VSEBINA – CONTENTS
PREGLEDNI ^LANKI – REVIEW ARTICLES
Kinetic study and characterization of borided AISI 4140 steel
[tudij kinetike in karakterizacija boriranega jekla AISI 4140
M. Keddam, M. Ortiz-Domínguez, O. A. Gómez-Vargas, A. Arenas-Flores, M. Á. Flores-Rentería, M. Elias-Espinosa,
A. García-Barrientos . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 665
Kinetics of the precipitation in austenite HSLA steels
Kinetika izlo~anja v avstenitnih HSLA-jeklih
E. Kalinowska-Ozgowicz, R. Kuziak, W. Ozgowicz, K. Lenik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 673
Combined influence of V and Cr on the AlSi10MgMn alloy with a high Fe level
Vzajemni vpliv V in Cr na zlitino AlSi10MgMn z visoko vsebnostjo Fe
D. Bolibruchová, M. @ihalová . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 681
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
A reliable approach to a rapid calculation of the grain size of polycrystalline thin films after excimer laser crystallization
Zanesljiv na~in hitrega izra~una velikosti zrn v polikristalni tanki plasti po UV-laserski kristalizaciji
C. C. Kuo . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 687
Effects of different stirrer-pin forms on the joining quality obtained with friction-stir welding
Vpliv razli~nih oblik vrtilnih konic na kvaliteto spoja pri tornem vrtilnem varjenju
H. Basak, K. Kaptan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 693
Monitoring early-age concrete with the acoustic-emission method and determining the change in the electrical properties
Pregled sve`ega betona z metodo akusti~ne emisije in dolo~anjem sprememb elektri~nih lastnosti
L. Pazdera, L. Topolar, M. Korenska, T. Vymazal, J. Smutny, V. Bilek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 703
Effect of the aggregate type on the properties of alkali-activated slag subjected to high temperatures
Vpliv vrste agregata na lastnosti z alkalijo aktivirane `lindre, izpostavljene visokim temperaturam
P. Rovnaník, Á. Dufka. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 709
Microstructure evolution of advanced high-strength TRIP-aided bainitic steel
Razvoj mikrostrukture naprednega visokotrdnostnega bainitnega jekla z uporabo TRIP
A. Grajcar. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 715
Effect of electric current on the production of NiTi intermetallics via electric-current-activated sintering
Vpliv elektri~nega toka pri izdelavi intermetalne zlitine NiTi s sintranjem, aktiviranim z elektri~nim tokom
T. Yener, S. Siddique, F. Walther, S. Zeytin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 721
Control of soft reduction of continuous slab casting with a thermal model
Kontrola mehke redukcije pri kontinuirnem litju slabov s termi~nim modelom
J. Stetina, P. Ramik, J. Katolicky . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 725
Optimization of operating conditions in a laboratory SOFC testing device
Optimizacija obratovalnih razmer laboratorijske gorivne celice SOFC
T. Skalar, M. Lubej, M. Marin{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 731
Production of shaped semi-products from AHS steels by internal pressure
Izdelava polproizvodov iz AHS-jekel, oblikovanih z notranjim tlakom
I. Vorel, H. Jirková, B. Ma{ek, P. Kurka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 739
Composite-material printed antenna for a multi-standard wireless application
Tiskana antena iz kompozitnega materiala za ve~standardno brez`i~no uporabo
T. Alam, M. R. I. Faruque, M. T. Islam, N. Misran . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 745
Friction and wear behaviour of ulexite and cashew in automotive brake pads
Odpornost proti trenju in obrabi avtomobilskih zavornih oblog z uleksitom in prahom iz indijskega oreha
Ý. Sugözü, Ý. Mutlu, A. Keskin. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 751
Diffusion kinetics and characterization of borided AISI H10 steel
Kinetika difuzije in karakterizacija boriranega jekla AISI H10
I. Gunes, M. Ozcatal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 759
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 49(5)663–844(2015)
MATER.
TEHNOL.
LETNIK
VOLUME 49
[TEV.
NO. 5
STR.
P. 663–844
LJUBLJANA
SLOVENIJA
SEP.–OCT.
2015
Application of the Taguchi method to optimize the cutting conditions in hard turning of a ring bore
Uporaba Taguchi-jeve metode za optimizacijo trdega stru`enja roba izvrtine
M. Boy, I. Ciftci, M. Gunay, F. Ozhan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 765
Thermal fatigue of single-crystal superalloys: experiments, crack-initiation and crack-propagation criteria
Toplotno utrujanje monokristalnih superzlitin: preizkusi, merila za nastanek in napredovanje razpoke
L. Getsov, A. Semenov, S. Semenov, A. Rybnikov, E. Tikhomirova . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 773
Investigating the effects of cutting parameters on the built-up-layer and built-up-edge formation during the machining of AISI
310 austenitic stainless steels
Preiskava vplivov parametrov rezanja na nastanek nakopi~ene plasti in nakopi~enega roba med stru`enjem avstenitnega nerjavnega
jekla AISI 310
M. B. Bilgin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 779
Mortar-type identification for the purpose of reconstructing fragmented Roman wall paintings (Celje, Slovenia)
Analiza ometov za rekonstrukcijo fragmentov rimskih stenskih poslikav (Celje, Slovenija)
M. Gutman, M. Lesar Kikelj, J. Kuret, S. Kramar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 785
Au-nanoparticle synthesis via ultrasonic spray pyrolysis with a separate evaporation zone
Sinteza nanodelcev zlata z ultrazvo~no razpr{ilno pirolizo z lo~eno cono izhlapevanja
P. Majeri~, B. Friedrich, R. Rudolf . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 791
Determination of the critical fraction of solid during the solidification of a PM-cast aluminium alloy
Dolo~anje kriti~nega dele`a strjene faze med strjevanjem v PM ulite aluminijeve zlitine
R. Kayikci, M. Colak, S. Sirin, E. Kocaman, N. Akar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 797
Microstructural evolution of Inconel 625 during hot rolling
Mikrostrukturni razvoj Inconela 625 med vro~im valjanjem
F. Tehovnik, J. Burja, B. Podgornik, M. Godec, F. Vode. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 801
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Thermophysical properties and microstructure of magnesium alloys of the Mg-Al type
Termofizikalne lastnosti in mikrostruktura magnezijevih zlitin tipa Mg-Al
P. Lichý, J. Beòo, I. Kroupová, I. Vasková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 807
Micro-encapsulated phase-change materials for latent-heat storage: thermal characteristics
Mikroenkapsulirani materiali s fazno premeno za shranjevanje latentne toplote: toplotne zna~ilnosti
M. Ostrý, D. Dostálová, T. Klubal, R. Pøikryl, P. Charvát. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 813
Ceramic masonry units intended for the masonry resistant to high humidity
Kerami~ni gradbeni elementi, namenjeni za zgradbe, odporne proti visoki vlagi
J. Zach, V. Novák, J. Hroudová, M. Sedlmajer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 817
Flame resistance and mechanical properties of composites based on new advanced resin system FR4/12
Negorljivost in mehanske lastnosti kompozitov na osnovi novih naprednih sistemov smol FR4/12
V. Rusnák, S. Rusnáková, L. Fojtl, M. @aludek, A. ^apka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 821
Effect of process parameters on the microstructure and mechanical properties of friction-welded joints of AISI 1040/AISI
304L steels
Vpliv procesnih parametrov na mikrostrukturo in mehanske lastnosti torno varjenih spojev jekel AISI 1040/AISI 304L
Ý. Kirik, N. Özdemýr . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 825
Influence of inoculation methods and the amount of an added inoculant on the mechanical properties of ductile iron
Vpliv metod modifikacije in koli~ine dodanega modifikatorja na mehanske lastnosti duktilnega `eleza
H. Avdu{inovi}, A. Gigovi}-Geki}, D. ]ubela, R. Sunulahpa{i}, N. Mujezinovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 833
One-step green synthesis of graphene/ZnO nanocomposites for non-enzymatic hydrogen peroxide sensing
Enostopenjska zelena sinteza nanokompozita grafen-ZnO za neencimatsko detekcijo vodikovega peroksida
S. S. Low, M. T. T. Tan, P. S. Khiew, H. S. Loh, W. S. Chiu . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 837
Material parameters for a numerical simulation of a compaction process for sintered double-height gears
Materialni parametri za numeri~no simulacijo postopka stiskanja sintranih dvovi{inskih zobnikov
T. Verlak, M. [ori, S. Glode` . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 841
M. KEDDAM et al.: KINETIC STUDY AND CHARACTERIZATION OF BORIDED AISI 4140 STEEL
665–672
KINETIC STUDY AND CHARACTERIZATION OF BORIDED AISI
4140 STEEL
[TUDIJ KINETIKE IN KARAKTERIZACIJA BORIRANEGA JEKLA
AISI 4140
Mourad Keddam1, Martín Ortiz-Domínguez2, Oscar Armando Gómez-Vargas3,
Alberto Arenas-Flores4, Miguel Ángel Flores-Rentería2, Milton Elias-Espinosa5,
Abel García-Barrientos6
1Département de Sciences des Matériaux, Faculté de Génie Mécanique et Génie des Procédés, USTHB, B.P. No. 32, 16111 El-Alia,
Bab-Ezzouar, Algiers, Algeria
2Universidad Autónoma del Estado de Hidalgo, Escuela Superior de Ciudad Sahagún-Ingeniería Mecánica, Carretera Cd. Sahagún-Otumba s/n,
Zona Industrial CP. 43990 Hidalgo, México
3Instituto Tecnológico de Tlalnepantla-ITTLA. Av., Instituto Tecnológico, S/N. Col. La Comunidad, Tlalnepantla de Baz. CP. 54070 Estado de
México, México
4Universidad Autónoma del Estado de Hidalgo-AACTyM, Carretera Pachuca Tulancingo Km. 4.5, Mineral de la Reforma, CP. 42184 Hidalgo,
México
5Instituto Tecnológico y de Estudios Superiores de Monterrey-ITESM Campus Santa Fe, Av. Carlos Lazo No. 100, Del. Álvaro Obregón, CP.
01389 D. F., México
6Universidad Autónoma del Estado de Hidalgo-CITIS, Carretera Pachuca Tulancingo Km. 4.5, Mineral de la Reforma, CP. 42184 Hidalgo,
México
keddam@yahoo.fr
Prejem rokopisa – received: 2014-02-11; sprejem za objavo – accepted for publication: 2014-10-08
doi:10.17222/mit.2014.034
In the present study, an alternative diffusion model was proposed for analyzing the growth of Fe2B layers formed on the AISI
4140 steel during the pack-boriding process. This model was based on solving the mass-balance equations for the Fe2B/Fe
interface to evaluate boron diffusion coefficients through the Fe2B layers in a temperature range of 1123–1273 K. The boride
incubation time for the Fe2B phase was included in the present model. The suggested model was validated experimentally at a
temperature of 1253 K for a treatment time of 5 h.
Furthermore, the generated boride layers were analyzed with light microscopy, scanning electron microscopy (SEM), energy
dispersive X-ray spectroscopy (EDS) and X-ray diffraction analysis (XRD). In addition, a contour diagram was also proposed as
a function of treatment time and temperature. On the basis of our experimental results, the boron activation energy for the AISI
4140 steel was found to be 189.24 kJ mol–1.
Keywords: boriding, incubation time, diffusion model, growth kinetics, activation energy, adherence
V tej {tudiji je predlagan alternativni model difuzije za analiziranje rasti Fe2B-plasti, ki nastane na jeklu AISI 4140 pri boriranju
v {katli. Ta model temelji na re{itvi ena~be ravnote`ja mas na stiku (Fe2B/Fe) pri oceni koeficienta difuzije bora skozi Fe2B-sloje
v temperaturnem obmo~ju 1123–1273 K. Inkubacijski ~as borida za Fe2B-fazo je bil vklju~en v predstavljeni model. Predlagani
model je bil eksperimentalno ocenjen pri temperaturi 1253 K in ~asu obdelave 5 h.
Poleg tega so bili izdelani sloji analizirani na svetlobnem mikroskopu, z vrsti~nim elektronskim mikroskopom (SEM),
energijsko disperzijsko rentgensko spektroskopijo (EDS) in z rentgensko difrakcijsko analizo (XRD). Dodatno je bil predlagan
{e konturni diagram kot funkcija ~asa obdelave in temperature. Na podlagi eksperimentalnih rezultatov je bilo ugotovljeno, da je
aktivacijska energija bora pri jeklu AISI 4140 189,24 kJ mol–1.
Klju~ne besede: boriranje, ~as inkubacije, model difuzije, kinetika rasti, aktivacijska energija, adherenca
1 INTRODUCTION
Boriding is a well-known thermochemical treatment
in which boron (because of its relatively small size)
diffuses into the metal substrate to form hard borides. As
a result of boriding, properties such as wear resistance,
surface hardness and corrosion resistance are improved.1
Boriding can be carried out with boron in different states
such as solid powder, paste, liquid, gas and plasma. The
most frequently used method is pack boriding owing to
its technical advantages.2 Generally, the commercial
boriding mixture is composed of boron carbide (B4C) as
a donor, KBF4 as an activator and silicon carbide (SiC) as
a diluent to control the boriding potential of the medium.
The boriding treatment requires temperatures ranging
from 800 °C to 1000 °C. Usually the treatment time
varies between 0.5 h and 12 h producing a boride layer
whose thickness depends on the boriding parameters
(time and temperature). The morphology of the boride
layer is affected by the presence of alloying elements in
the matrix. Saw-tooth-shaped layers are obtained in
low-alloy steels or Armco iron whereas in high-alloy
steels, the interfaces are smooth. According to the Fe-B
phase diagram,3 two iron borides can be formed (FeB
and Fe2B).
A monophase Fe2B layer with a tooth-shaped mor-
phology is generally suitable for industrial applications
because of the difference between the specific volume
and the coefficients of the thermal expansion of iron
Materiali in tehnologije / Materials and technology 49 (2015) 5, 665–672 665
UDK 669.058:669.781:544.034 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 49(5)665(2015)
boride and the substrate.4,5 The boron-rich phase FeB is
not preferred since FeB is more brittle and less tough
than Fe2B.4,5 Furthermore, the brittleness of FeB layers
causes a spalling when a high normal or tangential load
is applied.
The modeling of the boriding kinetics is considered
as a suitable tool to select the optimized parameters for
obtaining a desired boride layer of the treated material
for its practical use in the industry. In particular, many
models reported in the literature were used for analyzing
the kinetics of the Fe2B layers grown on different sub-
strates6–16, with and without the boride incubation times.
In the current work, an alternative diffusion model,
based on solving the mass-balance equation for the
Fe2B/substrate interface, was proposed to simulate the
kinetics of the Fe2B layers grown on the AISI 4140 steel.
In the present model, the boride incubation time was
independent of the temperature. The pack-borided AISI
4140 steel was characterized by means of the following
techniques: light microscopy, scanning electron micro-
scopy and XRD. Based on the experimental data, the
boron activation energy was also evaluated when pack
boriding the AISI 4140 steel in a temperature range of
1123–1273 K.
2 KINETIC MODEL
The model considers the growth of the Fe2B layer on
a substrate saturated with boron atoms as illustrated in
Figure 1. The f(x,t) function represents the boron
distribution in the iron matrix before the nucleation of
the Fe2B phase. t 0
Fe B2 corresponds to the incubation time
required to form the Fe2B phase when the matrix reaches
the state of being saturated with boron atoms. C up
Fe B2
represents the upper limit of the boron content in Fe2B
(60 · 103 mol m–3), C low
Fe B2 is the lower limit of the boron
content in Fe2B (59.8 · 103 mol m–3) and the point x(t = t)
= v represents the Fe2B layer thickness.17,18
The term C ads
B represents the effective adsorbed boron
concentration during the boriding process.19 In Figure 1,
a C C1 = −up
Fe B
low
Fe B
2 2 defines the homogeneity range of the
Fe2B layer, a C C2 = −low
Fe B
0
2 is the miscibility gap15,16 and
C0 is the boron solubility in the matrix considered as
zero.3 The following assumptions are considered for the
diffusion model:
– The growth kinetics is controlled by the boron diffu-
sion in the Fe2B layer.
– The Fe2B iron boride nucleates after a specific incu-
bation time.
– The boride layer grows because of the boron diffu-
sion perpendicular to the specimen surface.
– Boron concentrations remain constant in the boride
layer during the treatment.
– The influence of the alloying elements on the growth
kinetics of the layer is not taken into account.
– The boride layer is thin compared to the sample
thickness.
– A uniform temperature is assumed throughout the
sample.
– A planar morphology is assumed for the phase inter-
face.
