear for the non-hardfaced and differently hardfaced teeth were measured. Taking into
account both technical and economic factors, the most suitable hardfacing technology was determined.
Keywords: hardfacing, wear, loader bucket teeth, hardness, microstructure, friction coefficient
Izvr{ena je bila analiza podalj{anja zdr`ljivosti delov na gradbenih strojih s posebno pozornostjo na trdem navarjanju na zobeh
zajemalke. V prvem delu tega ~lanka so analizirani tribolo{ki procesi, ki se pojavijo pri teh strojih. Pri kopa~ih so deli, ki se
obrabljajo, izdelani iz pogojno varivega litega `eleza, ki zahteva posebno tehnologijo nana{anja, zato so bile izvr{ene {tevilne
raziskave, da bi dobili najprimernej{o tehnologijo. V eksperimentalnem delu ~lanka je predstavljena optimalna tehnologija
nana{anja trdih slojev na zobe zajemalke in predstavljen je postopek za ro~no varjenje. Na vzorce so bili naneseni razli~ni trdi
sloji z razli~nimi tehnikami in tehnologijami in preu~evane so bile mikrostrukture in mikrotrdote trdih nanosov. Preiskani so bili
posebno pripravljeni vzorci za tribolo{ke preizkuse. Rezultati eksperimentov so omogo~ili izbiro najprimernej{e tehnologije
nana{anja trdega sloja in njeno uporabo na realnih delih. Zobje zajemalke z vertikalnim, horizontalnim ali satastim nanosom
trdega sloja so bili name{~eni skupaj z novimi, neobdelanimi zobmi na nakladalnik in redno je bilo spremljano njihovo vedenje
med delom. Po dolo~enem ~asu je bila izmerjena obraba zob z razli~nimi nanosi in obraba neobdelanih zob. Z upo{tevanjem
tehni~nih in ekonomskih faktorjev je bila dolo~ena najprimernej{a tehnologija nana{anja trdih plasti.
Klju~ne besede: trdo navarjanje, obraba, zobje zajemalke nakladalnika, trdota, mikrostruktura, koeficient trenja
1 INTRODUCTION
During operation, certain parts of the road-construc-
tion machinery are exposed to different abrasive
materials that cause most of the damage of the parts in
direct contact with the stone aggregate causing abrasive
wear. Hard and sharp-edged particles of stone materials
are highly abrasive and they damage the working parts of
bucket teeth.
The working parts exposed to the abrasive wear due
to occasional medium-impact loads include: the bucket
teeth of loaders, trenchers and excavators, the blades of
concrete- and asphalt-cutting devices, the blades and rip-
pers of bulldozers and graders, the leading rings and
blades of rock-drill bits, the spindles of screw conveyors,
etc. The greatest abrasive wear occurs on bucket teeth.
For that reason, our experimental investigations were
conducted on the loader bucket teeth.
The studies of the causes for the damage of some
parts of machines and devices revealed that in more than
50 % of the cases the damage was the result of tribo-
logical processes under more or less regular operating
conditions1–6. The damaged parts can be either replaced
with new ones or, in most cases, they can be hardfaced.
Both reparatory and production hardfacing reduce
downtime and costs because new parts are expensive.
Hardfacing is economically justified especially in the
cases of large-sized parts or when there are many equal
parts. However, there are occasions when reparatory
hardfacing has to be performed regardless of the costs,
for example, when unique machines and devices have to
be repaired or no spare parts are available7–11.
2 DAMAGE DUE TO TRIBOLOGICAL CAUSES
Wear is generally considered to be the result of
friction or a combination of friction, thermal, chemical,
electrochemical and other factors on the elements of a
tribo-mechanical system. When studying friction, one
Materiali in tehnologije / Materials and technology 49 (2015) 1, 165–172 165
UDK 621.791.92:539.375.6:539.92 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(1)165(2015)
has to consider the factors dominant in each particular
case, such as the material, working-surface properties,
the contact-surface quality and properties, properties of
the medium between the contact surfaces, characteristics
of the relative motion between the working surfaces, the
load, the temperature, the quantity and properties of the
particles produced due to the wear, etc.1–6
During the surface contact between two tribo-ele-
ments, elastic and plastic deformations occur. They
depend on the load intensity, friction conditions, material
mechanical properties and micro-geometry of the contact
surfaces. When two rough surfaces interact, a momen-
tary loss of the contact may occur due to the micro-
roughness caused by their elastic and plastic deforma-
tions. The process of micro-wear involves a plurality of
such micro-deformations and a destruction of the surface
roughness peaks. Experimental investigations1–6 showed
that the process of the final wear is in fact a fatigue
process. Different authors give different classifications of
the wear, but all of them are based on the way the contact
between two bodies is realized. Therefore, the following
types of wear can be distinguished: adhesive, abrasive,
erosive, fatigue, cavitation, vibrational and corrosive
wear. Since the aim of this research was to study and
estimate also the filler materials used for hardfacing the
parts exposed to the abrasive wear, the abrasive wear
under moderate to medium impact loads was considered.
According to several authors1–6, abrasive wear
accounts for approximately half of all the wear. The
elements most exposed to abrasive wear are the parts of
construction, mining and agricultural machinery, ele-
ments of transport devices, working parts of the equip-
ment in metallurgy, some parts of tool machines, parts of
railway and tram equipment, impellers of hydraulic and
gas turbines, oil-well drilling bits, and parts of the equip-
ment for sandblasting.
Experimental investigations showed that the resistan-
ce to abrasive wear is linearly dependent on the metal
mechanical properties1,4. Therefore, the wear behaviour
of a metal can be predicted on the basis of its mechanical
properties, primarily the hardness. The penetration depth
of foreign particles in the coupled machine elements is
inversely proportional to the hardness of the surface
layers. However, the resistance to wear of the alloys of
the same hardness can vary depending on the chemical
composition and structure of the alloy. Hence, the wear
is influenced not only by the hardness, but also by the
shape, the size and the structural component arrange-
ment. There is no agreement on the most favourable type
of structure in relation to the resistance to abrasive wear.