The initial and boundary conditions for the diffusion
problem are represented as:
t = 0, x > 0, with: CFe2B [x(t), t = 0] = C0  0 (1)
Boundary conditions:
[ ]C x t t v t t CFe B Fe B Fe B upFe B2 2 2 2( ) ,= = = =0 0 0 (2)
(the upper boron concentration is kept constant)
for C ads
B > 60 · 103 mol m–3:
[ ]C x t t v t t CFe B low
Fe B
2
2( ) ,= = = = (3)
(the upper boron concentration is kept constant)
for C ads
B < 59.8 · 103 mol m–3
v0 is a thin layer with a thickness of  5 nm that formed
during the nucleation stage,20 thus v0 ( 0) is used when
compared to the thickness of the Fe2B layer (v). The
mass-balance equation at the Fe2B/substrate interface
can be formulated with Equation (4) as follows:
C C C
A v
J x v t t
up
Fe B
low
Fe B
Fe B
2 2
2
2
d
+ −⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟ ⋅ =
= = =
2 0
( )
( , )( ) ( , )( )A t J x v v t t A t⋅ − = + = ⋅d d dFe
(4)
where A(=1.1) is defined as the unit area and C0 repre-
sents the boron concentration in the matrix. The fluxes
JFe2B and JFe are obtained from the Fick’s first law as:
M. KEDDAM et al.: KINETIC STUDY AND CHARACTERIZATION OF BORIDED AISI 4140 STEEL
666 Materiali in tehnologije / Materials and technology 49 (2015) 5, 665–672
Figure 1: Schematic boron-concentration profile through the Fe2B
layer
Slika 1: Shematski prikaz profila koncentracije bora skozi Fe2B-plast
[ ]
[ ]
J x t t v t t
D C x t t v t t
x
Fe B
Fe B Fe B
2
2 2
( ) , )
( ) , )
= = = =
= −
= = =⎧
⎨
⎩
∂
∂
⎫
⎬
⎭ =x v
(5)
and:
[ ]
[ ]
J x t t v v t t
D C x t t v v t t
x
Fe
Fe Fe
d
d
( ) , )
( ) , )
= = + = =
= −
= = + =⎧
⎨
⎩
∂
∂
⎫
⎬
⎭ = +x v vd
(6)
The term JFe is zero since the boron solubility in the
matrix is very low ( 0 mol m–3).3
Thus, Equation (4) can be written as:
C C C x t
t
D
x t v
up
Fe B
low
Fe B
Fe B
2 2
2
2
d
d
+ −⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟ =
= −
=
2 0 ( )
( )
[ ]∂
∂
C x t t t t
x
x t v
Fe B2
( ),
( )
= =
=
(7)
The Fick’s second law of diffusion relating to the
change in the boron concentration through the Fe2B layer
with the time t and the distance x(t) is given with
Equation (8) :
[ ] [ ]∂
∂
∂
∂
C x t t
t
D
C x t t
x
Fe B
Fe B
Fe B2
2
2
( ), ( ),
=
2
2 (8)
When applying the boundary conditions proposed in
Equations (2) and (3), the solution of Equation (8) takes
the following form:
[ ]C x t t C
C C
v
D t
Fe B up
Fe B low
Fe B
up
Fe B
Fe B
2
2
2 2
2
erf
2
( ), = +
−
⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟
⋅
⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟erf 2 Fe B2
x
D t
(9)
By substituting the derivative of Equation (9) with
respect to the distance x(t) in Equation (7), Equation (10)
is obtained:
C C C v
t
D
t
C
up
Fe B
low
Fe B
Fe B up
Fe B
2 2
2
2
2
d
d
+ −⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟ =
= ⋅
2 0
π
( )−
⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟
⋅ −
⎛
⎝
⎜
⎞
⎠
⎟
C
v
D t
v
D t
low
Fe B
Fe B
Fe B
2
2
2
erf
2
exp
4
2 (10)
for 0  x  v.
Substituting the expression of the parabolic growth
law (v D t= 2 Fe B2 ) in Equation (10), Equation (11) is
deduced:
C C C
C C
up
Fe B
low
Fe B
up
Fe B
low
Fe B
2 2
2 2
2
+ −⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟ =
=
−
2
1
0

π erf( )

⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟ −exp( )
2
(11)
The normalized growth parameter () for the
Fe2B/substrate interface can be estimated numerically
using the Newton-Raphson method. It is assumed that
expressions C up
Fe B2 , C low
Fe B2 and C0 do not depend signifi-
cantly on the temperature (in the considered temperature
range).18
A schematic representation of the square of the layer
thickness against the linear time (v2 = 42DFe2Bt =
42DFe2B[tv + t 0
Fe B2 (T)]) is depicted in Figure 2. tv(= t –
t 0
Fe B2 ) is the effective growth time of the Fe2B layer and t
is the treatment time.
3 EXPERIMENTAL PROCEDURE
3.1 Boriding process
The AISI 4140 steel was used in this experimental
study. It had a nominal chemical composition of
0.38–0.43 % C, 0.75–1.00 % Mn, 0.80–1.10 % Cr,
0.15–0.30 % Si, 0.040 % S, and 0.035 % P. Samples had
a cubic shape with dimensions of 10 mm × 10 mm × 10
mm. Prior to the boriding process, the specimens were
polished, ultrasonically cleaned in an alcohol solution
and deionized water for 15 min at room temperature,
dried and stored under clean-room conditions. The sam-
ples were packed along with a Durborid fresh powder
mixture in a closed cylindrical case (AISI 304L). The
used powder mixture had an average size of 30 μm. The
powder-pack boriding process was carried out in a
conventional furnace under a pure argon atmosphere in
the temperature range of 1123–1273 K. Four treatment
times ((2, 4, 6 and 8) h) were selected for each tempe-
rature. After the completion of the boriding treatment,
the container was removed from the furnace and slowly
cooled to room temperature.
3.2 Microscopical observations of boride layers
The borided samples were cross-sectioned for me-
tallographic examinations using a LECO VC-50 preci-
sion cutting machine. The cross-sectional morphology of
the boride layers was observed with an Olympus GX51
light microscope in a clear field. Figure 3 shows
cross-sectional views of light images of the Fe2B layers
of the AISI 4140 steel formed at a temperature of 1223
K over different process durations.
The resultant microstructures of the Fe2B layers
appear to be very dense and homogenous, exhibiting
saw-tooth morphologies. Since the growth of a saw-tooth
boride layer is a controlled diffusion process with a
highly anisotropic nature, higher temperatures and/or
M. KEDDAM et al.: KINETIC STUDY AND CHARACTERIZATION OF BORIDED AISI 4140 STEEL
Materiali in tehnologije / Materials and technology 49 (2015) 5, 665–672 667
longer times stimulated the Fe2B crystals to make contact
with the adjacent crystals and forced them to retain an
acicular shape.21 It is seen that the thickness of the Fe2B
layer increased with an increase in the boriding time
(Figure 3) because the boriding kinetics is influenced by
the treatment time. For the kinetic study, the boride-layer
thickness was automatically measured with the aid of the
MSQ PLUS software. To ensure the reproducibility of
the measured layer thicknesses, fifty measurements were
collected in different sections of the samples of the
borided AISI 4140 steel to estimate the Fe2B layer thick-
ness, defined as the average value of the long boride
teeth.22–24 All the thickness measurements were taken
from a fixed reference on the surface of the borided AISI
4140 steel, as illustrated in Figure 4.
The phases of the boride layers were investigated
with X-ray diffraction (XRD) equipment (Equinox 2000)
using Co-K radiation with a wavelength of 0.179 nm.
The elemental distribution within the cross-section of a
boride layer was determined with electron dispersive
spectroscopy (EDS) equipment (JEOL JSM 6300 LV),
from the surface.
4 RESULTS AND DISCUSSIONS
4.1 SEM observations and EDS analyses
A cross-sectional view of the SEM micrograph of the
AISI 4140 steel borided at 1273 K for 6 h is shown in
Figure 5a. The boride layer grown on the substrate has a
saw-tooth morphology. The needles of Fe2B, with
different lengths, are visible on the SEM micrograph,
penetrating into the substrate. This typical morphology is
responsible for a good adhesion to the substrate. EDS
results obtained with SEM are shown in Figures 5b and
5c. These results indicate the presence of two elements at
the surface: Fe and Cr (Figure 5a). The results also show
the possible dissolution of chromium in Fe2B. In fact, the
atomic radius of Cr (= 0.166 nm) is about the same and
larger than that of Fe (= 0.155 nm), and it can then be
expected that Cr dissolved in the Fe sublattice of Fe2B.
The obtained result in Figure 5c indicates the presence
of the following elements: Fe, C, Cr, Si and Mn in the
vicinity of the Fe2B/substrate interface. It is seen that two
elements (carbon and silicon) are not dissolved in Fe2B,
being displaced towards the substrate. Silicon and boron
may form complex phases such as FeSi0.4B0.6, FeSiB and
boron cementite (Fe3B0.67C0.33).25
4.2 X-ray diffraction analysis
Figure 6 displays the XRD pattern recorded on the
surface of the AISI 4140 steel borided at a temperature
of 1273 K for a treatment time of 8 h. The diffraction
peaks relative to the Fe2B phase are easily identified.
Peaks with very small intensities are also observed for
M. KEDDAM et al.: KINETIC STUDY AND CHARACTERIZATION OF BORIDED AISI 4140 STEEL
668 Materiali in tehnologije / Materials and technology 49 (2015) 5, 665–672
Figure 2: Schematic representation of the square of the layer
thickness against the treatment time
Slika 2: Shematski prikaz kvadrata debeline plasti, odvisno od ~asa
obdelave
Figure 4: Schematic diagram illustrating the procedure for estimating
the boride-layer thickness on the AISI 4140 steel
Slika 4: Shematski diagram, ki prikazuje postopek za ugotavljanje
debeline borirane plasti na jeklu AISI 4140
Figure 3: Light micrographs of the boride layers formed on the
surface of AISI 4140 steel treated at 1223 K for variable times: a) 2 h,
b) 4 h, c) 6 h and d) 8 h
Slika 3: Svetlobni posnetki boriranih plasti, nastalih na povr{ini jekla
AISI 4140 pri temperaturi 1223 K in razli~nem trajanju: a) 2 h, b) 4 h,
c) 6 h in d) 8 h
the FeB phase but the Fe2B phase is dominant with
high-peak intensities. The CrB phase is also visible in
the XRD pattern recorded at the surface of the borided
AISI 4140 steel. In the experimental study about the
boriding of the AISI 4140 steel in molten borax, boric
acid and a ferro-silicon bath performed by Sen et al.26,
the XRD study showed the presence of FeB, Fe2B and
CrB at the surface of a sample borided at 1223 K for 6 h.
In addition, Ulutan et al.27 also identified, with an XRD
analysis, the same phases at the surface of the AISI 4140
steel (at 1000 °C for 6 h) after the powder-pack boriding.
Crystals of the Fe2B type orientate themselves with
the z-axis perpendicular to the surface. Consequently, the
peaks of the Fe2B phase belonging to the crystallo-
graphic planes and having a deviation from zero of the l
index, showed increased intensities in the X-ray diffrac-
tion spectra.28
The growth of boride layers is a controlled diffusion
process with a highly anisotropic nature. In the case of
the Fe2B phase, the [001] crystallographic direction is
the easiest path for the boron diffusion in Fe2B because
of the tendency of the boride crystals to grow along the
direction of minimum resistance, perpendicular to the
external surface. As the metal surface is covered, an
increasing number of the Fe2B crystals come in contact
with the adjacent crystals, being forced to grow inside
the metal and retaining the acicular shape.21 In the pow-
der-pack boriding, the active boron is supplied by the
powder mixture. To form a Fe2B phase on any borided
steel, a low boron potential is required as reported in the
reference works4,29, while a high amount of active boron
in a powder mixture gives rise to a bilayer configuration
consisting of FeB and Fe2B.
4.3 Estimation of the boron activation energy
The growth kinetics of the Fe2B layers formed on the
AISI 4140 steel was used to estimate the boron diffusion
coefficient through the Fe2B layers by applying the
suggested diffusion model. The  parameter is then
determined by solving the mass-balance equation for the
Fe2B/substrate interface (Equation (11)) using the
Newton-Raphson method. Table 1 lists the estimated
value of boron diffusion coefficient in Fe2B at each
temperature along with the squared value of normalized
growth parameter  determined from Equation (11).
Figure 7 depicts the time dependence of the squared
value of the Fe2B layer thickness. The slopes of the
straight lines in this figure provide the values of the
growth constants (= 4 2 DFe B2 ) for each boriding tempe-
rature. The values of the boron diffusion coefficients in
the Fe2B layers can be determined by knowing the value
of the  parameter. The boride incubation times for Fe2B
can also be deduced from the straight lines displayed in
Figure 7 by extrapolating them to the boride-layer thick-
ness of zero.
To estimate the boron activation energy for the AISI
4140 steel, it is necessary to plot the natural logarithm of
M. KEDDAM et al.: KINETIC STUDY AND CHARACTERIZATION OF BORIDED AISI 4140 STEEL
Materiali in tehnologije / Materials and technology 49 (2015) 5, 665–672 669
Figure 6: XRD pattern obtained at the surface of the borided AISI
4140 steel treated at 1273 K for 8 h
Slika 6: Rentgenogram, dobljen na povr{ini boriranega jekla AISI
4140 po 8 h na 1273 K
Figure 5: a) SEM micrograph of the cross-section of the AISI 4140
steel borided at 1273 K for 6 h, b) EDS spectrum of the surface of a
borided sample and c) EDS spectrum of the interface of a borided
sample
Slika 5: a) SEM-posnetek prereza boriranega jekla AISI 4140 po 6 h
na 1273 K, b) EDS-spekter boriranega vzorca na povr{ini in c)
EDS-spekter boriranega vzorca na stiku z osnovo
the boron diffusion coefficient for Fe2B versus the reci-
procal temperature following the Arrhenius equation
(Figure 8). A linear fitting was assumed to obtain the
temperature dependence of the boron diffusion coeffi-
cient for Fe2B with a correlation factor of 0.9935:
D
RTFe B2
4 kJ mol
= ×
−⎛
⎝
⎜
⎞
⎠
⎟−
−
151 10
189 22
1
. exp
.
(12)
where R = 8.3144621 J mol–1 K–1 and T is the absolute
temperature in Kelvin.
Table 2 shows a comparison of the boron activation
energies for some borided steels.26,30–32 The found value
of the boron activation energy (= 189.24 kJ mol–1) for the
AISI 4140 steel is slightly different from the value
reported in26 due to the boriding conditions (using the
liquid-boriding method).
4.4 Validation of the diffusion model
The present model was validated by comparing the
experimental value of the Fe2B layer thickness with the
predicted result at a temperature of 1253 K for a treat-
ment time of 5 h using Equation (13):
v
C C
C C
=
−
+
⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟ ⋅
−4
π
up
Fe B
low
Fe B
up
Fe B
low
Fe B
2 2
2 2
exp( 

2 )
( )erf
D tFe B2 (13)
Figure 9 shows the optical image of the boride layer
formed at 1253 K after 5 h of treatment.
M. KEDDAM et al.: KINETIC STUDY AND CHARACTERIZATION OF BORIDED AISI 4140 STEEL
670 Materiali in tehnologije / Materials and technology 49 (2015) 5, 665–672
Figure 8: Temperature dependence of the boron diffusion coefficient
for Fe2B
Slika 8: Temperaturna odvisnost koeficienta difuzije v Fe2B
Table 2: Comparison of the boron activation energies for some
borided steels
Tabela 2: Primerjava aktivacijskih energij bora pri nekaterih boriranih
jeklih
Material Boron activationenergy (kJ mol–1) References
AISI 5140
AISI 4340
AISI 1040
AISI 51100
AISI 4140
AISI 4140
223
324
118.8
106.0
215
189.24
30
30
31
32
26
Present study
Figure 7: Evolution of the squared value of Fe2B layer thickness as a
function of boriding time
Slika 7: Razvoj vrednosti kvadrata debeline plasti Fe2B v odvisnosti
od ~asa boriranja
Table 1: Squared value of normalized growth parameter and boron
diffusion coefficients for Fe2B as a function of boriding temperature
Tabela 1: Kvadratna vrednost normaliziranih parametrov rasti in
koeficientov difuzije bora v Fe2B v odvisnosti od temperature bori-
ranja
Temperature
(K) Type of layer
2
(Dimen-
sionless)
42DFe2B
(μm2 s–1)
1123
Fe2B 1.7×10–3
1.51×10–1
1173 4.11×10–1
1223 9.12×10–1
1273 16.3×10–1
Figure 9: Light micrograph of the cross-section of the borided AISI
4140 treated at 1253 K for 5 h
Slika 9: Svetlobni posnetek prereza boriranega jekla AISI 4140 po 5 h
na 1253 K
Table 3 gives a comparison between the experimental
value of the Fe2B layer thickness and the one predicted
on the basis of Equation (13). A good agreement was
obtained between the experimental value of the Fe2B
layer thickness and the predicted one for the AISI 4140
steel borided at 1253 K for 5 h.
Table 3: Predicted and estimated values of the Fe2B layer thickness
obtained at 1253 K for a treatment time of 5 h
Tabela 3: Predvidene in dobljene vrednosti za debelino plasti Fe2B po
5 h na 1253 K
Temperature
(K)
Type of
layer
Boride-layer
thickness (μm)
estimated by
Eq. (13)
Experimental
boride-layer
thickness (μm)
1253 Fe2B 154.48 158.12±10.43
4.5 Future exploitation of the simulation results
This kinetic approach can be used as a tool to deter-
mine the Fe2B layer thickness as a function of boriding
parameters (time and temperature) for the AISI 4140
steel. Equation (13) predicts the Fe2B layer thickness for
any temperature and boriding time. An iso-thickness
diagram was plotted as a function of the temperature and
exposure time as shown in Figure 10.
The results of Figure 10 can serve as a powerful tool
to select the optimum value of the Fe2B layer thickness
in relation with the potential applications of the borided
AISI 4140 steel at industrial scale.
As a rule, thin layers (e.g., 15–20 μm) are used to
protect against adhesive wear (in the cases of chipless-
shaping and metal-stamping dies and tools), whereas
thick layers are recommended for combating abrasive
wear (extrusion tooling for plastics with abrasive fillers
and pressing tools for the ceramic industry). In the case
of low-carbon steels and low-alloy steels, the optimum
boride-layer thicknesses range from 50 μm to 250 μm.
Finally, this model can be extended to predict the growth
kinetics of a bilayer configuration (FeB + Fe2B) grown
on any boride steel.
5 CONCLUSIONS
The AISI 4140 steel was pack borided in the tempe-
rature range of 1123–1273 K over the treatment times
varying from 2 h to 8 h. The Fe2B layers were formed on
the AISI 4140 steel substrate. A mathematical model was
suggested to estimate the boron diffusion coefficients for
the Fe2B layers. The boron activation energy for the AISI
4140 steel was found to be 189.24 kJ mol–1. This value
was compared with the data reported in the literature.
The validity of the diffusion model was examined by
comparing the experimental value of the Fe2B layer
thickness obtained at 1253 K after 5 h of treatment with
that predicted by the model.
Finally, an iso-thickness diagram was proposed to be
used as a tool to select the optimum boride layer thick-
ness in relation with the industrial use of this steel grade.