Some authors believe that the austenite-carbide is the
most favourable, while others prefer the martensite-car-
bide structure1–6. This disagreement stems from a great
variety of abrasive-wear types and a wide range of work-
ing conditions1.
Abrasive materials penetrate the metal-surface layers
causing surface damage. How deep the abrasive particles
penetrate into the metal depends on the shape and hard-
ness of the grains. For example, sharp-edged particles of
relatively soft abrasive materials can cause greater dama-
ge than rounded particles of hard abrasive materials.
Investigations1 showed that the bigger the size of the
particles, the more abrasive they are. Other factors of
influence include the length of the wear path, the specific
pressure, the relative humidity and the chemical environ-
ment.
The wear-resistance level is influenced by each
structural component of steel, but the level of influence
depends on its hardness and its presence in the structure.
Abrasive wear primarily depends on the possibility of
abrasive materials to penetrate into the surface of steel
and on the strength of the bonds between the structural
components at the metal grain boundaries. This means
that the quenched and tempered steel is more resistant
than the steel with a ferrite-pearlite structure. Experi-
ments showed that pure martensite structures, even those
of a lower hardness, exhibit a greater resistance than
martensite-carbide structures. Furthermore, if the mar-
tensite amount is reduced at the expense of the retained
austenite in a martensite-carbide structure, the resistance
to wear is increased despite a decrease in the hardness. It
is the austenite-carbide structure that has the highest
resistance to wear and not the martensite-carbide one, as
might have been expected because of the hardness. This
is due to the fact that the grain-boundary bonds are stron-
ger in austenite-carbide structures than in a combination
of martensite and carbide. In other words, abrasive parti-
cles pull carbides more easily out of a martensite matrix
than out of an austenite base, whose crystal-lattice para-
meter is similar to the carbide-lattice parameter1–6.
The change in operating conditions brings about a
change in the intensity and mechanisms of abrasive wear.
These conditions include the properties of abrasive
materials, the shape and size of grains, the type of bonds,
the specific pressure at the contact surfaces, the relative
sliding speed, the wear-path length, the humidity, the
chemical aggressiveness of the working environment,
etc.
3 SELECTION OF THE PROCEDURE, FILLER
MATERIAL AND HARDFACING TECHNOLOGY
The role of bucket teeth (Figure 1a) is to separate
and crush stone material and to protect the bucket against
wear. The number and shape of teeth depends on the
application (trench digging, gravel digging, material
loading, etc.) and the size of the bucket, varying from
three (trencher buckets) to fifteen teeth (buckets of
bigger loaders). They can be installed in two ways:
bolted directly to a bucket or their holder can be welded
to a bucket and the teeth are bolted onto the holder. Teeth
are made of cast steel or cast iron; their mass ranges
from 3 kg to 15 kg per piece.
V. LAZI] et al.: SELECTION OF THE MOST APPROPRIATE WELDING TECHNOLOGY FOR HARDFACING ...
166 Materiali in tehnologije / Materials and technology 49 (2015) 1, 165–172
Two problems were noticed during the operation of
the bucket teeth: the tougher and less hard teeth are
subjected to plastic deformation and intensive wear,
while the teeth made of brittle materials of high hardness
break at higher impact loads (Figure 1b). A special
problem is encountered with the construction machines
operating in quarries or open-pit mines of ore and coal,
where a piece of a broken tooth can get into the crushers
and cause damage. Broken bucket teeth can also get into
the equipment used for asphalt and concrete production
and cause major damage and downtime.
An analysis of the chemical composition of the teeth
material showed they were made of cast steel ^L3134
(JUS) (Table 1). A broken tooth was cut to be used for
an investigation of the hardness and microstructure of the
base material. The hardness ranged from 340 HV1 to
420 HV1, while the microstructure was martensite with
retained austenite. This microstructure was not resistant
enough because the teeth wore out quickly, especially
under the conditions of severe abrasion. For that reason,
instead of installing new teeth, it was decided to try to
extend the service life of both worn and new teeth with
reparatory and production hardfacing.
Hardfacing was performed with the manual arc-weld-
ing method using a high-alloy rutile electrode with the
properties from Table 2, as #1. According to the manu-
facturer, the layers hardfaced with this electrode exhibit a
high resistance to intensive abrasive wear and can
withstand moderate impact loads during operation. The
layers are very hard and can be grinded. This electrode is
especially suited for hardfacing the parts of the machi-
nery exposed to the metal-to-mineral wear. They include
bulldozer blades, excavator bucket teeth, excavator sho-
vels, parts of conveyors and crushing machines for
various applications, blades and parts of mixers, etc. The
reparatory hardfacing of these parts should be produced
with the interlayer deposition using an electrode (Table
2). The hardfacing parameters, given in Table 3, were
chosen in accordance with the recommendations in11–15.