NOMENCLATURE
v is the boride-layer thickness (m)
tv is the effective growth time of the Fe2B layer (s)
t is the treatment time (s)
t 0
Fe B2 is the boride incubation time (s)
C up
Fe B2 represents the upper limit of the boron content in
Fe2B (60 · 103 mol m–3)
C low
Fe B2 is the lower limit of the boron content in Fe2B
(59.8 · 103 mol m–3)
C ads
B is the adsorbed boron concentration in the boride
layer (mol m–3)
a C C1 = −up
Fe B
low
Fe B
2 2 defines the homogeneity range of the
Fe2B layer (mol m–3)
a C C2 = −low
Fe B
0
2 is the miscibility gap (mol m–3)
C0 is the terminal solubility of the interstitial solute
( 0 mol m–3)
CFe2B [x(t)] is the boron concentration profile in the Fe2B
layer (mol m–3)
v0 indicates the initial Fe2B layer (m)
 is the normalized growth parameter for the Fe2B/sub-
strate interface (it has no physical dimensions)
DFe2B denotes the diffusion coefficient of boron in the
Fe2B phase (m2 s–1)
Ji [x(t)], (with i = Fe2B and Fe) are the fluxes of boron
atoms at the Fe2B/substrate interface boundary (mol
m–2 s–1)
M. KEDDAM et al.: KINETIC STUDY AND CHARACTERIZATION OF BORIDED AISI 4140 STEEL
Materiali in tehnologije / Materials and technology 49 (2015) 5, 665–672 671
Figure 10: Iso-thickness diagram describing the evolution of Fe2B
layer as a function of boriding parameters
Slika 10: Diagram enakih debelin, ki opisuje nastanek Fe2B-plasti v
odvisnosti od parametrov boriranja
Acknowledgements
The work described in this paper was supported by a
grant of CONACyT and PROMEP, México. Also, the
authors wish to thank Ing. Martín Ortiz Granillo, the
Director of Escuela Superior de Ciudad Sahagún which
is part of Universidad Autónoma del Estado de Hidalgo,
México, for providing all the facilities necessary to
accomplish this research work.
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672 Materiali in tehnologije / Materials and technology 49 (2015) 5, 665–672
E. KALINOWSKA-OZGOWICZ et al.: KINETICS OF THE PRECIPITATION IN AUSTENITE HSLA STEELS
673–679
KINETICS OF THE PRECIPITATION IN AUSTENITE HSLA
STEELS
KINETIKA IZLO^ANJA V AVSTENITNIH HSLA-JEKLIH
El¿bieta Kalinowska-Ozgowicz1, Roman Kuziak2, Wojciech Ozgowicz3,
Klaudiusz Lenik1
1Lublin University of Technology, Fundamentals of Technology, Nadbystrzycka Str 38, 20-618 Lublin, Poland
2Institute for Ferrous Metallurgy, K. Miarki Str 12-14, 44-100 Gliwice, Poland
3Silesian University of Technology, Mechanical Engineering Faculty, Konarskiego Str 18a, 44-100 Gliwice, Poland
kalinowska-ozgowicz@tlen.pl
Prejem rokopisa – received: 2014-07-30; sprejem za objavo – accepted for publication: 2014-11-06
doi:10.17222/mit.2014.139
Presented are the results of an investigation of the kinetics of the static precipitation in HSLA steels with microalloying
additions of Nb, V, Ti and high N between the consecutive hot plastic-deformation cycles. The hot plastic deformation was
carried out using a torsional plastometer and a Gleeble thermomechanical simulator. The kinetics curves of the precipitation
process in the examined steels were determined based on the changes in the values of strain m in the flow curves during
continuous and interrupted deformations. The precipitation process in the austenite of the examined steels was investigated
using transmission electron microscopy of the test specimens quenched immediately after hot deformation and held isothermally
for various times. The precipitation effects inherited in the martensite were analysed with electron diffraction and X-ray
microanalysis.
Keywords: hot plastic deformation, kinetics of precipitation, HSLA steels, structure, Gleeble simulator
Predstavljeni so rezultati preiskave kinetike stati~nega izlo~anja v HSLA-jeklih, mikrolegiranih z Nb, V, Ti in visokim N med
zaporednimi cikli vro~e plasti~ne deformacije. Vro~a plasti~na deformacija je bila izvr{ena s torzijskim plastometrom in z
Gleeble-jevim termomehanskim simulatorjem. Krivulje kinetike procesa izlo~anja v preiskovanih jeklih so bile ugotovljene na
podlagi sprememb vrednosti napetosti m in krivulj te~enja med kontinuirnimi in prekinjenimi deformacijami. Proces izlo~anja v
avstenitu preiskovanih jekel je bil preiskovan s presevno elektronsko mikroskopijo preizkusnih vzorcev, kaljenih takoj po vro~i
deformaciji in razli~no dolgem izotermnem zadr`evanju. Podedovani u~inki izlo~anja v martenzitu so bili analizirani z
difrakcijo elektronov in rentgensko mikroanalizo.
Klju~ne besede: vro~a plasti~na deformacija, kinetika izlo~anja, HSLA-jekla, struktura, Gleeble-jev simulator
1 INTRODUCTION
Chemical compositions and structural changes of
HSLA steels during a hot plastic deformation are the
basic sources of the diversity of the mechanical pro-
perties of finished products. Shaping the mechanical
properties of this group of steels takes place by inducing
the relevant strengthening mechanisms and by develop-
ing proper phase compositions of steels.1–3 A decisive
role in shaping the mechanical properties of micro-
alloyed steels is played by the precipitates of secondary
interstitial phases (MX-type carbo-nitrides formed as a
result of alloying the microalloying elements of Nb, Ti,
V and N into steel, i.e., carbides and nitrides as well as
the products of their intersolubility – carbonitrides. The
stability of these phases depends mainly on their
chemical compositions and solution temperatures.4–8 The
knowledge of the basic thermodynamic data of the
analysed compounds makes it possible to use them in
hot-working processes such as: controlled rolling, rolling
with controlled recrystallisation, finish rolling and
thermomechanical treatments.9–11 HSLA microalloyed
steels with ferritic-pearlitic structures obtain relatively
high mechanical properties due to the grain refinement
and precipitation hardening.
2 EXPERIMENTAL PROCEDURE
The studies of the kinetics of the precipitation were
carried out on HSLA microalloyed constructional steels
with ferritic-pearlitic structures obtained from industrial
Materiali in tehnologije / Materials and technology 49 (2015) 5, 673–679 673
UDK 669.14:669.112.227.1:539.37 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 49(5)673(2015)
Table 1: Chemical compositions of investigated steels in mass fractions, w/%
Tabela 1: Kemijska sestava preiskovanih jekel v masnih dele`ih, w/%
Steel type
Concentration of the elements in w/%
C Mn Si P S Nb V Ti N Al Cu
B2 0.15 1.03 0.25 0.018 0.009 0.017 0.050 – 0.0070 0.004 –
G1 0.16 1.48 0.29 0.030 0.017 0.037 0.002 0.004 0.0098 0.010 0.049
heats. The chemical compositions of the examined steels
are given in Table 1.
For steels B2 and G1 the studies of the kinetics of the
precipitation of carbides and carbonitrides were carried
out and PTT (precipitation-temperature-time) diagrams
were developed for various austenite deformation condi-
tions and isothermal holding times of the following
deformation sequences.
A static precipitation in the non-deformed austenite
of the B2 steel was performed in a torsional plastometer
and test specimens of 6 mm in diameter and 10 mm in
gauge length were used. The specimens were heated at 3
°C/s to the austenitizing temperature of 1150 °C, held for
360 s, cooled to a deformation temperature of 900 °C at
approximately 10 °C/s and held isothermally for 6 s to
10800 s, then deformed at a rate of 4.4 s–1 until a failure
(torsion) took place. The heating was performed in the
plastometer’s resistance-heat furnace, in an argon atmo-
sphere. The austenitizing and deformation temperatures
were measured with a radiation pyrometer with an
accuracy of ± 1 °. The cooling of the test specimens to
room temperature, immediately after the hot deforma-
tion, was performed with a water jet from the plasto-
meter’s hydraulic system. An analysis of the value of
deformation corresponding to the maximum stress on the
 –  curves as a function of the isothermal holding time,
enabled us to determine3 the characteristic times of the
beginning (Ps) and the end (Pf) of the precipitation
process in the non-deformed austenite of the B2 steel.
For the G1 steel, the precipitation process was inve-
stigated in the non-deformed state and after a high-
temperature deformation of the austenite using the cyclic
axisymmetric-upsetting and stress-relaxation methods.
The investigations were performed on a Gleeble 3800
simulator using cylindrical test specimens of 7 mm in
diameter and 8.4 mm in length. The specimens were
heated at 3 °C/s to an austenitizing temperature of 1200
°C, held for 3 s, cooled to a deformation temperature of
1100 °C at approximately 1.0 °C/s and held isothermally
for 2 s to 1800 s, then deformed at 1.0 s–1 to  = 1. In the
stress-relaxation method, the precipitation process was
investigated in the austenite deformed to  = 0.2 at 800
°C to 1100 °C after an austenitization at 1200 °C, keep-
ing the same strain parameters as for the cyclic compres-
sion test. The heating of the test specimens in the
simulator was performed with the resistance method and
in an argon atmosphere. To reduce the coefficient of
friction, the faces of the test specimens were coated with
a graphite-and-tantalum foil.
The microstructures of the G1 steel after the selected
hot deformation and precipitation stages were deter-
mined with transmission electron microscopy (TEM)
using the selective electron diffraction and energy-dis-
persive X-ray spectroscopy (EDX) methods. In the
investigations, an S/TEM Titan 80-300 scanning electron
microscope by FEI, fitted with a field electron gun
(XFEG), an energy-dispersion spectrometer (EDS), an
external energy filter for imaging (EFTEM) and for
spectroscopy (EELS) as well as a system of BF/ADF/
HAADF detectors for the operations of the scanning
system were used. Observations were performed within
the scope of 245–300 kV, in a classic TEM system with a
spatial resolution of the network image below 0.10 μm
and using the microscope’s field emission gun to allow
the electron beam to be focused up to 0.3 nm during the
observation of high-resolution images (HRTEM). Micro-
structures were observed on the foils and small lamellas.
The foils were obtained by mechanical thinning and ion
polishing using a Gatan PIPS polishing machine, while
the lamella-formed preparations of 20 μm × 8 μm were
cut out with a gallium-focused ion beam from a specific
cross-sectional area of the test specimen and then
thinned with this ion beam to a thickness of approxima-
tely 50–70 nm using a FEI FIB Quanta 3D200 machine.
The lamella-preparation process was observed in situ
using the scanning electron microscope.
The procedure of a phase identification based on
electron diffraction was supported with the ELDYF
software for indexing and simulating electron diffraction
patterns.
3 RESULTS AND DISCUSSION
The results of the kinetics of the static precipitation
in the microalloyed constructional steels (B2 and G1) are
shown in Figures 1 to 6. In particular, for the B2 steel,
the kinetics of the static precipitation in the austenite not
deformed by the hot torsion was determined (Figure 1),
while, for the G1 steel, the precipitation process was
determined for the austenite in the non-deformed state
and after a high-temperature deformation (Figures 2 to
5). It was found that an increase in the holding time
before the deformation resulted in a distinct reduction in
the values of m and the corresponding maximum
stresses. This effect is equivalent to the reduction in the
degree of strain hardening of the austenite and results
from the growing share of the static precipitation during
E. KALINOWSKA-OZGOWICZ et al.: KINETICS OF THE PRECIPITATION IN AUSTENITE HSLA STEELS
674 Materiali in tehnologije / Materials and technology 49 (2015) 5, 673–679
Figure 1: Kinetics of the static precipitation in non-deformed
austenite of microalloyed steel B2, determined with a hot-deformation
test
Slika 1: Kinetika stati~nega izlo~anja v nedeformiranem avstenitu
mikrolegiranega jekla B2, dolo~ena s preizkusom vro~e deformacije
the isothermal holding. Thereby, the potential for dyna-
mic precipitation at the stage of hardening is decreasing,
which, in turn, results in a reduced degree of inhibition
of dynamic recrystallization and the observed changes in
m for the values of max. Therefore, in the absence of
precipitation before the deformation, the constant and
maximum value of m is obtained (Figures 1 and 2). It is
connected with the maximum solubility of Nb in the
austenite, potentially inhibiting the nucleation process of
dynamic recrystallisation so that the precipitation pro-
cess leads to a stabilisation of the m values at the mini-
mum level, observed at the stage of the stress increase in
the flow curves. The end of the first plateau and the
beginning of the next one on the kinetic curves were
adopted as the times of the beginning (Ps) and the end
(Pf) of the static precipitation in the examined micro-
alloyed steels.
The static precipitation process in the B2 steel after
the austenitization at 1150 °C and isothermal holding at
900 °C starts after approximately 740 s and ends after
approximately 6200 s (Figure 1). Before the start of the
precipitation in the B2 steel, the value of m is 1.51 and
after its completion it is m = 1.11. During the precipi-
tation the relationship between strain m and the isother-
mal holding time can be described with a linear function
(m = –0.43 · lg  + 2.67) whose intersection points deter-
mine the times of the beginning and the end of the
precipitation process.
In the case of the isothermal holding of microalloyed
steel G1 at 1100 °C after the austenitization at 1200 °C,
the relationship between m and the time is described in a
similar way. The first stage, the period of up to 47 s, is
described as m = 0.53, while the second stage (to appro-
ximately 1380 s) is described with the linear function
m = –0.1 lg  + 0.7. When the isothermal holding  >
1380 s, the third section with a stabilised value is des-
cribed with the formula m = 0.37. The inflexion points
of the function determine the times of the beginning
(Ps = 47 s) and the end (Pf = 1380 s) of the static precipi-
tation in the non-deformed austenite of the examined
microalloyed steel at the holding temperature of 1100 °C
(Figure 2).
The static precipitation process was also analysed in
the microalloyed steel G1 using the two-stage com-
pression method after the austenitization at 1200 °C and
the pre-strain of the test specimens ( = 0.2) before the
isothermal holding at 1100 °C. The diagram of the
changes in m during the isothermal holding is also cha-
racterised by three stages (Figure 3). In the initial stage,
there is a constant value, m = 0.42, then this value
changes linearly to settle again at a strain level of m =
0.62 for a longer holding time. The characteristic
inflexion points of the diagram determine similarly the
time of the beginning and the end of the static precipi-
tation in the non-deformed austenite. The process takes
place during the isothermal holding after the plastic
deformation from approximately 13 s to approximately
28 min.
For comparison, the static precipitation in the micro-
alloyed steel G1 was also investigated, with the stress-
relaxation method, after the austenitization at 1200 °C
and the preliminary compression of the test specimens
within a temperature range of 800–1100 °C, using the
Gleeble thermodynamic simulator. The times of the
beginning and the end of the static-precipitation process
were determined analytically and shown graphically in
diagrams (Figures 4 to 6). The values of the pre-strain
( = 0.2 and 0.4) were found to have a slight impact on
the time to the initiation of the precipitation (Ps was
approximately 17 s) at 900 °C, while for the maximum
value of the strain ( = 0.6) the corresponding time was
twice as high. The static precipitation at this temperature
ends, regardless of the size of the pre-strain, after
approximately 3 min (Figure 5). The precipitation at a
lower deformation temperature (800 °C) is evident in the
distinct plateau on the stress-relaxation curve and it takes
place within the period from Ps of approximately 30 s to
Pf of approximately 600 s (Figure 5). The studies imply
that for the austenite deformed before the isothermal
E. KALINOWSKA-OZGOWICZ et al.: KINETICS OF THE PRECIPITATION IN AUSTENITE HSLA STEELS
Materiali in tehnologije / Materials and technology 49 (2015) 5, 673–679 675
Figure 3: Kinetics of the static precipitation in deformed austenite of
microalloyed steel G1, determined with a hot-compression test
Slika 3: Kinetika stati~nega izlo~anja v deformiranem avstenitu
mikrolegiranega jekla G1, dolo~ena z vro~im tla~nim preizkusom
Figure 2: Kinetics of the static precipitation in non-deformed auste-
nite of microalloyed steel G1, determined with a hot-compression test
Slika 2: Kinetika stati~nega izlo~anja v nedeformiranem avstenitu
mikrolegiranega jekla G1, dolo~ena z vro~im tla~nim preizkusom
holding of the G1 steel, the time to the beginning of the
static precipitation is clearly shorter than the correspond-
ing time for the non-deformed austenite. The kinetic
analysis also revealed that the times of the end of the
static-precipitation process in both cases (non-deformed
austenite and the austenite deformed before the iso-
thermal holding) are comparable.
The analysis of the results of the kinetics obtained for
the examined microalloyed steel G1, considering the
austenite before the isothermal holding as well as the
conditions of the high-temperature deformation, allowed
PTT curves to be determined (Figure 6). The static-pre-
cipitation curves obtained for the austenite of the G1
steel in the investigated conditions of hot deformation
indicate that the precipitation process is the fastest at the
isothermal-holding temperature of 900 °C. The times of
the initiation and the end of the precipitation are 17 s and
210 s, respectively. The times of the initiation of the
precipitation at both higher (1000 °C) and lower
(800 °C) temperatures are longer, 22 s and 30 s, respec-
tively. The end of the precipitation process at 800 °C
takes place after approximately 10 min.
The precipitation process in the G1 steel was inve-
stigated on the test specimens quenched in water imme-
diately after the hot deformation and isothermal holding
for various periods. The phase-precipitation effects
E. KALINOWSKA-OZGOWICZ et al.: KINETICS OF THE PRECIPITATION IN AUSTENITE HSLA STEELS
676 Materiali in tehnologije / Materials and technology 49 (2015) 5, 673–679
Figure 6: Results of the studies of the kinetics of the static precipi-
tation (PTT) for microalloyed steel G1 subject to plastic deformation
with the hot-compression method
Slika 6: Rezultati {tudija kinetike stati~nega izlo~anja (PTT) za
mikrolegirano jeklo G1 po plasti~ni deformaciji z metodo vro~ega
stiskanja
Figure 7: a) Microstructure of martensite lath with Fe3C and TiN
precipitates, b) image in the dark field, c) diffraction, d) diffraction
solution; G1 steel, TA = 1200 °C, Tdef = 1100 °C,  = 1.0 s–1,  = 0.2,
w (water) = 30 s
Slika 7: a) Mikrostruktura latastega martenzita z Fe3C in izlo~ki TiN,
b) posnetek v temnem polju, c) difrakcija, d) difrakcija indeksirana;
jeklo G1, TA = 1200 °C, Tdef = 1100 °C,  = 1,0 s–1,  = 0,2, w (voda)
= 30 s
Figure 5: Stress-relaxation curve for microalloyed steel G1 subject to
a hot-compression test at 800 °C
Slika 5: Krivulja zmanj{evanja napetosti pri mikrolegiranem jeklu G1
po vro~em tla~nem preizkusu pri 800 °C
Figure 4: Stress-relaxation curves for microalloyed steel G1 subject to
a hot-compression test at 900 °C; Ps,f – times of the beginning and end
of the static precipitation process
Slika 4: Krivulje zmanj{evanja napetosti pri mikrolegiranem jeklu G1
pri vro~em tla~nem preizkusu na 900 °C; Ps,f – ~asi do za~etka in kon-
ca stati~nega procesa izlo~anja
inherited in the martensite were analysed with electron
diffraction and X-ray microanalysis.
The isothermal holding of the examined steel after
the hot deformation results in static-precipitation pro-
cesses. After a deformation at 1100 °C and an isothermal
holding for 30 s, dispersive precipitates with varied sizes
and morphologies were revealed in the martensite-lath
matrix with a high dislocation density (Figure 7).
In the martensite laths a TiN precipitation was
revealed in the dark field (Figure 7b) of the reflection
from plane (111)TiN. There are also reflections from Fe3C
in the diffraction pattern, indicating the likelihood of a
formation of lower bainite bands formed due to cooling.
The NbC particles with a spherical form and a size of
approximately 60 nm were also revealed in the marten-
site laths and at the former austenite-grain boundaries. In
addition, at the martensite-lath boundaries, the traces of
the retained austenite were revealed too. Similar results
of the phase identification for microalloyed steels are
also quoted in12–15.
An increase in the isothermal-holding time after the
plastic deformation of the examined steel up to approxi-
mately 3 h results in a formation of complex titanium
and niobium carbides with diverse morphologies and
different chemical compositions on the bainite and mar-
tensite laths (Figure 8). A microanalysis of the poly-
gon-shaped (Ti, Nb)C particle area of approximately 120
nm revealed a high concentration of Nb (Figure 8b,
point 1), whereas in the precipitates of 30 nm (Figure
8b, point 3), comparable niobium and titanium contents
were found.
The HRTEM investigations of the precipitation in the
martensitic matrix of the G1 steel austenitized at 1200 °C
and deformed at 900 °C, with a comparable strain rate
(1.0 s–1) and similar isothermal-holding times, but with
higher pre-strains ( = 0.6), allowed a determination of
the morphology of the observed particles and their
probable identification with the electron-diffraction
method.
In the test specimens held for 44 s and cooled in
water, there are titanium-containing nanometric precipi-
tates of approximately 7 nm in the martensite-lath
E. KALINOWSKA-OZGOWICZ et al.: KINETICS OF THE PRECIPITATION IN AUSTENITE HSLA STEELS
Materiali in tehnologije / Materials and technology 49 (2015) 5, 673–679 677
Figure 9: a) NbC spherical precipitate in martensite lath, b) image in
the dark field, c) diffraction, d) diffraction solution; G1 steel, TA =
1200 °C, Tdef = 900 °C,  = 1.0 s–1,  = 0.6, w(water) = 10800 s
Slika 9: a) Okrogli izlo~ek NbC v martenzitni lati, b) posnetek v
temnem polju, c) difrakcija, d) difrakcija indeksirana; jeklo G1, TA =
1200 °C, Tdef = 900 °C,  = 1,0 s–1,  = 0,6, w(voda) = 10800 s
Figure 8: a) Microstructure of martensite lath with (Ti,Nb)C precipi-
tates, b) microanalysis of micro-areas (1, 2, 3); G1 steel, TA = 1200
°C, Tdef = 1100 °C,  = 1.0 s–1,  = 0.2, w (water) = 10800 s
Slika 8: a) Mikrostruktura martenzitne late s (Ti,Nb)C-izlo~ki, b) mi-
kroanaliza mikropodro~ij (1, 2, 3); jeklo G1, TA = 1200 °C, Tdef =
1100 °C,  = 1,0 s–1,  = 0,2, w (voda) = 10800 s
matrix. Using diffraction analysis and following the air
cooling of the specimens in ferrite, precipitates of a
similar size were observed and identified as TiC. After
the hot deformation at 1100 °C and the same holding
time (44 s) and air cooling, spheroidal precipitates (Ti,
Nb) of approximately 15 nm were found in the examined
steel. The electron microphotographs showed that the
precipitates in the martensitic matrix, especially after the
deformation at 900 °C, are not evenly distributed and
that in many micro-areas the spacing of the particles was
less than 100 nm.
An increase in the holding time up to 185 s following
the deformation at 900 °C results in a formation of NbC
precipitates with a size between 20 nm and approxima-
tely 30 nm in the recrystallised austenite matrix (Figure
9). Niobium carbide precipitates were identified with
electron diffraction. Their shape and distribution were
revealed in the dark field of the reflection of (101)NbC
(Figures 9b to 9d). These observations indicate that the
precipitation-hardening process is not uniformly distri-
buted within the entire volume of a test specimen. It can
also be assumed that the distribution of the precipitates
in the solution matrix – ’ – reflexes the distribution of
the preferred nucleation sites in the hot-deformed auste-
nite substructure. However, the confirmation of this
effect is difficult as the austenite substructure is subject
to degradation during the phase transformation that takes
place due to quench hardening. Micro-bands of twins in
the martensitic matrix and precipitates in the form of
discrete particles at the boundaries between these bands
were identified. The chemical composition of these
precipitates does not differ from that of the precipitates
in the matrix.
It needs to be emphasised that Nb forms carbides
and nitrides with equal lattices (NaCl type) that are
completely intersoluble. The similarity of the lattices of
NbC and NbN formed in the examined steel makes their
precise identification difficult. As NbC and NbN form
isomorphous phases in the solid solution, these precipi-
tates can be regarded as carbonitrides. Such suggestions
are made frequently.16 In addition, the results of the
chemical analyses of the Nb(C, N) precipitates extracted
from the deformed austenite in the steels with micro-
alloying elements confirmed that Nb carbonitride is rich
in carbon. Compared with a stoichiometric composition,
it is similar to that of NbC.17
4 CONCLUSIONS
The results obtained with the investigations and the
analysis support the following conclusions:
1. In the examined HSLA steels microalloyed with Nb,
Ti, V and N, the precipitation of carbonitrides in both
deformed and non-deformed austenite after the iso-
thermal holding corresponds to PTT (precipitation-
temperature-time) curves.
2. The precipitation of NbC, TiN and (Ti, Nb)C during
a plastic deformation of microalloyed steels (B2, G1)
with torsion and compression in a temperature range
of 800–1100 °C takes place in the torsion test for up
to about 5 s.
3. The static-precipitation process in microalloyed steel
B2, determined by means of hot-torsion tests after an
austenitization at a temperature of 1150 °C and iso-
thermal holding at 900 °C, starts after about 740 s
and finishes after about 6200 s. A similar process of
the precipitation of steel G1, hot compressed after an
austenitization at 1200 °C and isothermal holding at
1100 °C takes place in the course of holding from
about 50 s to about 1400 s.
4. The static precipitation of the G1 steel examined with
the compression and stress-relaxation method within
the temperature range of 800–1100 °C depends
mainly on the pre-strain temperature and size. In the
investigated deformation conditions the maximum
isothermal rate of this process was found at the tem-
perature of 900 °C. The times of the beginning and
the end of the static-precipitation process were about
17 s and about 210 s, respectively.
5. The microstructure of the ’ solution matrix is
formed by a partially twined martensite lath with a
high dislocation density, while the precipitates with
equal frequencies of the particles of single NbC and
TiC carbides, NbN and TiN nitrides and also of com-
plex (Ti, Nb)C carbides were identified with electron
diffraction and energy-dispersive X-ray spectrometry
(EDX).
6. The diameters of the precipitates revealed in the G1
steel range between approximately 7 nm and approxi-
mately 120 nm, and their growth is proportional to
the isothermal holding time after the hot deformation.
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2 R. Kuziak, T. Bold, Y. W. Cheng, Microstructure control of
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Materiali in tehnologije / Materials and technology 49 (2015) 5, 673–679 679

D. BOLIBRUCHOVÁ, M. @IHALOVÁ: COMBINED INFLUENCE OF V AND Cr ON THE AlSi10MgMn ALLOY ...
681–686
COMBINED INFLUENCE OF V AND Cr ON THE AlSi10MgMn
ALLOY WITH A HIGH Fe LEVEL
VZAJEMNI VPLIV V IN Cr NA ZLITINO AlSi10MgMn Z VISOKO
VSEBNOSTJO Fe
Dana Bolibruchová, Mária @ihalová
Department of Technological Engineering, Faculty of Mechanical Engineering, University of @ilina, Univerzitná 8215/1, 010 26 @ilina,
Slovakia
maria.zihalova@fstroj.uniza.sk
Prejem rokopisa – received: 2014-07-31; sprejem za objavo – accepted for publication: 2014-10-09
doi:10.17222/mit.2014.146
Removing impurities from aluminium alloys increases the final cost of castings. The most common impurity in Al-Si based
alloys is iron, whose removal is the main problem for secondary Al-Si alloys. Several techniques of eliminating iron effects have
been introduced, but the most common one is the use of iron "correctors" that change the morphology of iron-based
intermetallic phases. The use of the correctors like Co and Mn is well described but other elements such Ni, Cr and V need to be
studied more extensively. In the present paper, the influence of V and the combined influence of V and Cr are analysed. The aim
of the article is to determine the microstructure, the mechanical properties and the solidification behaviour of the AlSi10MgMn
cast alloy with a high iron level treated with V and Cr.
Keywords: AlSi10MgMn, iron correctors, combined effect, vanadium, chromium
Odstranjevanje ne~isto~ iz aluminijevih zlitin pove~uje kon~ne stro{ke pri ulitkih. Najbolj pogosta ne~isto~a v zlitinah na osnovi
Al-Si je `elezo. Uvedenih je bilo ve~ tehnik za odstranitev vpliva `eleza, vendar je najpogostej{a uporaba "korektorjev" `eleza,
ki spremenijo morfologijo intermetalnih faz na osnovi `eleza. Uporaba korektorjev, kot sta Co in Mn, je dobro opisana, pri
drugih elementih, kot na primer Ni, Cr in V, pa so potrebne {e dodatne raziskave. V ~lanku je analiziran vpliv V in vzajemni
vpliv V in Cr. Namen ~lanka je dolo~iti mikrostrukturo, mehanske lastnosti in vedenje pri strjevanju livarske zlitine
AlSi10MgMn z visoko vsebnostjo `eleza in te zlitine, obdelane z V in Cr.
Klju~ne besede: AlSi10MgMn, korektorji `eleza, vzajemni vpliv, vanadij, krom
1 INTRODUCTION
Aluminium has been acquiring increasing signifi-
cance over the past few decades due to its excellent
properties and a diversified range of applications.1 Over
the years, aluminium alloys have been specially
developed to meet the increasing demands of today’s
industry. This development has resulted in the production
of smaller and light-weight components that comply
with property, environmental and other specifications.
Recently, the appeal for recycling the resources has
become more intensive with the increasing public
awareness of the need to preserve the materials and
energy and to protect the environment. However, the
increasing use of recycled aluminium casting alloys
requires a strict procedure of removing the harmful
effects of impurity elements.2
The most common impurity in aluminium Al-Si
alloys is iron. Iron is highly soluble in liquid aluminium
and its alloys, but it has a very low solubility in the solid
state and so it tends to combine with other elements to
form intermetallic-phase particles of various types.3–5
The morphologies of iron-containing intermetallics are
often found to cause negative effects, where the
plate-like phases, such as -Al5FeSi, are considered to be
more harmful than the script-type phases, such as
-Al15(FeMn)3Si2.5
When Fe is present in excess of specified levels,
various methods are available to reduce its harmful
influence. The conventional method is to add some
chemical "correctors" to change the morphology from
the platelet -Al5FeSi (a brittle form) to the globular or
script forms (less brittle forms). The globular or script
Fe-rich phases are considered not to lead to brittleness.
Alternative methods of iron reduction are described in
various references.6,7 Manganese is the most commonly
used and the least expensive element for Fe neutralisa-
tion in Al-Si alloys.2 If an iron amount exceeds a value
of the mass fraction w = 0.45 %, the recommended
addition of Mn should not be lower than half of the iron
amount;7 however, script-like particles containing Mn are
still observed at the Mn-to-Fe ratio of 0.17.8 Other iron
correctors also include Co, Cr, Mo, Ni, V and Be.3,4,6,7 Cr
improves the strength at indoor and higher temperatures
and mildly deteriorates the elongation. The presence of
Cr phases (CrFe)4Si4Al13 and (CrFe)5Si8Al2 can increases
the brittleness.9,10 Vanadium refines the grains of
aluminium alloys. Together with Ti and Mo, it increases
the hot-cracking resistance and decreases the porosity.9
The influence of V occurs with an amount of 0.05–0.15
Materiali in tehnologije / Materials and technology 49 (2015) 5, 681–686 681
UDK 669.715:621.74.046:66.046.51 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 49(5)681(2015)
% of the melt mass. V reduces the length of -phase
platelets Al5FeSi and, together with Ni and T, it has a
significant influence on improving the mechanical
properties (Rm and A5). Similarly to Fe and Co, vanadium
has a negative influence on the fluidity of Al-Si alloys.9
There is also a significant beneficial influence of vana-
dium on the hardness.10
Until now a large number of experiments focusing on
the influence of individual iron correctors have been
performed9,10 but few papers about a combined influence
of iron correctors have been published. Petrík et al.11
evaluated the properties of an Al-Si alloy containing Fe,
Ni and Mn. Kumari et al.2 investigated the effects of
individual and combined additions of Be, Mn, Ca and Sr
on the solidification behaviour of the structure and
mechanical properties of the AlSi7Mg0.3Fe0.8 alloy.
These papers refer to the beneficial influences of the
combined additions of the selected iron correctors;
however, there is very little evidence of an influence of
combined vanadium-and-chromium additions on Al-Si
alloys.
In this paper the influences of vanadium and of an
addition combining V and Cr are analysed in the
AlSi10MgMn alloy with an elevated iron amount.
2 MATERIALS AND METHODOLOGY
Aluminium alloy AlSi10MgMn was used as the
experimental material. The chemical composition of the
used alloy is presented in Table 1. In the first stage of
the experiments, the commercial AlSi10MgMn alloy was
alloyed with iron. Iron was added in the amount of
40000 μg/g using the AlFe10 master alloy. After obtain-
ing the alloy with a high Fe mass content (1 % of Fe), V
and Cr were added. Vanadium was added in the mass
fraction of 0.2 % and the combined addition of V and Cr
included 0.2 % of V and 0.5 and 1 % of Cr. The required
amounts of V and Cr were added as the AlV10 and
AlCr20 master alloys. For each experiment, the alloy
was melted in a graphite crucible using an electric-resis-
tance furnace. The melts were not further modified, grain
refined or purified. The melts were poured into perma-
D. BOLIBRUCHOVÁ, M. @IHALOVÁ: COMBINED INFLUENCE OF V AND Cr ON THE AlSi10MgMn ALLOY ...
682 Materiali in tehnologije / Materials and technology 49 (2015) 5, 681–686
Table 1: Chemical composition of AlSi10MgMn alloy
Tabela 1: Kemijska sestava zlitine AlSi10MgMn
Element Si Mg Mn Fe Ti Zn Cu V Cr Al
Amount (w/%) 10.220 0.277 0.108 0.448 0.046 0.029 0.047 0.010 0.006 balance
Figure 1: a) AlSi10MgMn + w(Fe) 1 %, b) detail of Fe phases, etched
in 20 mL H2SO4 + 100 mL H2O
Slika 1: a) AlSi10MgMn + w(Fe) 1 %, b) detajl Fe-faz, jedkano v 20
mL H2SO4 + 100 mL H2O
Figure 2: a) AlSi10MgMnFe + w(V) 0.2 %, b) detail of Fe phases,
etched in 20 mL H2SO4 + 100 mL H2O
Slika 2: a) AlSi10MgMnFe + w(V) 0,2 %, b) detajl Fe-faz, jedkano v
20 mL H2SO4 + 100 mL H2O
nent moulds preheated to 200 °C after reaching a
temperature of (760 ± 5) °C. From the poured castings,
specimens for the tensile testing, hardness testing and
microstructure evaluation were prepared. The solidifica-
tion behaviour of the melts was analysed with the meas-
uring equipment containing NiCr-Ni thermocouples.
3 RESULTS
3.1 Microstructure
Figure 1 shows the typical microstructure of the
AlSi10MgMn alloy with an addition of a mass fraction
1 % of Fe. It can be seen that platelets of the -Al5FeSi
phases are present in the interdendritic regions as well as
precipitated along with the eutectic Si. The needles of
the -Al5FeSi phase are evenly dispersed in the alloy and
only small amounts of script-like particles are present.
The microstructure of the alloy with the V addition is
shown in Figure 2. The needle-like phases present in the
microstructure were of a different colour after the
etching compared to alloy containing 1 % of Fe. The
distribution and dimensions of the needles are similar to
those of the alloy without the vanadium addition.
The combined addition of V and Cr caused the
occurrence of sludge phases in the alloys’ microstruc-
tures (Figures 3 and 4). In the alloy with a higher Cr
level, a higher amount of sludge phases is present as the
sludge-phase formation is mostly influenced by the Cr
level in the alloy.7 Iron-based intermetallic phases are
smaller after both additions of Cr (w = 0.5 % and 1.0 %).