V. LAZI] et al.: SELECTION OF THE MOST APPROPRIATE WELDING TECHNOLOGY FOR HARDFACING ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 165–172 167
Figure 1: Loader bucket teeth: a) new parts, b) damaged part
Slika 1: Zobje zajemalke: a) novi deli, b) po{kodovani del
Table 1: Chemical composition, mechanical properties and notation of steel ^L31345,8
Tabela 1: Kemijska sestava, mehanske lastnosti in oznaka jekla ^L31345,8
Chemical composition, w/% Mechanical properties Relation toother standards
C Si Mn P S RmMPa
ReH
MPa
A5
%
Hardness
HV1 Microstructure DIN
Prescribed 0.45 0.50 1.80 0.040 0.040 780–930 390 7 340–430* Martensite-bainite andretained austenite
GS-36MnS
Analyzed 0.35 0.20 1.65 0.020 0.020 – – – 340–420 Martensite-bainite andretained austenite
Table 2: Properties of filler materials5,8
Tabela 2: Lastnosti dodajnega materiala5,8
Electrode notation Chemical composition, w/%
Current
type
Mechanical
properties of
the hardfaced
layer
Application
FIPROM Jesenice DIN C Si Mn Cr Ni
1 ABRADUR58
DIN 8555
E 10-UM-60-GR
3.6 – – 32 –
~
= (+)
57–62 HRC
For hardfacing the parts
exposed to severe abrasive
wear from minerals in the
cold state
2 INOX B18/8/6
DIN 8556
E 18 8 6 Mn B 20+
0.12 0.8 7 19 9 = (+) KV > 80 J Interlayer austenite basicelectrode
Table 3: Parameters of hardfacing with the MMA welding method
Tabela 3: Parametri trdega navarjanja z metodo varjenja MMA
Electrode notation Electrode core
diameter
de/mm
Welding current
I/A
Voltage
U/V
Hardfacing
speed
vz/(cm/s)
Input heat
ql/(J/cm)FIPROM Jesenice DIN
1 ABRADUR 58 E 10-UM-60-GR 3.25 130 25 0.124 20968
2 INOX B 18/8/6 E 18 8 6 Mn B 20+ 3.25 100 24 0.136 14118
4 EXPERIMENTAL INVESTIGATIONS ON
MODELS AND BUCKET TEETH
4.1 Model investigations
4.1.1 Investigations of the hardness and microstruc-
ture of the hardfaced-layer zones
The aim of these experimental investigations was to
establish the optimum hardfacing technology. The speci-
mens were hardfaced in a single pass or several passes
(layers) (Figures 2a to 2c), either with or without pre-
heating. Hardfaced samples were cut into metallographic
samples for tribological investigations (Figure 2d). Their
hardness was measured along three different directions
(see detail "A" in Figure 2d) and the microstructure of
characteristic hardfaced-layer zones was estimated. The
hardfaced-layer hardness was 551 to 742 HV1, while the
microstructure was estimated as martensite-ledeburite
with retained austenite and excreted carbides at the grain
boundaries. The microstructure of the heat-affected zone
(HAZ) was martensite with transitions into interphase
structures, and the hardness was from 465 HV1 to 613
HV1. The microstructure of the interlayer was auste-
nite-carbide, while its hardness was about 482 HV1. The
hardness distribution and the microstructures of the cha-
racteristic zones are shown in Figure 3 for the three-
layer hardfacing and in Figure 4 for the two-layer hard-
facing with preheating5–8.
4.1.2 Tribological investigations
Tribological investigations were performed for a
block-on-disk contact, on a TPD-93 tribometer (Figure
5) made at the Faculty of Engineering in Kragujevac.
The aim of these investigations was to evaluate the
resistance to wear of the base materials and deposited
layers. Three prismatic samples (two from the hardfaced
layer and one from the base material) were prepared for
tribological investigations (6.5 mm × 15 mm × 10 mm).
V. LAZI] et al.: SELECTION OF THE MOST APPROPRIATE WELDING TECHNOLOGY FOR HARDFACING ...
168 Materiali in tehnologije / Materials and technology 49 (2015) 1, 165–172
Figure 4: Hardness distribution and microstructures of characteristic
hardfaced-layer zones (two layers, Tp = 200 °C)
Slika 4: Razporeditev trdote in mikrostruktura podro~ij zna~ilnih trdih
navarov (dva sloja, Tp = 200 °C)
Figure 2: Order of hardfaced-layer deposition: a) 1st layer, b) 2nd layer, c) 3rd layer, d) metallographic sample (block)
Slika 2: Zaporedje trdih navarov: a) 1. sloj, b) 2. sloj, c) 3. sloj, d) vzorec za metalografijo
Figure 3: Hardness distribution and microstructures of characteristic
hardfaced-layer zones (interlayer + two layers, To = 20 °C)
Slika 3: Razporeditev trdote in mikrostruktura podro~ja trdega navara
(vmesna plast + dva sloja, To = 20 °C)
During the investigations, the line block-on-disk contact
was realized. The external variables of the tested samples
were the contact forces, the sliding speed and the
lubricant (motor oil GLX 2 SAE 15-W-40)4,5,8.
Prior to the investigations, the surface topography of
blocks and discs was assessed with the computer
measuring system Talysurf 6. The force FN = 300 N and
sliding speed vkl = 1 m/s were adopted. During the con-
tact time of  60 min, the friction-coefficient variation
was registered (Figure 6). When the contact was termi-
nated, the surface topography was assessed again and the
wear-scar widths of the blocks were measured (Figure
7). The wear-scar width was measured using a universal
microscope UIM-21, with a magnification of 50 times4–9.
Based on the results obtained with tribological inve-
stigations, the macroscopic and microscopic damages of
the damaged blocks from Figure 8 were assessed.
4.2 Hardfacing of real parts with different methods of
hardfaced-layer deposition
The measurements of the wear-scar widths showed
that hardfaced layers have a significantly higher resi-
stance to wear (especially those hardfaced without pre-
heating) than the base material. This illustrates the
importance of selecting the right hardfacing technology
and filler materials for excavator/loader bucket teeth.