3.2 Mechanical properties
An examination of the ultimate tensile strength
(UTS) and elongation of the alloys was performed in line
with the EN ISO 6892-1 standard. The ultimate tensile
strength and elongation of the alloys are shown in
Figure 5. The addition of Fe to the commercial
AlSi10MgMn alloy caused a decrease in the UTS and
elongation (sample No. 2). The highest improvement in
both UTS and elongation was observed after the vana-
dium addition (w = 0.2 %). The combined addition of V
and Cr decreased the alloy’s tensile properties. The
tensile strength of the alloy with the combined addition
is even lower than that of the alloy containing a high iron
amount. The same is true of the elongation in the case of
the addition of V and Cr in the mass fractions of 0.2 %
and 0.5 %. The hardness of the alloys was evaluated with
the Brinell hardness-measuring method in line with the
EN ISO 6506-1 standard. In Figure 6 it can be seen that
after the additions of all the investigated elements,
Brinell hardness increased. The most significant increase
D. BOLIBRUCHOVÁ, M. @IHALOVÁ: COMBINED INFLUENCE OF V AND Cr ON THE AlSi10MgMn ALLOY ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 681–686 683
Figure 4: a) AlSi10MgMnFe + w(V) 0.2 % + w(Cr) 1.0 %, b) detail of
Fe phases, etched in 20 mL H2SO4 + 100 mL H2O
Slika 4: a) AlSi10MgMnFe + w(V) 0,2 % + w(Cr) 1,0 %, b) detajl
Fe-faz, jedkano v 20 mL H2SO4 + 100 mL H2O
Figure 3: a) AlSi10MgMnFe + w(V) 0.2 % + w(Cr) 0.5 %, b) detail of
Fe phases, etched in 20 mL H2SO4 + 100 mL H2O
Slika 3: a) AlSi10MgMnFe + w(V) 0,2 % + w(Cr) 0,5 %, b) detajl
Fe-faz, jedkano v 20 mL H2SO4 + 100 mL H2O
(42 %) in the alloy’s hardness occurred for the V-alloyed
sample.
3.3 Thermal analysis
A cylindrical sample with a thermocouple placed in
the middle of the casting was used for obtaining the ther-
mal curve. Figures 7 to 10 show the cooling curves and
their first derivatives and Table 2 lists the alloys, thermal
arrests and phases formed in the analysed alloys. The
first derivative is a measure of the instantaneous cooling
rate along the cooling curve and is used to indicate the
presence of minor slope changes on the curves.2 Figure
7 shows the first thermal-arrest point of the
AlSi10MgMn alloy with the mass fraction 1.0 % of Fe at
591.2 °C, where the formation and growth of Al nuclei
occur. After this arrest, the temperature of the solidifying
alloy continues to decrease. The second thermal arrest
occurs at 575 °C. This is caused by the latent heat of fu-
sion of the iron intermetallic phases. The last thermal-
arrest point occurs at 566.7 °C, corresponding to the
eutectic-Si formation and growth. The cooling curve
does not indicate the thermal arrest at the temperature
corresponding with the formation of the Mg2Si phase. It
might have been caused by a very low Mg level in the
measuring place and solidification conditions that do not
result in the formation of the Mg2Si phase. The thermal
arrests corresponding with the primary -Al after the
additions of V and also V with Cr occur at higher tempe-
ratures. The eutectic temperature after the combined
addition of V and Cr decreased compared to that of the
D. BOLIBRUCHOVÁ, M. @IHALOVÁ: COMBINED INFLUENCE OF V AND Cr ON THE AlSi10MgMn ALLOY ...
684 Materiali in tehnologije / Materials and technology 49 (2015) 5, 681–686
Figure 6: Brinell hardness of AlSi10MgMn alloy: sample No. 1 –
commercial AlSi10MgMn alloy, sample No. 2 – AlSi10MgMn +
w(Fe) 1 %, sample No. 3 – AlSi10MgMnFe + w(V) 0.2 %, sample No.
4 – AlSi10MgMnFe + w(V) 0.2 % + w(Cr) 0.5 % and sample No. 5 –
AlSi10MgMnFe + w(V) 0.2 % + w(Cr) 1.0 %
Slika 6: Brinellova trdota zlitine AlSi10MgMn: vzorec {t. 1 –
komercialna zlitina AlSi10MgMn, vzorec {t. 2 – AlSi10MgMn +
w(Fe) 1 %, vzorec {t. 3 – AlSi10MgMnFe + w(V) 0,2 %, vzorec {t. 4 –
AlSi10MgMnFe + w(V) 0,2 % + w(Cr) 0,5 % in vzorec {t. 5 –
AlSi10MgMnFe + w(V) 0,2 % + w(Cr) 1,0 %
Figure 7: Cooling curve and its first derivative of AlSi10MgMn alloy
with w(Fe) 1.0 %
Slika 7: Ohlajevalna krivulja in njen prvi odvod zlitine AlSi10MgMn
z w(Fe) 1,0 %
Figure 9: Cooling curve and its first derivative of AlSi10MgMnFe
alloy with w(V) 0.2 % and w(Cr) 0.5 %
Slika 9: Ohlajevalna krivulja in njen prvi odvod zlitine
AlSi10MgMnFe z w(V) 0,2 % in w(Cr) 0,5 %
Figure 5: Tensile strength and elongation of AlSi10MgMn alloy:
sample No. 1 – commercial AlSi10MgMn alloy, sample No. 2 –
AlSi10MgMn + w(Fe) 1 %, sample No. 3 – AlSi10MgMnFe + w(V)
0.2 %, sample No. 4 – AlSi10MgMnFe + w(V) 0.2 % + w(Cr) 0.5 %
and sample No. 5 – AlSi10MgMnFe + w(V) 0.2 % + w(Cr) 1.0 %
Slika 5: Natezna trdnost in raztezek zlitine AlSi10MgMn: vzorec {t. 1
– komercialna zlitina AlSi10MgMn, vzorec {t. 2 – AlSi10MgMn +
w(Fe) 1 %, vzorec {t. 3 – AlSi10MgMnFe + w(V) 0,2 %, vzorec {t. 4
- AlSi10MgMnFe + w(V) 0,2 % + w(Cr) 0,5 % in vzorec {t. 5 –
AlSi10MgMnFe + w(V) 0,2 % + w(Cr) 1,0 %
Figure 8: Cooling curve and its first derivative of AlSi10MgMnFe
alloy with w(V) 0.2 %
Slika 8: Ohlajevalna krivulja in njen prvi odvod zlitine
AlSi10MgMnFe z w(V) 0,2 %
alloy with the Fe addition and the difference between the
eutectic temperatures of the alloy containing mass
fractions 0.5 % and 1.0 % of Cr and the alloy without V
and Cr is 17.9 °C and 14 °C, respectively. This leads to
the formation of a finer eutectic Si compared to the alloy
with the iron addition.
Table 2: Thermal arrests and phases formed in AlSi10MgMn alloy
with a high iron level without and with V and Cr additions
Tabela 2: Toplotni zastoji in faze, ki nastajajo v AlSi10MgMn zlitini z
visoko vsebnostjo `eleza ter brez dodatka V in Cr ali z njima
Alloy Thermal arrest(Temperature (°C)) Phases formed
AlSi10MgMn +
w(Fe) 1.0 %
591.2
575.0
566.7
Primary -Al
Iron intermetallics
Eutectic Si
AlSi10MgMnFe +
w(V) 0.2 %
597.7
576.0
563.2
Primary -Al
Iron intermetallics
Eutectic Si
AlSi10MgMnFe +
w(V) 0.2 % + w(Cr)
0.5 %
599.4
575.5
548.8
Primary -Al
Iron intermetallics
Eutectic Si
AlSi10MgMnFe +
w(V) 0.2 % + w(Cr)
1.0 %
603.4
574.1
552.7
Primary -Al
Iron intermetallics
Eutectic Si
4 DISCUSSIONS
The microstructure of the AlSi10MgMn alloy after
the iron addition in the mass fractions of 1.0 % contains
a large amount of iron-based intermetallic phases, mostly
in the form of platelets. After the addition of w(V) = 0.2
%, the platelet-like particles were still present but after
the same etching technique as used with the iron-treated
alloy, the needles obtained a different colour. This might
have been due to a different chemical composition of the
iron-based platelets. The combined addition of V and Cr
caused the formation of sludge phases in the alloy. As
the size and amount of the sludge phases are mostly
influenced by Cr, more of them were observed after the
addition of V and Cr in the mass fractions of 0.2 % and
1.0 %, respectively.
The formations of different intermetallic phases were
reflected on the alloy’s mechanical properties. The UTS
of the alloy after the Fe addition decreased compared to
the commercial AlSi10MgMn alloy. The decrease in the
UTS is connected with a higher number of the brittle
needle-like intermetallics in the iron-treated alloy. The
addition of V (w = 0.2 %) improved the tensile strength
of the iron-polluted alloy, but its value was still lower
than for the commercial alloy. Vanadium is known as a
grain-refining element and this characteristic might have
been one of the reasons for the tensile-strength improve-
ment. Also, it is possible that the chemical interaction
between the iron-based phases and vanadium caused the
change in the phase colour after the etching, leading to a
decrease in the deleterious effect of such phases on the
UTS. The combined addition of V and Cr in both mass
fractions (0.2 % V + 0.5 % Cr and 0.2 % V + 1.0 % Cr)
leads to a significant decrease in the alloy’s UTS. The
reason for the decreased tensile properties is the forma-
tion of hard and brittle sludge phases. These phases are
brittle and cannot withstand the same stress as the ductile
aluminium matrix.
The elongation of the alloys was influenced similarly
by the analysed elements. The highest elongation was
found for the commercial alloy (1.38 %), followed by the
alloy treated with iron and vanadium (1.24 %). Brinell
hardness was positively influenced by every used
element. The highest Brinell hardness was found for the
alloy containing Fe (w = 1.0 %) and V (w = 0.2 %). The
reason for the increase in the hardness value after the
addition of iron is probably a higher amount of iron
intermetallic phases with a high value of microhardness.
The same effect on Brinell hardness might have occurred
due to the presence of sludge phases after the combined
addition of V and Cr. The microhardness of the sludge
phases can reach 800–1000 HV 7 while the aluminium-
matrix microhardness is only 60–100 HV. The increase
in the hardness after the V addition might have been
caused by a high number of iron intermetallics and also
by the strengthening of the aluminium matrix due to
dissolved vanadium.
The thermal analysis of the alloys did not show any
significant differences between the analysed alloys. It
was found that the thermal arrest of the primary Al for-
mation occurs at higher temperatures after the addition
of V and Cr compared to the alloy containing the mass
fraction 1.0 % of Fe.
5 CONCLUSION
Several conclusions can be drawn from the obtained
results of the individual and combined effects of V and
Cr on the AlSi10MgMn alloy with a high iron level:
• The iron addition (w = 1.0 %) to the commercial
AlSi10MgMn alloy caused the formation of a high
amount of iron-based platelet particles. The tensile
properties (UTS and elongation) were decreased but
the hardness increased.
D. BOLIBRUCHOVÁ, M. @IHALOVÁ: COMBINED INFLUENCE OF V AND Cr ON THE AlSi10MgMn ALLOY ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 681–686 685
Figure 10: Cooling curve and its first derivative of AlSi10MgMnFe
alloy with w(V) 0.2 % and w(Cr) 1.0 %
Slika 10: Ohlajevalna krivulja in njen prvi odvod zlitine
AlSi10MgMnFe z w(V) 0,2 % in w(Cr) 1,0 %
• The addition of V (w = 0.2 %) to the iron-containing
alloy led to changes in the mechanical properties.
The UTS and elongation were positively affected and
the measured values almost reached the values of the
commercial alloy. V also caused Brinell hardness to
increase to the highest value (91 HBW) compared to
the other analysed alloys.
• The combined influence of V and Cr led to the for-
mation of sludge phases. These phases decreased the
mechanical properties compared to the vanadium-
treated alloy. The UTS and Brinell hardness were
almost the same at the chromium levels of w = (0.5
and 1.0) %. The elongation of the samples was the
lowest with the combined V and Cr addition of the
mass fractions 0.2 % and 0.5 %, respectively.
• The thermal analysis showed that alloying elements
V and Cr do not have a significant influence on the
solidification behaviour of the AlSi10MgMn alloy
with a high iron level.
• The best combination of the mechanical and foundry
properties was obtained with the vanadium addition
in the mass fraction of 0.2 %.
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(2004), 73–79
D. BOLIBRUCHOVÁ, M. @IHALOVÁ: COMBINED INFLUENCE OF V AND Cr ON THE AlSi10MgMn ALLOY ...
686 Materiali in tehnologije / Materials and technology 49 (2015) 5, 681–686
C. C. KUO: A RELIABLE APPROACH TO A RAPID CALCULATION OF THE GRAIN SIZE ...
687–691
A RELIABLE APPROACH TO A RAPID CALCULATION OF THE
GRAIN SIZE OF POLYCRYSTALLINE THIN FILMS AFTER
EXCIMER LASER CRYSTALLIZATION
ZANESLJIV NA^IN HITREGA IZRA^UNA VELIKOSTI ZRN V
POLIKRISTALNI TANKI PLASTI PO UV-LASERSKI
KRISTALIZACIJI
Chil-Chyuan Kuo
Department of Mechanical Engineering, Ming Chi University of Technology, No. 84, Gungjuan Road, Taishan Taipei Hsien 243, Taiwan
jacksonk@mail.mcut.edu.tw
Prejem rokopisa – received: 2013-10-08; sprejem za objavo – accepted for publication: 2014-11-27
doi:10.17222/mit.2013.237
Excimer laser crystallization (ELC) is the most commonly employed technology for fabricating low-temperature polycrystalline
silicon (LTPS) thin films. The grain size of polycrystalline thin films after ELC is usually determined with a manual calculation,
which includes certain disadvantages, i.e., human error and is time-consuming and exhausting. To mitigate these disadvantages,
a high-efficiency approach to calculating the grain size of polycrystalline thin films automatically is proposed. It was found that
the selected-boundary-definition approach is a promising candidate for calculating the grain size of polycrystalline thin films.
The savings in the analysis time is up to 75 %. The average error rate of the measurement can be controlled within 8.33 %.
Keywords: low-temperature polycrystalline silicon, automatic grain-size analysis, excimer laser crystallization
UV-laserska kristalizacija (ELC) je najbolj pogosto uporabljena tehnologija za izdelavo nizkotemperaturne polikristalne tanke
plasti silicija (LTPS). Velikost zrn v polikristalni tanki plasti po ELC se navadno dolo~a z ro~nim izra~unom, ki pa ima nekatere
pomanjkljivosti, kot je ~love{ka napaka, in je ~asovno potratna in utrujajo~a. Za ubla`itev teh pomanjkljivosti je predlagan zelo
u~inkovit, avtomati~en na~in za izra~un velikosti zrn v polikristalni tanki plasti. Ugotovljeno je, da je pribli`ek k selektivnemu
dolo~anju meje obetajo~ na~in za izra~un velikosti zrn polikristalnih tankih plasti. Prihranek ~asa za analizo je do 75 %.
Povpre~ni odmik napake pri meritvah je okrog 8,33 %.
Klju~ne besede: nizkotemperaturni polikristalni silicij, avtomatska analiza velikosti zrn, UV-laserska kristalizacija
1 INTRODUCTION
High-performance complementary metal-oxide-semi-
conductor (CMOS) circuits on glass are essential for the
system-on panel (SOP) technology, which has potential
applications in various information devices including
cell-phones, laptop computers and large-size flat panel
television sets. Polycrystalline silicon (poly-Si) thin
films have been widely used as CMOS gates, thin-film
transistors (TFTs), solar cells and various other applica-
tions in semiconductor-device technology. Excimer laser
crystallization (ELC) is an industrial technique used for
preparing poly-Si thin films on commercially available,
inexpensive glass substrates for the development of
high-performance TFTs in active-matrix flat panel dis-
plays.1–6 A rapid deposition of the laser-energy density,
on a nanosecond time scale, onto the surface region of
the an amorphous-silicon (a-Si) thin film leads to its
melting and recrystallisation into a poly-Si thin film,
while keeping the glass substrate at a low temperature.
The final quality of the device depends significantly on
the phase-transformation mechanisms which need to be
manipulated precisely for obtaining poly-Si thin films
with a large grain size and a good uniformity. The
phase-transformation mechanisms of a-Si thin films have
been extensively investigated using an in-situ optical
diagnostic technique during ELC in the previous stu-
dies.7–17 Numerous researches have been done for fabri-
cating large-grained poly-Si thin films because the
performance of TFTs is significantly affected by the size
of the poly-Si thin films after ELC.18–23 Until now, how-
ever, the grain size of the poly-Si thin films after ELC
has usually been determined with a manual calculation.
This approach included certain disadvantages such as
human error and was time-consuming and exhausting.
Therefore, a high-efficiency approach is proposed in this
work for calculating the grain size of polycrystalline thin
films efficiently and accurately using the Image-Pro
software.24,25
2 EXPERIMENT
Figure 1 shows a schematic illustration of the experi-
mental set-up for ELC. The sample has a stacked struc-
ture consisting of a thick 300 nm SiO2 capping layer and
a thick 90 nm a-Si layer formed on a thick 0.7 mm
non-alkali glass substrate (Corning 1737). All the films
were prepared with plasma-enhanced chemical vapor de-
position (PECVD). These samples were then dehydroge-
Materiali in tehnologije / Materials and technology 49 (2015) 5, 687–691 687
UDK 532.6:669.782:620.186 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)687(2015)
nated with a thermal treatment at 500 °C for 2 h to
reduce the hydrogen content in order to prevent the
ablation caused by a sudden hydrogen eruption during
ELC.26 The samples were then held by self-closing
tweezers at the end of a cantilever beam fixed to an x – y
precision translation stage. The x- and y-axis displace-
ments of the two stages can be accurately manipulated
(resolution = 0.625 μm). The movement of the focusing
lens mounted onto a z-axis stage was precisely controlled
to adjust the desired excimer laser fluences for
crystallization. The pulsed excimer laser-energy levels
were monitored using a laser power meter (Vector H410
SCIENTECH). The variation in the pulse-to-pulse
excimer laser energy was found to be less than 5 %. The
a-Si thin films were irradiated with an excimer laser
beam (	 = 351 nm, repetition rate = 1 Hz, LAMBDA
PHYSIK COMPex 102) with laser fluences ranging from
100 mJ/cm2 to 500 mJ/cm2. A stainless-steel slit (2 mm ×
15 mm) located in the optical path of the excimer laser
was employed to transform the incident Gaussian beam
into a rectangular beam spot with a better than ± 10 %
energy variation. All the experiments were performed at
ambient temperature and pressure.
After ELC, the microstructural analyses of the
annealed poly-Si thin films were carried out using field
emission scanning electron microscopy (FE-SEM) with
JEOL JSM-6500F. Before the FE-SEM observation, the
crystallized silicon films were Secco-etched in order to
highlight the grain boundaries (GBs) and intra-grain
defects.27 The acceleration electron beam energy for
FE-SEM was 15 kV (a resolution of 1.5 nm). Six ap-
proaches (count, auto-split, watershed split, limited
watershed split, boundary definition and selected-boun-
dary definition) were employed for calculating the grain
size of the poly-Si thin films.
3 RESULTS AND DISCUSSION
Figure 2 shows a typical SEM micrograph of the
poly-Si thin films after ELC. The grain size can be
determined from the longest length inside the grain
boundary. To determine the best approach to replace the
tradition manual-calculation method, a test SEM micro-
graph was selected to be investigated. Figure 3 shows
the grain-size calculation result using the manual-calcu-
lation approach. The total number of the grain size of
poly-Si was 24. The largest grain size, the smallest grain
size and the average grain size were (333.3, 26.6 and
152.2) mm, respectively. Figure 4 shows the grain-size
calculation result using six different approaches. Figure
5 shows the variation in the counts of the grain size for
seven different calculation approaches. The average error
rate for the approaches of count, auto-split, watershed
split, limited watershed split, boundary definition and
selected-boundary definition was (21.32, 19.33, 20.76,
20.76, 16.29 and 8.33) %, respectively. As one can see,
the selected-boundary-definition approach provides the
lowest average error rate in the grain-size calculation
compared with the tradition manual-calculation approach
in this case. The average error rates for the approaches of
count, auto-split, watershed split and limited watershed
split were higher than those of the approaches of boun-
dary definition and selected-boundary definition because
the Image-Pro software cannot precisely evaluate the
C. C. KUO: A RELIABLE APPROACH TO A RAPID CALCULATION OF THE GRAIN SIZE ...
688 Materiali in tehnologije / Materials and technology 49 (2015) 5, 687–691
Figure 3: Grain-size calculation result using manual-calculation
approach
Slika 3: Izra~un velikosti zrn z ro~nim {tetjem
Figure 2: SEM micrograph of poly-Si thin films after ELC
Slika 2: SEM-posnetek poli-Si tanke plasti po ELC
Figure 1: Schematic illustration of the experimental set-up for ELC
Slika 1: Shematski prikaz eksperimentalnega sestava za ELC
C. C. KUO: A RELIABLE APPROACH TO A RAPID CALCULATION OF THE GRAIN SIZE ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 687–691 689
Figure 6: SEM micrograph of poly-Si thin film of case study 1 28
Slika 6: SEM-posnetek poli-Si tanke plasti pri {tudiju primera 1 28
Figure 5: Variation in the counts of the grain size for seven different
calculation approaches
Slika 5: Razlike v izra~unu velikosti zrn pri sedmih razli~nih na~inih
izra~una
Figure 4: Grain-size calculation result using six different approaches of: a) count, b) auto-split, c) watershed split, d) limited watershed split, e)
boundary definition and f) selected-boundary definition
Slika 4: Rezultat izra~una velikosti zrn s {estimi razli~nimi na~ini: a) {tetje, b) avtomatska razdelitev, c) razdelitev po razvodnicah, d) omejena
razdelitev po razvodnicah, e) definicija mej in f) selektivna definicija mej
grain boundary of a SEM micrograph. To reduce the
average error rate of the measurement, the two ap-
proaches of boundary definition and selected-boundary
definition were further applied. The average error rate of
the measurement was still not acceptable, though the
boundary-definition approach can reduce the average
error rate of the measurement. Finally, the selected-
boundary-definition approach was applied. The selec-
ted-boundary-definition approach provides the best
accuracy of the grain-size calculation because the grain
boundary of the SEM micrograph was traced first and
then calculated using the Image-Pro software.
To evaluate the accuracy of the selected-boundary-
definition approach, two case studies were applied to
investigate the average error rate. Figure 6 shows a SEM
micrograph of the poly-Si thin film of case study 1.28
Figure 7 shows the variation in the counts of the grain
size for two different calculation approaches in case
study 1. The average error rate of the measurement was
only 1.28 %. In this case, the total time for calculating
the grain size with the manual calculation was approxi-
mately 16 h. However, the total calculating time was
drastically reduced to approximately 4 h using the
selected-boundary-definition approach. The saving in the
analysis time was up to 75 %. Figure 8 shows a SEM
micrograph of the poly-Si thin film of case study 2.
Figure 9 shows the variation in the counts of the grain
size for two different calculation approaches in case
study 2.28 The average error rate of the measurement was
only 4.03 %. In this case, the total time for calculating
the grain size with the manual calculation was approxi-
mately 12 h. However, the total calculating time was
drastically reduced to approximately 3 h using the
selected-boundary-definition approach. The saving in the
analysis time was up to 75 %. It is worth noting that the
average error rate of the measurement was obviously
smaller than for the test sample because the grain
boundary was clear for the two SEM micrographs. Thus,
a SEM micrograph with a clear grain boundary is critical
for calculating the grain size with the Image-Pro soft-
ware when the Secco etching is employed29–31.
As discussed above, the Image-Pro software is a
powerful tool for analyzing the grain size of the poly-Si
thin films after ELC. The saving in the analysis time is
up to 75 % and the average error rate of the measurement
can be controlled within 8.33 % when using the compu-
ter-calculation approach compared with the manual-cal-
culation approach.
4 CONCLUSIONS
A simple and highly efficient approach for calculat-
ing the grain size of the poly-Si thin films after ELC was
successfully demonstrated. A SEM micrograph with a
clear grain boundary is critical for calculating the grain
size with the Image-Pro software. The selected-boun-
dary-definition approach was proved to be a promising
candidate for calculating the grain size of the poly-Si
thin films efficiently and accurately. The saving in the
analysis time was up to 75 % and the average error rate
of the measurement can be controlled within 8.33 %
when compared with the manual-calculation approach.
C. C. KUO: A RELIABLE APPROACH TO A RAPID CALCULATION OF THE GRAIN SIZE ...
690 Materiali in tehnologije / Materials and technology 49 (2015) 5, 687–691
Figure 9: Variation in the counts of the grain size for two different
calculation approaches in case study 2
Slika 9: Razlike v izra~unu velikosti zrn pri dveh razli~nih na~inih
izra~una, za primer 2
Figure 7: Variation in the counts of the grain size for two different
calculation approaches in case study 1
Slika 7: Razlike v izra~unu velikosti zrn pri dveh razli~nih na~inih
izra~una za primer 1
Figure 8: SEM micrograph of poly-Si thin film of case study 2 28
Slika 8: SEM-posnetek poli-Si tankih plasti pri {tudiju primera 2 28
Acknowledgements
This work was financially supported by the National
Science Council of Taiwan under contract nos. NSC
102-2221-E-131-012 and NSC 101-221-E-131-007. The
skillful technical assistance of Mr. Kun-Wei Wu of the
Ming Chi University of Technology is gratefully
acknowledged.
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C. C. KUO: A RELIABLE APPROACH TO A RAPID CALCULATION OF THE GRAIN SIZE ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 687–691 691

H. BASAK, K. KAPTAN: EFFECTS OF DIFFERENT STIRRER-PIN FORMS ON THE JOINING QUALITY ...
693–701
EFFECTS OF DIFFERENT STIRRER-PIN FORMS ON THE
JOINING QUALITY OBTAINED WITH FRICTION-STIR WELDING
VPLIV RAZLI^NIH OBLIK VRTILNIH KONIC NA KVALITETO
SPOJA PRI TORNEM VRTILNEM VARJENJU
Hudayim Basak1, Kadir Kaptan2
1Gazi University, Technology Faculty, Department of Industrial Design Engineering, Ankara, Turkey
2Gazi University, Technical Education Faculty, Department of Mechanical Education, Ankara, Turkey
hbasak@gazi.edu.tr
Prejem rokopisa – received: 2014-04-02; sprejem za objavo – accepted for publication: 2014-10-15
doi:10.17222/mit.2014.062
In this study, 20 different pins that could be employed in friction-stir welding were manufactured. These pins and the Al7075 T6
material were joined with the friction-stir-welding technique. Tensile tests, microstructure and microhardness measurements,
scanning-electron-microscopy (SEM) and energy-dispersion-spectrometry (EDS) analyses were conducted on the components
joint by friction-stir welding. Designs of the weld and the tip that achieve the best performance were determined.
Keywords: friction-stir welding, stirrer-pin forms
V tej {tudiji je bila napravljena primerjava 20 razli~nih konic, ki se uporabljajo pri tornem vrtilnem varjenju. Konice in Al7075
T6 material so bili spojeni s tornim vrtilnim varjenjem. Pri komponentah, spojenih z vrtilno tornim varjenjem, so bili izvr{eni
natezni preizkusi, pregledana mikrostruktura in izmerjena mikrotrdota. Uporabljena je bila vrsti~na elektronska mikroskopija
(SEM) in energijsko disperzijska rentgenska spektroskopija (EDS). Dolo~ena je bila tudi oblika zvara in konice, kjer so bile
dose`ene najbolj{e lastnosti.
Klju~ne besede: torno vrtilno varjenje, oblika konic
1 INTRODUCTION
Friction-stir welding (FSW) is a welding method
undergoing continuous development that was originally
introduced in England (Cambridge) in 1991,1,2 and went
through a major development in the 1990s.3,4 Newly
developed materials generally require modern joining
techniques. With the development of new alloys in the
past twenty years, major advances were also made
regarding the welding of these materials. Friction-stir
welding (FSW) was successfully employed in the weld-
ing of aluminum alloys that are difficult to weld with the
traditional melt-welding methods, especially those that
are subjected to age hardening.5,6
The problems arising from joining aluminum and
Al-alloys with the traditional melt-welding methods sti-
mulated the researchers to develop new joining methods.
When joining aluminum alloys (especially those sub-
jected to age hardening) with the melt-welding methods,
a high heat input may culminate in these materials
leading to a high thermal expansion and crack formation
in the welding seam due to wide solidification intervals.7
In addition, it is known that the high heat input in arc
welding leads to the formation of low melting-point
phases at the grain boundaries in the region under the
effect of the heat in aluminum alloys and, therefore,
cracks at the grain boundaries of the region during the
solidification. Another problem encountered when
joining aluminum alloys subjected to age hardening with
the melt-welding methods is a reduction in the hardness
and strength due to the solution of hardening precipitates
in the welding seam and an excessive aging of the heat-
effected zone.8–10 This situation results in mechanical
disconformities in the welding zone. Due to these speci-
fic reasons, solid-state welding methods (diffusion, fric-
tion and friction-stir welding) provide huge advantages
for the welding of these materials.
Friction-stir welding, which is a solid-state welding
method, is an appropriate alternative welding method for
joining the materials whose welding construction is diffi-
cult, due to its unique advantages like a brief welding
duration, minimum surface preparation and ease of auto-
mation.11 With this method, no protective gas, additional
welding metal or opening of the welding grooves on the
components to be welded are necessary. The welding
quality of the welded joints obtained with friction-stir
welding does not show variation with respect to the type
of material.10 Even though the method is suitable for
several materials, its most significant area of application
is the welding of aluminum and aluminum alloys. Alu-
minum alloys joined with this method exhibit decent
mechanical properties.12 The FSW method can be com-
fortably applied in all positions.13,14 In the studies con-
ducted for aluminum materials, it was found that appro-
priate joining can be performed for plate thicknesses
between 1–75 mm.15
The tool-pin profile and welding parameters play
major roles in achieving the quality of friction-stir weld-
Materiali in tehnologije / Materials and technology 49 (2015) 5, 693–701 693
UDK 621.791/.792:669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)693(2015)
ing. Friction-stir-welding studies used different para-
meters and stirrer pins made for these studies. The effect
of the tool-pin profile on mechanical properties was the
subject of a few studies.16–18 The effect of the tool-pin
profile on the bonding and mechanical properties was
investigated together with certain parameters such as the
rotational speed,16 welding speed19 shoulder diameter20,21
and axial force.22 Also, in recent years, some studies
have been made on the optimization of the tool geometry
using statistical analyses and finite-element methods.23,24
2 EXPERIMENTS AND MATERIAL
The Al7075 material was chosen to be employed in this
study. For the FSW method, this material was chosen due
to its superior hardness, tensile strength, joining capability
and resistance properties. In addition, the Al7075 material
is used in many industrial areas (especially in the space and
aeronautical industry) thanks to its mechanical properties.
Stirrer-pin profiles were formed by taking screw-thread
types (metric (triangular), trapezoidal, circular, square and
saw-shaped) into consideration. Along with that, in the case
when no screw thread is cut, triangular, cylindrical, conical
H. BASAK, K. KAPTAN: EFFECTS OF DIFFERENT STIRRER-PIN FORMS ON THE JOINING QUALITY ...
694 Materiali in tehnologije / Materials and technology 49 (2015) 5, 693–701
Figure 1: Drawing of a manufactured stirrer pin
Slika 1: Risba izdelane konice
1.
Square screwed
stirrer pin
2.
Circular screwed
stirrer pin
3.
Trapezoidal screwed
stirrer pin
4.
Triangular screwed
stirrer pin
5.
Saw-type screwed
stirrer pin
6.
3-groove square
screwed stirrer pin
7.
3-groove circular
screwed stirrer pin)
8.
3-groove triangular
(metric) screwed
stirrer pin
9.
3-groove trapezoidal
screwed stirrer pin
10.
3-groove saw-type
screwed stirrer pin
11.
4-groove square
screwed stirrer pin
12.
4-groove circular
screwed stirrer pin
13.
4-groove triangular
(metric) screwed
stirrer pin
14.
4-groove trapezoidal
screwed stirrer pin
15.
4-groove saw-type
screwed stirrer pin
16.
Unscrewed conical
profiled stirrer pin
17.
Unscrewed 3-groove
circular profiled
stirrer pin
18.
Unscrewed 4-groove
circular profiled
stirrer pin
19.
Unscrewed
triangular profiled
stirrer pin
20.
Unscrewed square
profiled stirrer pin
Figure 2: Manufactured stirrer pins
Slika 2: Izdelane vrste konic
and square pin profiles were created in order to compare
the joining qualities of the pin profiles. Stirrer pins were
manufactured on the basis of the drawing given in Figure
1. The manufacturing principle for all the stirrer pins was
taken into account in accordance with this drawing. As it is
theoretically known, screw types are divided into two
groups: connection and movement types. The triangular
screw type is the most frequently used connection bolt.
Since the profile gradient in triangular profile screws in-
creases the frictional resistance (μ' = tan 
' = μ/(cos (/2)),
the friction per unit area for triangular threads is at the
highest level. As for the movement bolts, the friction per
unit area is very low.25 For the profiles designed on the
basis of this theoretical information, we can predict that,
of all the pin designs, the highest joining-strength quality
is achieved with the triangular screw. We can also state
that low joining strengths are achieved with the move-
ment-type screws (circular and square). The AISI/SAE
4340 (DIN 1.6582, EN 34CrNiMo6) material was
selected as the stirrer-pin material and it was subjected to
a heat treatment for a hardness of 50–52 HRC. The
stirrer pins were made ready for use with the FSW
method. The manufactured stirrer pins are presented in
Figure 2.
2.1 Friction-stir-welding method
The friction-stir welding process is shown in Figure
3. The FSW method is a type of welding that was deve-
loped using the friction-stir-welding technique. This
welding method consists of moving an immersed shoul-
der pin (probe) revolving at a high speed over two plates
butted to each other at a particular speed.26 Even though
aluminum alloys are usually difficult to weld, this me-
thod can be used for enhancing the welding perfor-
mance. Researchers are focused on joining aluminum
alloys. This welding method was applied on aluminum
materials thinner than 1 mm and thicker than 35 mm and
it was found that very good mechanical properties were
achieved.27–31
The parameters used in the tests are given in Table 1.
Table 1: Test parameters
Tabela 1: Parametri preizkusa
Material Aluminum alloy (Al7075 T6),a (80 × 120) mm prismatic component
Stirrer pins 20 stirrer pins created on the basis of
screw types
Feed rate 28–40 mm/min
Number of
revolutions
710–1200 r/min
Bench used for
burnishing process
Taksan 40T1500, a CNC milling bench
with a vertical processing center
During the construction stage of friction-stir welding,
the values of the movement rate and rotational speed
were determined separately for different stirrer pins. The
reason for this is the fact that a rotational speed yielding
very good results for one pin does not provide for an effi-
cient joining for another pin. Therefore, the values that
yield good welding results were determined in terms of
increasing or decreasing the movement and rotational
speeds. In this way, the welding operation was carried
out (Figure 4).
The friction-stir welding was started with ungrooved
stirrer pins and the welding was performed with square,
circular, saw-type, trapezoidal and triangular screwed
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Figure 4: Performing the welding operation
Slika 4: Izvedba varjenja
Figure 3: a) Tool approaching the workpiece, b) adjusting the tool
before the welding, c) drawing of the tool30,31
Slika 3: a) Orodje se pribli`uje obdelovancu, b) prilagoditev orodja
pred varjenjem, c) risba orodja30,31
stirrer pins. In the welding operations involving the
circular, trapezoidal and triangular screwed 3-groove
pins, the components were welded at a rotational speed
of 1200 r/min and a movement rate of 28 mm/min. In the
cases of the 4-groove pins, fractures occurred in the
triangular and saw-type screwed pins. A good joint was
achieved at the rotational rate of 1200 r/min and the
movement rate of 28 mm/min.
Lastly, the joining operation was conducted with
untapped stirrer pins. The joining could not be achieved
with these pins when carried out on the basis of standard
workpiece values. In order to ensure the joining, the
movement rate was reduced to 28 mm/min and the rota-
tion was increased to 1200 r/min according to our
previous experiences and the joint was achieved in this
manner. Afterwards, the process was finalized by per-
forming joining operations with the 4-groove stirrer pins.
Since the triangular and saw-type threaded stirrer
pins fractured at the selected standard rotational speed
and movement rate, the stirrer-pin rotational speed was
reduced to 710 r/min and the movement rate to 28
mm/min and the joining was performed in this manner;
lastly, the untapped stirrer pins were used for the joining
operation. When the joining operation with these pins
was conducted at the standard workpiece values, the
joining did not take place. In order to ensure the joining,
the movement rate was 28 mm/min and the rotational
speed was 1200 r/min.32
2.2 Tensile test
For the entire FSW using different stirrer pins, three
tensile-test samples were extracted from the welding
samples joined with the FSW method. The samples were
cut to the desired dimensions with the erosion machine
and prepared for testing. The tensile-test samples were
prepared according to standards TSE 138 (EN
10002-1).33 They were milled in a ZWICK tensile-testing
device. During the test, the pulling speed was determined
as 5 mm/min and the frontload as 10 N.