After the samples had been hardfaced and tribologi-
cally investigated, the optimum technology was selected
and applied to the real parts. Taking into consideration
the limitations imposed by the base-material thickness,
the hardfaced-layer height and electrode diameter, the
V. LAZI] et al.: SELECTION OF THE MOST APPROPRIATE WELDING TECHNOLOGY FOR HARDFACING ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 165–172 169
Figure 7: Wear scar – block No. 3 (after the contact of 60 min)
Slika 7: Brazgotina zaradi obrabe – blok {t. 3 (po stiku 60 min)
Figure 5: Tribometer TPD-93 and other measuring equipment for the
performance of tribological tests
Slika 5: Tribometer TPD-93 in druga merilna oprema za izvajanje
tribolo{kih preizkusov
Figure 8: Wear of block Nos. 1, 2 and 3
Slika 8: Obraba na bloku {t. 1, 2 in 3
Figure 6: Friction-coefficient variation during the contact of 60 min
(blocks 1, 2 and 3)
Slika 6: Spreminjanje koeficienta trenja med stikom 60 min (blok 1, 2
in 3)
V. LAZI] et al.: SELECTION OF THE MOST APPROPRIATE WELDING TECHNOLOGY FOR HARDFACING ...
170 Materiali in tehnologije / Materials and technology 49 (2015) 1, 165–172
Figure 12: Hardfaced teeth after 3200 h of operation: a) front surface, b) back surface
Slika 12: Zob s trdim navarom po 3200 h dela: a) zgornja povr{ina, b) spodnja povr{ina
Figure 11: Hardfaced teeth after 1600 h of operation: a) front surface, b) back surface
Slika 11: Zob s trdim navarom po 1600 h dela: a) zgornja povr{ina, b) spodnja povr{ina
Figure 10: Bucket-tooth mounting sequence: a) front surface, b) back surface, c) teeth surface after 360 h of operation
Slika 10: Zaporedje namestitve zob zajemalke: a) zgornja povr{ina, b) spodnja povr{ina, c) zobje po 360 h dela
Figure 9: Hardfaced teeth with the layers deposited: a) longitudinally, b) horizontally and c) in a honeycomb pattern
Slika 9: Zob s trdim navarom, polo`enim: a) vzdol`no, b) vodoravno in c) satasto
three-layer hardfacing (with an interlayer) was chosen.
Hardfaced layers were deposited at the front and back
sides – vertically over a tooth (Figure 9a), horizontally
across a tooth (Figure 9b) or in a honeycomb pattern
(Figure 9c). The width of the hardfaced layers was from
10 mm to 12 mm and the height was about 4 mm. Then,
the hardfaced teeth were installed onto the central part of
the bucket with the most intensive abrasive wear (Figure
10a, teeth 4, 5 and 6). The end teeth mounted on both
sides were new (1, 2, 3 and 7, 8, 9).
The wear area was monitored and inspected after
360 h (Figures 10a to 10c), after 1600 h (Figures 11a
and 11b) and finally after 3200 h of operation (Figures
12a and 12b)5,8.
5 DISCUSSION
The teeth mass was measured before they were
installed. After the teeth were used under real operating
conditions for about 3200 hours, they were dismounted
and their mass was measured again (Table 4). Figure 13
gives the diagrams of mass losses for the new (non-hard-
faced) and hardfaced teeth.
The examination of the damaged-tooth geometries
and the measurements of the material depth in the wedge
parts of the teeth revealed that the wear degree of the
lower (back) surfaces was about 3 times higher than that
of the upper (front) surfaces of the teeth. This can be
ascribed to the friction of the back surfaces in contact
with the abrasive material5,8.
The greatest resistance to the abrasive wear (loading
of crushed stone material) was exhibited by the teeth
with the hardfaced layers applied vertically, followed by
the teeth with the hardfaced layers deposited in a honey-
comb structure, while the teeth with the horizontal
hardfaced layers showed the least wear resistance. The
best hardfacing patterns for the other conditions of the
abrasive wear require a further investigation5,8.
On the basis of the knowledge acquired in practice
and from the literature, the criterion for a worn-tooth
replacement was established, stating that teeth should be
replaced after 6400 h of operation (two seasons) or when
the mass loss is about 20 %5,8,16,17. Considering the last
rows of Table 4 and Figure 13, one can notice that the
hardfaced teeth were worn significantly less than the new
ones, by about 2 to 4 times. Thus, one can conclude that
the expected service life of the hardfaced teeth would be
2–4 times longer than the service life of the
non-hardfaced teeth, depending on the applied hard-
facing method.
The techno-economic analysis4,5,7,15,16 showed that the
production and reparatory hardfacing of the working
parts of construction machinery is cost-effective. Espe-
cially the production hardfacing of parts was studied in
detail. Taking into consideration the costs for repair, the
costs for new parts and the service life, it can be con-
cluded that the savings reach up to 225 %, while for the
repair hardfacing, they reach up to, or even exceed,
300 %5,8.
6 CONCLUSION
Through the experimental investigations of the
hardness and microstructure of hardfaced-layer characte-
ristic zones, as well as tribological investigations of
hardfaced layers, the optimum hardfacing technology
was determined.
Based on these models and other experimental inve-
stigations, as well as the tests of real parts in operating
conditions, it was concluded that the optimum produc-
V. LAZI] et al.: SELECTION OF THE MOST APPROPRIATE WELDING TECHNOLOGY FOR HARDFACING ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 165–172 171
Table 4: Results of mass measurements before and after the teeth were put into operation for 3200 h
Tabela 4: Merjenje mase zob pred 3200 h dela in po njem
Tooth type New (non-hardfaced) teeth Hardfaced teeth with the layersdeposited* New (non-hardfaced) teeth
Tooth number 1 2 3 4 5 6 7 8 9
Tooth mass before the
operation, kg 8.60 8.58 8.62 9.02 9.05 9.62 8.67 8.60 8.64
Tooth mass after the
operation, kg 7.75 7.72 7.84 8.80 8.65 9.22 7.62 7.74 7.82
Mass loss, kg 0.85 0.86 0.78 0.22 0.50 0.40 0.95 0.86 0.82
Mass-loss percentage,
% 9.88 10.02 11.09 2.44 5.46 4.16 11.08 10.00 9.49
*Tooth #4 hardfaced vertically; tooth #5 hardfaced horizontally; tooth #6 hardfaced in a honeycomb structure
Figure 13: Diagram of mass losses for the non-hardfaced and hard-
faced teeth
Slika 13: Prikaz izgube mase za neobdelane in trdo navarjene zobe
tion-hardfacing technology should consist of the follow-
ing steps:
• Proper selection of the filler material (the right type
of electrodes);
• Selection of the hardfacing parameters (welding
current, voltage, welding speed, input heat);
• Number and order of deposited layers;
• Direction of the hardfaced-layer deposition;
• Thermal treatment prior to hardfacing (preheating) or
deposition of a buffer interlayer.