2.3 Microstructure, scanning electron microscopy
(SEM) and energy dispersion spectrometry (EDS)
For the purpose of conducting a microstructural anal-
ysis, the welding samples joined with the screwed stirrer
pins that display the highest and the lowest tensile
strengths (triangular and circular screwed pins: samples
2, 4, 7, 8, 13) were selected. The samples were prepared
with the standard metallographic sample-preparation
method. At the end of this procedure, the samples were
seared with a macro-searing machine.
A Prior light microscope was employed for scanning
the microstructures and its images were transferred into a
computer environment. In order to obtain the profiles of
the microhardness distributions from the welded regions
towards the main metal in the samples, loads of 100 g
were applied on the samples with adjusted micro-
structures, and their Vickers (HV0.1) microhardness levels
were measured with a Shimadzu microhardness device.
The measurements were repeated at least three times,
within certain intervals, and the graphs were obtained
after taking the arithmetic averages and standard devi-
ations. In addition, in order to determine the coating
status of the stirrer-pin samples according to the pin geo-
metry, a scanning electron microscope was employed.
For this purpose, a Jeol 6060 LV SEM device was used.
To prove that the coated material was welded, an EDS
analysis was conducted using the IXRF (X-ray fluore-
scence) system.
3 RESULTS AND DISCUSSION
3.1 Tensile-rest results and discussion
Duplicates of the tensile-test samples prepared
according to the standards were created. After this stage,
the sample components were ready for the tensile test.
Significant variations in the tensile-stress values of the
used stirrer pins depending on their geometries were
observed. The samples were usually ruptured from the
welded zone to the main metal. The rupture points of
samples 2, 4 and 13 are shown in Figure 5.
The welded connections created with the FSW ope-
rations using the screwed pins exhibit higher tensile
strengths compared to those of the unscrewed pins.
The obtained tensile values are displayed as a graph
in Figure 6. In this figure, stirrer pin “0” indicates the
unwelded Al7075 material.
The samples formed with the screwed stirrer pins
exhibited tensile-strength values close to that of the main
component. The highest tensile strength was achieved by
the sample component formed with the triangular
(metric) screwed stirrer pin. The lowest tensile strength
was exhibited by the unscrewed conical profiled stirrer
pin. Among the sample components made with the
unscrewed stirrer pins, the highest tensile strength was
achieved with the square profile stirrer pin. The tensile
strengths of the samples acquired with the unthreaded
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Figure 5: Rupture points on the samples used in the tensile test
Slika 5: Mesto poru{itve pri nateznem preizkusu vzorcev
stirrer pins were well below the desired level. Therefore,
the unthreaded pins were not used in the other analysis.
The lowest tensile strength among the sample compo-
nents made with the screwed stirrer pins was achieved
with the circular screwed stirrer pin. On average, the
lowest tensile strength among the sample components
made with both the grooved and ungrooved screwed
stirrer pins was achieved with the square screwed stirrer
pin. In general, better tensile strengths were achieved
when grooves in the screwed stirrer pins were opened.
Only with the triangular screw does the ungrooved
screwed stirrer pin provide an exception.
In order for an FSW operation to be effective, the
surface of the stirrer pin needs to be large. For this
reason, it is believed that the parameter of the surface
area, which depends on the stirrer pin’s geometry,
directly affects the FSW operation. If the surface area
increases, the emerging quantity of heat is higher as
friction is more dominant. Thus, it can be stated that a
smoother stir can be obtained when joining two compo-
nents. However, for some samples of this study, possibly
due to the stirring geometry, the stir at the welding root
was not homogeneous. In such cases, a microstructure-
characterization study was also performed.
3.2 Microstructure and characterization of the welding
interface
The interface microstructures of the samples used in
the FSW process are shown in Figures 7a to 7f. In
sample 2 (the circular screwed stirrer pin), microfrac-
tures were found at the root of the weld. These fractures
probably caused the mixture not to be effective. Hence,
the tensile strengths of these samples turned out to be
low (Figure 7a). It is seen that the FSW operations
involving samples 4, 7 and 8 (triangular (metric) screwed
stirrer pin, 3-groove circular screwed stirrer pin and
3-groove triangular (metric) screwed stirrer pin, respec-
tively) were more successful. No problems such as
disintegration or cracking between the weld and the main
metal were observed (Figures 7b to 7d). However, the
formation of thin cracks seen on sample 13 (the 4-groove
triangular (metric) screwed stirrer pin) is thought to be
resulting from the thermal expansion (Figure 7f).
The samples prepared for the microstructural analysis
were used for microhardness measurements. During
these measurements, a 100 g (0.1 HV) pressure in terms
of HV (Vickers) was applied for 5 s. The hardness was
measured at 3 different zones, starting at the weld zone,
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Materiali in tehnologije / Materials and technology 49 (2015) 5, 693–701 697
Figure 6: Tensile-test values of the welding samples created with
FSW
Slika 6: Vrednosti, dobljene pri nateznem preizkusu vzorcev,
zvarjenih s FSW
Figure 8: Comparison of microhardness values for samples nos. 2 and
4
Slika 8: Primerjava vrednosti mikrotrdote vzorcev {t. 2 in 4
Figure 7: Microstructure images of samples joined with stirrer pins
nos. 2, 4, 7, 8, 12 and 13
Slika 7: Posnetki mikrostrukture vzorcev, spojenih z vrtilno konico {t.
2, 4, 7, 8, 12 in 13
at distances of (0.25, 0.5, 1, 2 and 3) mm. The deter-
mined microhardness average and the standard deviation
were calculated. Graphs of the determined values were
prepared for comparison.
In Figures 8 to 10 the graphs displaying the micro-
structure profiles of the samples subjected to the FSW
operation, measured at specific intervals from the weld
zone towards the main metal are given. These graphs
provide significant information about whether the FSW
operations were conducted correctly or not. In Figure 9,
the hardness profiles belonging to samples 7 and 8 relate
to the FSW operation conducted with the 3-groove cir-
cular screwed stirrer pin and 3-groove triangular (metric)
screwed stirrer pin. It is observed that the hardness
profiles of the samples exhibit similar tendencies after
both situations. No significant variation is observed in
the hardness tendencies from the weld zone towards the
main metal. This situation explicitly demonstrates that
the other mechanical properties of these samples may
also be superior. However, it is seen that the hardness
relatively increased in sample 8, made with the 3-groove
triangular (metric) screwed stirrer pin, at about a distance
of 0.25 mm from the weld zone. The potential reason
behind this situation may be the thinning of crystallized
grains as a consequence of excessive heat and sudden
cooling taking place at this region. As it is known, a
deformation of crystal grains of a material whose recry-
stallization mechanism was deformed, takes place as a
result of the formation of subgrain boundaries at the end
of dynamic recovery that occurs due to a polygonization
of dislocations when the material’s temperature reaches
between 1/3–2/3 of its melting temperature. Probably,
this mechanism took place in sample 8. At the same
time, the tensile values of samples 7 and 8 can be ex-
plained with their hardness profiles and they turn out to
be similar. In general, it is observed that the average
hardness tendencies of these samples are relatively lower
compared to the other samples. This may be considered
to be leading to the relatively lower tensile strengths of
these samples. The fact that the average hardness of
these samples turned out to be low can be attributed to a
larger temperature increase in these samples during the
FSW operation. It can be shown that the higher friction
provided by the stirrer pins used for these samples causes
the temperature of the samples to become higher.
Figure 8 shows the hardness profiles for samples 2
and 4 after the FSW operation was conducted with the
circular screwed stirrer pin and triangular (metric)
screwed stirrer pin, respectively. It was observed that the
distribution of the hardness for sample 4, the material
with which the triangular (metric) screwed stirrer pin
was used and which yielded the best result in the study,
is more homogenous compared to all the other samples.
This demonstrates that both the weld zone and the main
metal were more modified in terms of the microstructure
and chemical composition. This situation can also be
seen clearly in Figure 7b. When the hardness profile for
sample 2, whose FSW operation was done with the cir-
cular screwed stirrer pin, is examined, it is deduced from
the increasing hardness that in this sample, deformation
hardening might have occurred in the region near the
joining area (at the 0–0.5 mm interval). However, ex-
treme deformation hardening may lead to microfractures
like the ones that occurred in this sample due to an
excessively increased density of dislocations. These
microfractures, in turn, led to the sample’s low tensile
strength. It is likely that the decreasing effectiveness of
the FSW was due to inadequate heat. Along with the
inadequate level of the heat, an excessive plastic
deformation might have lead to deformation hardening.
In Figure 10, the hardness profiles for samples 12
and 13, whose FSW operations were conducted with the
4-groove circular screwed stirrer pin and 4-groove
triangular (metric) screwed stirrer pin, respectively, are
demonstrated. Significant drops and rises in sample 13,
in the regions near the joining area, stand out, while the
hardness profile of sample 12 is smoother compared to
sample 13. A possible reason behind this is thought to be
the capillary fractures arising from the thermal expansion
as seen from Figure 7f as well. For welded joints, this is
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698 Materiali in tehnologije / Materials and technology 49 (2015) 5, 693–701
Figure 10: Comparison of microhardness values for samples nos. 12
and 13
Slika 10: Primerjava vrednosti mikrotrdote vzorcev {t. 12 in 13
Figure 9: Comparison of microhardness values for samples nos. 7 and
8
Slika 9: Primerjava vrednosti mikrotrdote vzorcev {t. 7 in 8
perceived to be a critical problem. However, if no such
problem occurred in this sample, we could expect its
hardness profile to be higher and smoother and, corres-
pondingly, we could expect it to exhibit the highest
tensile strength.
In order to see the changes in the physical structures
of the stirrer pins after a welding operation, SEM images
of some stirrer pins (triangular and circular screwed)
were taken. The stirrer-pin images taken with SEM are
given in Figure 11.
When the SEM images were analyzed, as seen at the
first sight, a higher amount of the joining material was
found to be covering the circular screwed stirrer pin
compared to the triangular screwed stirrer pin. It was
also observed that the circular screwed stirrer pin demon-
strated the lowest strength among the screwed types
during the tensile test and the highest joining strength
was obtained with the triangular screwed stirrer pin. On
the basis of this observation, we can state that one reason
why the joining performed with the circular screwed
stirrer pin failed to display an adequate strength is the
fact that a portion of the material was being deposited on
the stirrer pin. In addition, it can be asserted that the
grooves employed on the stirrer pins yield higher joining
strengths than the ungrooved stirrer pins.
3.3 EDS analysis of a stirrer pin
For the purpose of making a partial estimation of the
stirrer-pin lifetime, an EDS analysis of a stirrer pin was
carried out (Figure 12). The EDS analysis was con-
ducted in the marked zone on the stirrer pin and the
materials found in it were determined.
After the analysis, no fragment was found to be
broken from the stirrer pin’s material. According to this
result, we can assume that the stirrer pin can have a long
lifetime. Also, we can presume that since the Al-alloy
that collects over the stirrer-pin cavities and grooves due
to the heat at each joining, a decrease in the performance
of the stirrer pin will not take place.
4 DISCUSSION
The internal structure formed at the weld zone after
the friction-stir welding consists of three different
regions: the recrystallized region (RCR), the thermo-
mechanically affected region (TAR) and the region under
the influence of heat.
It was observed that the hardness increased in the
weld zone of sample 2 (the circular screwed stirrer pin)
and the hardness distribution from the weld zone towards
the main metal decreased. However, there was no note-
worthy variation in the hardness distribution between the
weld zone and the main metal in sample 4 (the triangular
screwed stirrer pin). Homogenous distributions of the
microstructure and the hardness are essential for
achieving the superior mechanical properties expected
from the welded joints. On the other hand, the yield and
tensile strength are proportional to the hardness. It is
believed that even though the hardness level at the weld
zone for sample 4 turned out to be lower than that of
sample 2, the fact that its microstructure and hardness
distributions are more homogenous leads to an increase
in the sample’s tensile strength. The microfractures seen
at the welding root of sample 2 indicate that an effective
stir did not take place and it is thought that this caused
the tensile strength to be low.
Even though the hardness values for samples 7 and 8
were relatively a bit higher at the weld zones than at the
main metal, it can be stated that they exhibited homo-
genous hardness distributions. Moreover, the hardness
levels of both samples are similar in all the regions.
Therefore, it is understood that both pins yield results
similar to these properties.
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Figure 12: Stirrer pin on which EDS analysis was conducted
Slika 12: Vrtilna konica z ozna~enim mestom EDS-analize
Figure 11: SEM images of stirrer pin: a) no. 4 (triangular screwed), b)
2 (circular screwed), c) 8 (3-groove triangular screwed) and d) 13
(4-groove triangular screwed)
Slika 11: SEM-posnetki vrtilnih konic: a) {t. 4 (trikotni vijak), b) 2
(kro`ni navoj), c) 8 (trikotni vijak s tremi zarezami) in d) 13 (trikotni
vijak s {tirimi zarezami)
Even though the hardness distribution in sample 13
(the triangular screw stirrer pin) was expected to be
homogenous, or just slightly higher at the midpoint of
the weld zone, it was found to be higher in the region
close to the main metal. The probable reason for this is
an excessive increase in the temperature in the unmelted
region close to the main metal because the friction
stirring of the pin with this geometry is more dominant
and, along with that, the grain structure becomes finer
due to the influence of fast cooling. As for the hardness
distribution in sample 12 (the circular screw stirrer pin) it
is much more homogenous.
In general, it was observed that an increase in the
hardness in the weld zones or regions close to the joining
areas of the analyzed and discussed samples leads to an
increase in their tensile strengths.
It was observed that the tensile values for the samples
whose microstructures and microhardnesses demonstrate
homogenous distributions are similar to those of the
main metal. It was seen that the size and shape of the
regions formed as a result of joining depend on the
geometry of the stirrer pin. Hence, it can be concluded
that the geometry of the stirrer pin has a significant
effect on the mechanical behavior.
5 CONCLUSIONS
In this study, joining operations were carried out by
employing the friction-stir welding technique. On the
basis of the samples joined in this way, the stirrer pin
providing the best performance was determined with
tensile tests, microstructure characterizations, SEM and
EDS.
As a result of the conducted studies, the results given
below were obtained:
• It was observed that the welding seam made with the
FSW is smoother than the seams obtained with the
other methods.
• Among the stirrer pins, the best performance and
highest joining strength were achieved with the
triangular thread. It is suggested that this type is
preferred in stirrer-pin designs.
• Neither the joining surface nor the tensile strength
turned out to be at the desired level in the welds
performed with the unscrewed stirrer pins.
• Among the joints made with the screwed stirrer pins,
the poorest results were obtained with the circular
threaded stirrer pins.
• As we had theoretically predicted, the lowest
strengths were achieved with the conductive screws.
• Using the FSW method, a tensile-strength value close
to the one of the main metal can be achieved.
• In the stirrer-pin designs, screwed types should be
preferred as they facilitate the flow of the welding
seam.
• The pins with large friction surfaces facilitate more
homogenous and better joints.
• An increase in the hardness in the regions of the weld
zone or close to the joining area increases the ten-
sile-strength values.
• A larger friction surface of a stirrer pin increases the
tensile strength of the welded joint.
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Materiali in tehnologije / Materials and technology 49 (2015) 5, 693–701 701

L. PAZDERA et al.: MONITORING EARLY-AGE CONCRETE WITH THE ACOUSTIC-EMISSION METHOD ...
703–707
MONITORING EARLY-AGE CONCRETE WITH THE
ACOUSTIC-EMISSION METHOD AND DETERMINING THE
CHANGE IN THE ELECTRICAL PROPERTIES
PREGLED SVE@EGA BETONA Z METODO AKUSTI^NE EMISIJE
IN DOLO^ANJEM SPREMEMB ELEKTRI^NIH LASTNOSTI
Lubos Pazdera1, Libor Topolar1, Marta Korenska1, Tomas Vymazal1,
Jaroslav Smutny1, Vlastimil Bilek2
1Brno University of Technology, Faculty of Civil Engineering, Veveri 331/95, 602 00 Brno, Czech Republic
2ZPSV a.s., Uhersky Ostroh, Czech Republic
pazdera.l@fce.vutbr.cz
Prejem rokopisa – received: 2014-07-18; sprejem za objavo – accepted for publication: 2014-11-05
doi:10.17222/mit.2014.112
Concrete is the most popular building composite material (CM). Its long-term aging properties depend on the mixture, the
setting and the curing. When concrete is produced and used in the place where we want it to be for the whole of its life, it is very
sensitive and easily ruined. Curing is one of the things that we do to keep concrete protected during the first week or so of its
life: we maintain proper temperature (neither too hot nor too cold) and dampness. The acoustic-emission method and the
electrical-property measurement technique are applied to monitor early-age concrete. The relationships between the acoustic-
emission activity, the temperature and the electrical properties, e.g., the resistivity and the capacity of concrete at an early age
were studied in this research. W. Chen et al. studied the microstructure development of hydrating cement pastes during early
ages using non-destructive methods including the ultrasound P-wave propagation-velocity measurement and non-contact
electrical-resistivity tests. Nevertheless, there are not many references about the continuous study of the properties of concrete
during an early age. A long-time monitoring of concrete properties is necessary to determine its lifetime and quality. The
acoustic-emission method was proven to be a very advantageous tool for a non-destructive monitoring of structural micro-
changes during the lifetime of concrete. The differences between hardened concrete mixtures detected with acoustic emission
can partly determine their properties at the age of 28 d. The basic concrete property is its compressive strength after 28 d;
nevertheless, this can change over a long time period, thus a continuous measuring needs to be applied.