The right choice of the hardfacing technology and its
proper application ensure numerous advantages and
benefits over the installation of new non-hardfaced parts.
Hardfacing extends the service life by two to four times,
increases productivity, shortens downtime, reduces the
inventory of spare parts, i.e., reduces the production
costs in general.
Although hardfacing represents an almost unique
process, requiring the technology to be modified for each
particular part, the general procedure applicable to
similar parts of construction, mining and agricultural
machineries has been determined in this paper. This
technology can be applied to both manufacturing new
teeth (or parts in general) and repairing the worn ones.
The techno-economic benefits are the same.
It has also been shown that the choice of the hard-
facing technology is closely related to the complex
procedure of hardfaced-layer quality control. This means
that this type of work can be performed only in special-
ized facilities with adequate equipment and skilled staff.
In other words, successful hardfacing can be performed
only by expert teams specialized in the technical system
maintenance.
7 REFERENCES
1 S. Markovi}, Lj. Milovi}, A. Marinkovi}, T. Lazovi}, Tribological
Aspect of Selecting Filler Metal for Repair Surfacing of Gears by
Hard facing, Structural Integrity and life, 11 (2011) 2, 127–130
2 J. Dzubinski, A. Klimpel, Hard facing and thermal spraying, WNT,
Warsaw 1983, 7–19
3 M. Gierzynska, Friction life and lubrication in plastic metal process-
ing, WNT, Warsaw 1983, 35–39
4 V. Lazi}, Optimization of hard facing process from the aspect of
tribological characteristics of the hard faced layers and residual
stresses, PhD thesis, Faculty of Mechanical Engineering, Kragujevac,
2001
5 M. Mutavdzi}, Reparatory hard facing of machine parts and devices
in civil engineering, MS Thesis, Faculty of Mechanical Engineering,
Kragujevac, 2007
6 L. Gusel, R. Rudolf, N. Rom~evi}, B. Buchmeister, Genetic pro-
gramming approach for modelling of tensile strength of cold drawn
material, Optoelectron. Adv. Mater. Rapid Commun., 6 (2012) 3/4,
446–450
7 B. Nedeljkovi}, M. Babi}, M. Mutavd`i}, N. Ratkovi}, S. Aleksan-
drovi}, R. Nikoli}, V. Lazi}, Reparatory hard facing of the rotational
device knives for terrain leveling, Journal of the Balkan Tribological
Association, 16 (2010) 1, 46–75
8 M. Jovanovi}, V. Lazi}, M. Mutavd`i}, D. Adamovi}, Selection of
optimal technology for reparatory hard facing of bucket teeth, Weld-
ing and welded structures, 50 (2005) 1, 11–20
9 V. Lazi}, M. Mutavd`i}, D. Milosavljevi}, S. Aleksandrovi}, B. Ne-
deljkovi}, P. Marinkovi}, R. ^uki}, Selection of the most appropriate
technology of reparatory hard facing of working parts on universal
construction machinery, Tribology in Industry, 33 (2011) 1, 18–27
10 K. M. Mashloosh, T. S. Eyre, Abrasive wear and its application to
digger teeth, Tribology International, 18 (1985) 5, 259–266
11 J. D. Verhoeven, A. H. Pendray, H. F. Clark, Wear tests of steel knife
blades, Wear, 265 (2008) 7/8, 1093–1099
12 R. F. Smart, J. C. Moore, Materials selection for wear resistance,
Wear, 56 (1979) 1, 55–67
13 V. Lazi}, M. Jovanovi}, D. Milosavljevi}, M. Mutavd`i}, R. ^uki},
Choosing of the most suitable technology of hard facing of mixer
blades used in asphalt bases, Tribology in Industry, 30 (2008) 1–2,
3–10
14 Standards: JUS, ASTM, DIN, EN
15 Catalogues of filler material producers: S@ Fiprom Jesenice-Slove-
nia, Lincoln Electric-USA, Esab-Sweden
16 R. Wasserman, How to save millions by reducing inventories of spare
parts, 1st ed., Eutectic-Castolin Institute for the Advancement of
Maintenance and Repair Welding Techniques, New York 1971,
11–23
17 R. ^uki}, Techno-economic analysis of production and reparatory
hard facing of various parts of machine systems, PhD thesis, Uni-
versity of Belgrade, 2010
V. LAZI] et al.: SELECTION OF THE MOST APPROPRIATE WELDING TECHNOLOGY FOR HARDFACING ...
172 Materiali in tehnologije / Materials and technology 49 (2015) 1, 165–172
Z. SAMARD@IJA et al.: CHARACTERIZATION OF TiO2 NANOPARTICLES WITH HIGH-RESOLUTION FEG ...
CHARACTERIZATION OF TiO2 NANOPARTICLES WITH
HIGH-RESOLUTION FEG SCANNING ELECTRON MICROSCOPY
KARAKTERIZACIJA NANODELCEV TiO2 Z
VISOKOLO^LJIVOSTNO FEG VRSTI^NO ELEKTRONSKO
MIKROSKOPIJO
Zoran Samard`ija1, Domen Lapornik2, Ksenija Gradi{ek2, Dejan Verhov{ek2,
Miran ^eh1
1Jo`ef Stefan Institute, Department for Nanostructured Materials, Jamova cesta 39, 1000 Ljubljana, Slovenia
2Cinkarna Celje, d. d., Metalur{ko-kemi~na industrija Celje, Kidri~eva 26, 3001 Celje, Slovenia
zoran.samardzija@ijs.si
Prejem rokopisa – received: 2014-01-16; sprejem za objavo – accepted for publication: 2014-03-04
Ultrafine titanium dioxide powder (UF-TiO2) has many applications as a result of its nanometer particle size and
semiconducting electric properties. An upgraded method for synthesis of UF-TiO2 was developed, with the ability to control the
characteristics of the nanoparticles, i.e. their size, shape and crystal structure. In order to determine the correlation between the
various synthesis parameters and the final form of the TiO2 nanoparticles, their size and morphology were analyzed using
high-resolution field-emission-gun scanning electron microscopy (FEGSEM). The UF-TiO2 specimens for the FEGSEM
high-magnification observations were prepared using three different routes. Micrographs were recorded with an inside-the-lens
secondary electron detector and additionally using scanning transmission (STEM) mode. We found that coating the specimens
with carbon or Au-Pd did not improve the image quality. Carbon coating was inappropriate because it filled the gaps between
the nanoparticles giving the appearance of a smoothed, agglomerated surface. The TiO2 nanoparticles sputtered with Au-Pd
were observed to have better contrast; however, they were falsely enlarged due to the Au-Pd nano-layer around them. The best
results were obtained with the uncoated specimens, where TiO2 nanoparticles sized 5–10 nm were clearly visible at a
magnification of 300 000-times. With STEM mode we were able to distinguish the TiO2 nanoparticles relatively easily; however
with different image contrast than when using standard SEM mode.
Keywords: titanium dioxide, nanoparticles, FEGSEM
Prah titanovega dioksida z nanodelci (UF-TiO2) ima vsestranske aplikacije, ki izhajajo iz majhne velikosti delcev in njegovih
polprevodni{kih elektri~nih lastnosti. Nadgrajena je bila metoda za sintezo UF-TiO2 z mo`nostjo kontrole lastnosti nanodelcev:
njihove velikosti, oblike in kristalne strukture. Za dolo~anje vpliva razli~nih parametrov sinteze na kon~no obliko nanodelcev
TiO2 je bila preiskovana njihova velikost in oblika z visokolo~ljivostnim vrsti~nim elektronskim mikroskopom s katodo na
poljsko emisijo (FEGSEM). Vzorci UF-TiO2 so bili pripravljeni za analize FEGSEM na tri razli~ne na~ine. Posnetki so bili
narejeni z detektorjem sekundarnih elektronov, name{~enim v objektni le~i mikroskopa in dodatno v presevnem na~inu
(STEM). Ugotovljeno je, da prevodne plasti ogljika ali Au-Pd, ki so bile nanesene na povr{ino vzorcev, ne izbolj{ajo kakovosti
slike. Nanos amorfne plasti ogljika je bil neprimeren, ker je zapolnil prostor med nanodelci, dajajo~ vtis zglajene, aglomerirane
povr{ine. Nanodelci v vzorcu, ki je bil napr{en z Au-Pd, so bili vidni z bolj{im kontrastom, vendar so bili navidezno ve~ji, kot
so v resnici, zaradi dodane plasti Au-Pd okoli njih. Najbolj{i rezultati so bili dobljeni z opazovanjem nenapr{enega vzorca, kjer
so bili osnovni nanodelci TiO2 z velikostjo od 5 nm do 10 nm jasno vidni pri 300 000-kratni pove~avi. S STEM je relativno
enostavno razlo~iti nanodelce TiO2, vendar z druga~nim slikovnim kontrastom v primerjavi s standardnim SEM-na~inom.
Klju~ne besede: titanov dioksid, nanodelci, FEGSEM
1 INTRODUCTION
Ultrafine, nanometer-grade titanium dioxide (UF-
TiO2) is known for its versatile applications, originating
from its very high specific surface due to the nano-sized
particles and its semiconducting electric properties. The
range of use for TiO2 nanoparticles in its two crystalline
forms, anatase or rutile, is very wide: photocatalytic
applications, UV absorbing transparent coatings, plastics
additive, cosmetics UV blockers, photo-electrochromic
windows, DSSCs ("dye-sensitized solar cells") and many
more.1–3 The production of TiO2 material with a high
specific surface is one of the core development activities
in the factory Cinkarna Celje, with a strategic orientation
towards the production of UF-TiO2 in water-suspension
form. For that reason the gel-sol synthesis method was
employed and upgraded to obtain anatase and/or rutile
TiO2 nanoparticles.4 To study the influence of various
reaction parameters of the gel-sol synthesis on the cha-
racteristics of the final UF-TiO2 product it is necessary to
perform fast and reliable analyses of the samples of TiO2
powders that consist of nanoparticles. Namely, it is
necessary to have reliable data on the particle size and
the morphology of the ultrafine powders after each syn-
thesis protocol in order to establish the correct correla-
tion between the processing parameters and the mate-
rial’s characteristics according to a specific application.
A modern field-emission-gun scanning electron micro-
scope (FEGSEM) with a very-high-brightness cathode is
one of the analytical tools that can be used for the
characterization of nanoparticles. The very small, fo-
cused, electron-probe diameter in the FEGSEM makes it
possible to observe the specimens with an ultimate
resolution of about 1 nm.5 In this work we implemented
Materiali in tehnologije / Materials and technology 49 (2015) 1, 173–176 173
UDK 537.533:66.017:620.3 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(1)173(2015)
advanced methods of FEGSEM microscopy to study the
ultrafine, nano-sized TiO2 powders and to determine the
morphology and the size of both the agglomerates and
the constituent TiO2 nanoparticles.
2 EXPERIMENTAL
For the FEGSEM analyses the powder specimens
with ultrafine TiO2 nanoparticles were prepared in three
different ways. Initially, the TiO2 nanoparticles were
dispersed from properly diluted suspensions using an
ultrasonic device and then deposited and dried on poli-
shed aluminum holders. The first-type specimens were
uncoated, the second-type specimens were coated with
an amorphous carbon layer 4 nm and the third-type
specimens were coated with a layer 3 nm of Au-Pd alloy.
The coatings were applied in a Gatan PECS 682
ion-beam coating apparatus with specimen rotation up to
20 r/min, tilt up to 25° and a rocking speed of 10° s–1.
The specimens were observed using two FEGSEM
microscopes: a JEOL JSM-7600F and a Zeiss Sigma VP.
The FEGSEM experimental set-ups were adjusted and
optimized according to the individual specimen and the
desired final magnification, with an emphasis on the
preferential application of low-voltage high-resolution
microscopy.6 Thus the applied SEM accelerating volt-
ages were set in the range 2–10 kV with electron-beam
currents of 20–100 pA and working distances of 3–10
mm. An inside-the-lens secondary-electron detector was
used for the imaging. All the images were recorded
without noise-reduction tools. In addition, the SEM
scanning-transmission (STEM) mode was used at 30 kV
with samples that were prepared by the direct dispersion
of the nanoparticles on a thin carbon-membrane.
3 RESULTS AND DISCUSSION
The high-resolution FEGSEM micrographs of the
uncoated specimen of the ultra-fine TiO2 anatase powder
were recorded using a sequence of magnifications, as
shown in Figure 1. The rounded agglomerates of the
particles with a size of  1 μm remained stable, even
after applied ultrasonic dispersion, as displayed in
Figure 1a. The micrographs at higher magnifications of
100 000-times and 300 000-times (Figures 1b and 1c)
clearly revealed, angularly shaped, constituent anatase
nanocrystallites within the agglomerates. From these
images the size of the crystallites was easily estimated to
be up to 20 nm. Furthermore, exploiting the limits of the
microscope by operating under specially optimized ima-
ging conditions, we were able to record the images of the
basic nanocrystallites at an ultimate, maximum magni-
fication of 1 000 000-times, as presented in Figure 1d.
Dimensional measurements were then performed direct-
ly on the image and these confirmed that the size of the
anatase nanocrystallites is in the range 10–20 nm. More-
over, the shape of some nanocrystallites was recognized
as being similar to octahedra, which is typical for the
anatase crystal structure.7
Two samples of crystalline UF-TiO2, rutile and ana-
tase, were prepared employing identical sample prepa-
ration routes and then examined using the same FEG-
SEM set-up with a final magnification of 100 000-times.
The micrograph presented in Figure 2 shows rutile
nanoparticles in the form of "wheat-like" grains with a
size of 20–40 nm, also exhibiting a narrow size distri-
bution as a consequence of the inherent monodispersed
nature of the rutile precursor within the applied synthesis
procedure.8 In contrast, the anatase sample looks diffe-
rent, revealing the presence of nano-agglomerates that
consisted of very small, basic nanoparticles with a size
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174 Materiali in tehnologije / Materials and technology 49 (2015) 1, 173–176
Figure 2: The micrograph of the basic nanoparticles of rutile UF-TiO2
Slika 2: Posnetek osnovnih nanodelcev UF-TiO2 rutila
Figure 1: FEGSEM micrographs of the ultrafine anatase TiO2: a)
submicrometer-sized agglomerates of the nanoparticles, b), c) basic
anatase nanoparticles in the agglomerates at 100 000-times and
300 000-times magnifications, d) the image of the anatase nanocry-
stallites at highest accessible 1 000 000-times magnification
Slika 1: FEGSEM-posnetki zelo finega prahu anatasa TiO2: a) sub-
mikrometrsko veliki aglomerati nanodelcev, b), c) osnovni nanodelci
anatasa v aglomeratih, posneti pri pove~avah 100 000-krat in
300 000-krat, d) posnetek nanokristalitov anatasa pri najve~ji mo`ni
pove~avi 1 000 000-krat
below 10 nm, as shown in Figure 3. The size of these
agglomerates was between 20 nm and 40 nm only.
In order to get a more detailed view of the constituent
anatase nanoparticles the specimens were additionally
treated using the three preparation routes mentioned in
the experimental part. An ultra-high-magnification
300 000-times micrograph of the uncoated anatase speci-
men is shown in Figure 4 and clearly reveals the
authentic structure of the nano-agglomerates, which
consist of the basic TiO2 nanoparticles sized 5–10 nm. In
the case of the uncoated sample, surface charging was
avoided by operating at a properly selected accelerating
voltage, with lower beam currents and using faster scan-
ning speeds for the recording. A similar image of the
specimen coated with a Au-Pd nanolayer shown in
Figure 5 provides basically the same information as
given in Figure 4; however, due to applied coating the
agglomerates seem somewhat rounded and also the
nanoparticles are falsely slightly larger. The imaging of
the Au-Pd-coated anatase specimen was easier to per-
form than for the uncoated specimen, because the sample
was more stable under the electron-beam and was
sufficiently conductive to avoid charging artifacts, thus
allowing us to operate at higher beam currents. Also, the
image contrast and signal-to-noise ratio were better
using higher currents and because of the high secon-
dary-electron yield of the Au-Pd material. If one would
neglect these artificial changes to the morphology due to
the Au-Pd coating, these prepared samples are accepta-
ble, especially because of the easier FEGSEM high-
magnification imaging. In comparison with the uncoated
and the Au-Pd-coated samples the third anatase sample
that was coated with carbon appeared very different, as
shown in Figure 6. An amorphous carbon layer covered
the surface of the agglomerates, filling the gaps between
the nanoparticles. Therefore, the agglomerates look
Z. SAMARD@IJA et al.: CHARACTERIZATION OF TiO2 NANOPARTICLES WITH HIGH-RESOLUTION FEG ...
Materiali in tehnologije / Materials and technology 49 (2015) 1, 173–176 175
Figure 6: The high-resolution micrograph of nanometer-sized anatase
agglomerates; specimen coated with amorphous carbon layer
Slika 6: Visokolo~ljivostni posnetek nanometrskih aglomeratov
anatasa; vzorec, napr{en z amorfno plastjo ogljika
Figure 4: The high-resolution micrograph of the details of nano-
meter-sized anatase agglomerates with constituent TiO2 nanoparticles
sized 5–10 nm; uncoated specimen
Slika 4: Visokolo~ljivostni posnetek detajlov nanometrskih aglome-
ratov anatasa z osnovnimi nanodelci TiO2 z velikostjo 5–10 nm;
nenapr{en vzorec
Figure 5: The high-resolution micrograph of the details of nano-
meter-sized anatase agglomerates; specimen coated with Au-Pd layer
Slika 5: Visokolo~ljivostni posnetek detajla nanometrskih aglomera-
tov anatasa; vzorec, napr{en s plastjo iz zlitine Au-Pd
Figure 3: The micrograph of nanometer-sized agglomerates of anatase
UF-TiO2
Slika 3: Posnetek aglomeratov nanodelcev UF-TiO2 anatasa
falsely smooth and in this case it was not possible to
resolve the basic TiO2 nanoparticles.
A complementary bright-field (BF) STEM image of
agglomerated anatase nanoparticles is presented in
Figure 7. Because of the high applied accelerating volt-
age (30 kV) the issues related to accurate astigmatism
correction and precise electron-beam centering are less
pronounced, although they are very important in the
standard imaging mode. The contours of the agglomera-
tes are sharp and visible, thus allowing a straightforward
estimation of their size. The drawback of the STEM
observation is rapid sample contamination and the sub-
sequent loss of image contrast. Once the area of interest
is found, the STEM images have to be recorded as
quickly as possible so as to reduce the contamination and
to keep the image contrast at the desired level. The basic
anatase nanoparticles were also visible in STEM mode.
In this case dark-field (DF) imaging was used at higher
magnifications, as shown in Figure 8. The bright spot in
the upper-left corner of the image and the visible traces
of the grey frame arise from the contamination that
occurred within several seconds of exposure under the
electron-beam.
4 CONCLUSIONS
The morphology, size and size distribution of the
ultrafine titanium dioxide powders consisting of nano-
particles were successfully analyzed using a high-resolu-
tion FEGSEM with useful images attained at magnifica-
tions up to 1 000 000-times. By comparing the images of
the differently prepared samples we found that sputtered
conductive layers of either carbon or Au-Pd do not
improve the quality of the FEGSEM micrographs. The
best micrographs were obtained on genuine, uncoated
specimens where TiO2 nanoparticles sized 5–10 nm were
clearly revealed at a 300 000-times magnification. Car-
bon coating was found to be inappropriate because it
filled the gaps between the nanoparticles, making the
agglomerates artificially smooth. The nanoparticles in
the specimen coated with Au-Pd could be observed
easily with better contrast; however, they appeared
slightly bigger because of the coating. Using the STEM
mode we were able to see agglomerated anatase nano-
particles relatively easily, albeit with a different image
contrast. The best images for the characterization of the
morphology and the size of TiO2 nanoparticles were
recorded using the inside-the-lens secondary electron
detector and the appropriate FEGSEM set-ups optimized
for ultra-high-magnification observation.
Acknowledgements
This work was supported by the Slovenian Research
Agency (ARRS) within the project J2-4237 "Electron
microscopy and microanalysis of materials on submicro-
meter scale".
5 REFERENCES
1 Z. Zhang, C. C. Wang, R. Zakaria, J. Y. Ying, J. Phys. Chem. B, 102
(1998), 10871–10878
2 A. L. Linsebigler, G. Lu, J. T. Yates Jr., Chem. Rev., 95 (1995),
735–758
3 B. O’Regan, M. Gratzel, Nature, 353 (1991), 737–740
4 T. Sugimoto, X. Zhou, A. Muramatsu, J. Colloid Interface Sci., 259
(2003), 43–52
5 J. I. Goldstein, D. E. Newbury, D. C. Joy, C. E. Lyman, P. Echlin, E.
Lifshin, L. Sawyer, J. R. Michael, Scanning Electron Microscopy
and X-Ray Microanalysis, 3rd ed., Kluwer Academic Publishers,
New York 2003, 34
6 S. Asahina, T. Togashi, O. Terasaki, S. Takami, T. Adschiri, M.
Shibata, N. Erdman, Microscopy and Analysis, 26 (2012), S12–S14
7 Y. Li, T. J. White, S. H. Lim, J. Solid State Chem., 177 (2004),
1372–1381
8 J. Park, V. Priman, E. Matijevi}, J. Phys. Chem. B, 105 (2001),
11630–11635
Z. SAMARD@IJA et al.: CHARACTERIZATION OF TiO2 NANOPARTICLES WITH HIGH-RESOLUTION FEG ...
176 Materiali in tehnologije / Materials and technology 49 (2015) 1, 173–176
Figure 8: STEM dark-field image of nanometer-sized anatase agglo-
merates and nanoparticles
Slika 8: STEM-posnetek v temnem polju nanometrskih aglomeratov
in nanodelcev anatasa
Figure 7: STEM bright-field image of nanometer-sized anatase agglo-
merates
Slika 7: STEM-posnetek v svetlem polju nanometrskih aglomeratov
anatasa
Erratum
G. Novak, J. Koko{ar, A. Nagode, D. Steiner Petrovi~,
Correlation between the excess losses and the relative
permeability in fully finished non-oriented electrical
steels, Mater. Tehnol., 48 (2014) 6, 997–1001
The Editor-in-Chief declared by mistake the article as
professional article, correct is original scientific article,
as evident from the reviewers report.
A/Prof. Dr. Matja` Torkar
Editor-in-Chief
Materiali in tehnologije / Materials and technology 49 (2015) 1, 177 177