Keywords: concrete, acoustic emission, electrical properties, early age, lifetime, curing
Beton je najbolj priljubljen gradbeni kompozitni material (CM). Njegove dolgoletne lastnosti so odvisne od me{anice, polo`itve
in strjevanja. Ko je beton izdelan in ko ga polo`imo na mesto, kjer naj bi bil do konca trajnostne dobe, je zelo ob~utljiv in se
lahko po{koduje. Strjevanje je ena od stvari, ki zahteva za{~ito betona v prvih tednih njegovega trajanja s pravilno temperaturo
(niti pretoplo niti prehladno) in vla`nostjo. Metoda akusti~ne emisije in meritve elektri~nih lastnosti so bile uporabljene za
kontrolo mladega betona. Preiskovana je bila odvisnost med aktivnostjo akusti~ne emisije in temperaturo, elektri~nimi last-
nostmi, kot sta upornost in kapacitivnost mladega betona. W. Chen s sodelavci je {tudiral razvoj mikrostrukture mlade hidra-
tantne cementne zmesi z neporu{nimi metodami, vklju~no z merjenjem hitrosti napredovanja ultrazvo~nih P-valov, in naredil
preizkuse brezkontaktne elektri~ne upornosti. Vseeno ni veliko literaturnih virov o kontinuirnem spremljanju lastnosti mladega
betona. Za dolo~itev zdr`ljivosti in kvalitete betona je potrebna dolgotrajna kontrola. Metoda akusti~ne emisije se je izkazala kot
napredno orodje za neporu{no kontrolo mikrosprememb zgradbe med trajnostno dobo betona. Razlike, ugotovljene z akusti~no
emisijo med me{anicami betona pri strjevanju, so delno dolo~ile njihove lastnosti po 28 d. Osnovna lastnost betona je 28-dnevna
tla~na trdnost; vendar se ta lahko spreminja v dalj{em ~asovnem obdobju, zato se uporabljajo kontinuirne meritve.
Klju~ne besede: beton, akusti~na emisija, elektri~ne lastnosti, zgodnja starost, trajnostna doba, strjevanje
1 INTRODUCTION
Cement hydration plays a critical role in the
temperature development of early-age concrete due to
the heat generation1–4. It also has significant effects on
the material properties and performances at an early age,
including material strengths5, critical stresses6 and dis-
tresses like cracking7. Therefore, it is essential to capture
the hydration property and temperature development of
concrete at an early age to prevent premature failures. It
is known that the durability of cement-based products
and concrete structures is highly influenced by the early
stages of hydration8. A precise knowledge of the micro-
mechanical properties during the successive phases of
the hydration process provides information on the
concrete resistance and allows an assessment of its
durability9. Several non-destructive techniques have been
developed and applied in that respect, most of them
based on ultrasonic-wave measurements10. The recording
of passive energy using acoustic-emission techniques
was used to evaluate the structural activity in concrete
during the early ages, showing periods of intense micro-
structural changes during the curing process11.
Microstructural changes occurring in freshly poured
concrete during the curing were monitored on a labora-
tory scale using a combination of the acoustic-emission
technique and the monitoring of electrical properties12.
The acoustic-emission method is a passive-ultrasonic-
Materiali in tehnologije / Materials and technology 49 (2015) 5, 703–707 703
UDK 691:620.1:691.3:620.179.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)703(2015)
signal recording technique allowing an online monitor-
ing of the internal microstructural activity of young
concrete during the hydration process8. The acoustic-
emission technique, which is one of the non-destructive
evaluation techniques, is a useful method for investi-
gating local damage in concrete during its lifetime13. The
micro-changes in concrete structures can be easily esti-
mated, taking account of the number of acoustic-emis-
sion hits and counts of the acoustic-emission signal14.
Acoustic-emission signals are detected, due to micro-
cracks (i.e., structural changes in the material) by the
acoustic-emission sensors attached to the surface of a
concrete specimen. The acoustic-emission parameters
such as the number of hits, counts, duration time, ampli-
tude, energy and rise time are recorded by the acoustic-
emission measurement system15.
Impedance spectroscopy, called also dielectric spec-
troscopy, the monitoring of electrical properties and
nuclear-magnetic-resonance spectroscopy have become
promising techniques for probing cements and con-
cretes16. Cement paste is electrically conductive due to
its interconnected pore network filled with an aqueous
phase containing mobile ions such as Na+, K+ and OH–.17
Electrical conductivity and resistivity are effective
parameters used to monitor the microstructural features
of cement-based materials. The electrical conductivity of
cement-based materials is dependent on the porosity, the
pore connectivity and the conductivity of the pore solu-
tion, all of which are significant for the durability of the
material18. The impedance characteristics and spectra
recorded over a wide range of frequencies (if possible
from 100 mHz to 10 MHz or higher) provide new infor-
mation and an insight into the concrete microstructure
and hydration19. Electrical resistivity is related to the
reinforcement corrosion, moisture and heat transfer in
concrete20.
2 EXPERIMENTAL SET-UP
An acoustic-emission measuring system called
LOCation ANalyser (LOCAN) 320 made by PAC
(France) was used for the measurements. Wide-band
sensors (made by the 3S Sedlak Company) were used.
The sensor output was connected to pre-amplifiers (the
PAC Company) with a 2 kHz high-pass filter. The pro-
prietary PAC program package was used to record the
acoustic-emission parameters. The acoustic-emission
signals were taken by a LOCAN 320 acoustic-emission
localizer. The LOCAN 320 system was fine-tuned at a
well site and with a sonic-logging tool in the wellbore,
with the acoustic threshold set slightly above the back-
ground-noise level as measured by LOCAN 320 and the
internal gain set according to the manufacturer’s recom-
mendations. This device was designed to record and
evaluate ultrasonic signals. It was used to analyze and
store the acoustic-emission parameters. It can process
signals coming from four measuring points. Each
recorded event can be characterized with the following
acoustic-emission parameters: recording date and time,
(maximum) amplitude, acoustic-emission energy, count,
rise time and mean frequency21.
The impedance of the material under investigation
changes in accordance with the structural changes, parti-
cularly the water absorption and evaporation. A change
in the resistance is obvious. The capacity C of a
parallel-plate capacitor is computed with:
C f
S
dr
( ) = ⋅ ⋅ 0 (1)
where r is the relative permittivity, 0 is the vacuum
permittivity (8.854 · 10–12 F/m), S is the measuring-elec-
trode area, d is the distance and f is the frequency. The
resistance, R, of the electrodes is given by:
R f
d
S
( ) = ⋅
 (2)
where 
 is the resistivity. Microstructural changes in a
material make the material’s permittivity change. The
permittivity value can also be affected by macrocracks,
depending on the frequency22.
Two cylindrical steel electrodes of a diameter of
6 mm, buried 65 mm under the specimen surface, were
used for measuring the resistance. To measure the capa-
citance, two rectangular metal-plate electrodes with
dimensions of 25 mm × 45 mm were applied. All the
electrodes were fixed to a plastic slab so that their con-
stant configuration was guaranteed. The electrodes and
the temperature-sensor outputs were connected to an
automated measuring device. The measurements of the
capacitance, the temperature and the impedance were
started within 15 min after the mixture preparation. This
phase of the experiment was carried out using a TESLA
BM 595 RLCG bridge and a selector switch. The measu-
rement was carried out at selected points with frequen-
cies in the range from 100 Hz to 20 kHz. Each of the
electrical quantities (resistance, capacitance, tempera-
ture, etc.) was measured separately23.
L. PAZDERA et al.: MONITORING EARLY-AGE CONCRETE WITH THE ACOUSTIC-EMISSION METHOD ...
704 Materiali in tehnologije / Materials and technology 49 (2015) 5, 703–707
Figure 1: Experimental set-up before unmoulding, (samples S1, S2
tested with the acoustic-emission method and S3, S4 with impedance
spectroscopy)
Slika 1: Eksperimentalni sestav pred razdrtjem (vzorca S1, S2 preiz-
ku{ena z metodo akusti~ne emisije in S3, S4 z impedan~no spektro-
skopijo)
Sleeper concrete with a 32 mm crushed aggregate
attained a compression strength of 107 MPa. Two of the
specimens with dimensions of 100 mm × 100 mm × 400
mm were mould cast. Steel-mould upper sides were
covered with PE foil (Figure 1). The components of the
mixture are shown in Table 1. The mechanical properties
of the monitored samples are statistically processed in
Table 2. When the specimens were unmoulded, 24 h
after the mixing, one specimen was covered with foil and
the other one was not (Figure 2). In Figures 1 and 2,
samples S1 and S2 are tested with the acoustic-emission
method and samples S3 and S4 with impedance
spectroscopy; samples S1 and S4 are covered.
Table 1: Concrete mixture
Tabela 1: Me{anica betona
Composition Mass, m/kg
Cement CEM I 42.5R 390
Water 125
Plasticizer Stachement 2060 4
Sand 0/4 700
Coarse aggregate 16/32 1275
3 RESULTS
The acoustic-emission activities of both samples
during hydration, hardening and setting are shown in
Figures 3 and 4. After the unmoulding, i.e., at the age of
24 h, when one sample was wrapped and the other one
was left without any protection, the acoustic-emission
activities were different. The increase in the cumulative
acoustic-emission-count curve of the protected sample is
lower than that of the sample without protection during
the whole 120 d period, as shown in Figure 4. It should
be noted that the acoustic-emission signals were picked
up six hours after the mixtures were made.
The temperature of the samples after the unmoulding
was kept close to the ambient temperature (Figure 3).
The basic changes in the concrete structure during the
hydration relate to the temperature peak, in this case
after 12 h, as shown in Figure 4.24 The temperature
curves for both samples are the same as before the un-
moulding, i.e., slightly different25.
Electrical measurements show a higher resistivity and
a lower capacity, i.e., a higher capacitance of the speci-
men without protection after the unmoulding (Figure 5).
Different changes of both samples are also shown in
Figure 6. On the first day the frequency characteristics
of resistance (R) and capacitance (XC) are the same, but
thereafter they are different. There are three types of
curves described in the legend in Figure 6. Each curve,
as described on the curve at the 6th h, consists of eight
points whose frequencies from the left-hand side of the
horizontal axis are 100 Hz, 200 Hz, 400 Hz, 1 kHz,
L. PAZDERA et al.: MONITORING EARLY-AGE CONCRETE WITH THE ACOUSTIC-EMISSION METHOD ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 703–707 705
Figure 3: Dependence of cumulative acoustic emission (NC) and tem-
perature difference () on time (t)
Slika 3: Odvisnost kumulativne akusti~ne emisije (NC) in tempera-
turne razlike () od ~asa (t)
Figure 2: Experimental set-up after unmoulding (sample S1 and S4
are covered, S2 and S3 are uncovered)
Slika 2: Eksperimentalni sestav po razdrtju (vzorca S1 in S4 sta
pokrita in S2 ter S3 nepokrita)
Figure 4: Details of the curves plotted in Figure 3 for the first two
days
Slika 4: Detajl krivulje s slike 3 za prva dva dneva
2 kHz, 4 kHz, 10 kHz and 20 kHz 26. It is clear that the
shapes of the curves are different after the unmoulding,
therefore, the curves of both samples are the same on the
first day (in Figure 6 marked with downward-pointing
triangles).
Table 2: Main parameters of samples
Tabela 2: Glavni parametri vzorcev
Units Mean Deviation
Compressive strength MPa 107 5
Flexural strength MPa 10.0 0.2
Modulus of elasticity
(bending) GPa 37 3
Fracture toughness MPa m1/2 1.8 0.4
Fracture energy J/m2 200 10
4 CONCLUSION
The acoustic-emission method and impedance spec-
troscopy are powerful tools for determining the develop-
ment of microcracks during the hardening and setting of
concrete. It can be assumed that higher numbers of
microcracks cause higher numbers of acoustic-emission
events. The dependence of a cumulative acoustic-emis-
sion activity on the time shows the necessity of curing
the concrete during its whole lifetime, especially during
intense hydration and hardening. The big differences
between the acoustic-emission activities of wrapped and
unprotected samples clearly indicate the essentiality of
the high moistness of hardened concrete. The number of
microcracks in cement mortar affects the resulting me-
chanical properties of concrete.
Impedance, or dielectric, spectroscopy together with
acoustic emission can also help with the monitoring and
a detailed description of lifetime behaviours of concrete
specimens during hydration.
As the acoustic emission (Figure 3) and the impe-
dance (Figure 5) history of both samples are similar, the
number of microcracks is higher in the samples without
foil, which means that the curing of concrete during an
early age is necessary.
Acknowledgement
This paper has been worked out under the project No.
LO1408 "AdMaS UP – Advanced Materials, Structures
and Technologies", supported by Ministry of Education,
Youth and Sports under the "National Sustainability Pro-
gramme I" and under the project No. 13-18870S
"Assessment and Prediction of the Concrete Cover
Layers Durability" supported by Czech Science
Foundation.
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L. PAZDERA et al.: MONITORING EARLY-AGE CONCRETE WITH THE ACOUSTIC-EMISSION METHOD ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 703–707 707

P. ROVNANÍK, Á. DUFKA: EFFECT OF THE AGGREGATE TYPE ON THE PROPERTIES OF ALKALI-ACTIVATED ...
709–713
EFFECT OF THE AGGREGATE TYPE ON THE PROPERTIES OF
ALKALI-ACTIVATED SLAG SUBJECTED TO HIGH
TEMPERATURES
VPLIV VRSTE AGREGATA NA LASTNOSTI Z ALKALIJO
AKTIVIRANE @LINDRE, IZPOSTAVLJENE VISOKIM
TEMPERATURAM
Pavel Rovnaník, Ámos Dufka
Faculty of Civil Engineering, Brno University of Technology, Veveøí 95, 602 00 Brno, Czech Republic
rovnanik.p@fce.vutbr.cz
Prejem rokopisa – received: 2014-07-24; sprejem za objavo – accepted for publication: 2014-12-08
doi:10.17222/mit.2014.116
High temperatures present a risk of destruction for most silicate-based construction materials. Although these materials are not
flammable, they lose their properties due to a thermal decomposition. In contrast to ordinary Portland-cement-based materials,
alkali-activated slag exhibits a better thermal stability when exposed to temperatures up to 1200 °C. Due to its different porosity
it is less susceptible to spalling. However, the properties of the composites after a high-temperature treatment depend also on the
stability of the aggregates. The effects of two different types of the aggregate (quartz and chamotte) on the residual mechanical
properties and microstructure of the alkali-activated slag mortars exposed to 200–1200 °C is presented in this study. The results
showed an improved mechanical performance of the thermally stable chamotte aggregate at temperatures above 600 °C.
-quartz is transformed to -quartz at 573 °C, causing a volume instability and, consequently, a strength deterioration. Although
chamotte also contains some quartz phase, the reaction of mullite with the alkalis from the matrix leads to the formation of
albite and anorthite, making the material tougher and, thus, compensating the negative effect of quartz.
Keywords: alkali-activated slag, high temperatures, chamotte, quartz, strength, microstructure
Visoke temperature so tveganje za razpad ve~ine gradbenih materialov, ki temeljijo na silikatih. ^eprav ti materiali niso vnet-
ljivi, izgubijo svoje lastnosti zaradi termi~ne razgradnje. Nasprotno od materialov, ki temeljijo na portland cementu, z alkalijo
aktivirane `lindre izkazujejo bolj{o temperaturno obstojnost, ~e so izpostavljene temperaturam do 1200 °C. Zaradi razli~ne
poroznosti so manj ob~utljive za pojav lu{~enja. Vendar pa so lastnosti kompozita po visokotemperaturni obdelavi odvisne tudi
od stabilnosti agregatov. V {tudiji so predstavljeni vplivi dveh razli~nih vrst agregatov (kremen in {amota) na preostale
mehanske lastnosti in mikrostrukturo malte iz z alkalijo aktivirane `lindre, izpostavljene temperaturam 200–1200 °C. Rezultati
so pokazali izbolj{ane mehanske lastnosti toplotno stabilnega {amotnega agregata pri temperaturah nad 600 °C. -kremen se
transformira pri 573 °C v -kremen, kar povzro~i volumensko nestabilnost in posledi~no poslab{anje trdnosti. ^eprav {amot
vsebuje nekaj kremenove faze, reagiranje mulita z alkalijami iz osnove povzro~i nastanek albita in anortita, ki povzro~ita, da
material postane bolj `ilav in s tem kompenzira negativni vpliv kremena.
Klju~ne besede: z alkalijo aktivirana `lindra, visoke temperature, {amot, kremen, trdnost, mikrostruktura
1 INTRODUCTION
Portland cement is today’s most commonly used con-
struction material. It owes its widespread utilization to
its good mechanical properties, a low cost and a gener-
ally good performance. Nonetheless, Portland-cement
concrete is not an eternal material and its performance
can be affected by several degradation processes, one of
which involves high temperatures. Although concrete is
not flammable, when exposed to elevated temperatures it
partially decomposes and loses its mechanical proper-
ties.1–3 The second mostly discussed problem about Port-
land-cement concrete is spalling. This effect is mainly
related to the build-up of the pore pressure in conse-
quence of the vaporization of the physically/chemically
bound water which results in tensile loading of the
microstructure of heated concrete.4–6 These two pheno-
mena can be partially eliminated by utilizing alkali-
activated slag as a binder instead of ordinary Portland
cement.
Alkali-activated slag (AAS) is a hydraulic binder
based on finely ground, granulated blast-furnace slag.
Since the slag itself has no hydraulic properties it must
be activated with strong alkaline solutions, among which
hydroxides, carbonates and especially silicates play the
main roles. In contrast to the composites based on Port-
land cement, alkali-activated slag shows a much better
performance when exposed to very high temperatures.7–9
Similar to the case of Portland cement, a partial dehy-
dration and decomposition of the binder can be observed
up to 600 °C. The principle changes in the micro-
structure of alkali-activated slag occur between 600 °C
and 800 °C, when the dehydration of the C-A-S-H phase
is complete and new phases start to crystallize, among
which akermanite is dominant. Such a significant phase
transformation is reflected in the morphology, pore
distribution and, especially, in the improvement of the
mechanical properties of the alkali-activated slag.8 How-
ever, the properties of the composite materials with an
Materiali in tehnologije / Materials and technology 49 (2015) 5, 709–713 709
UDK 691:620.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)709(2015)
AAS matrix after a high-temperature treatment also
depend on the stability of the aggregates. The effects of
two different types of the aggregate (quartz and cha-
motte) on the residual mechanical properties and micro-
structure of the alkali-activated slag mortars exposed to
200–1200 °C is presented are this study.
2 EXPERIMENTAL PART
2.1 Materials
The tested material was prepared with the alkali acti-
vation of granulated blast-furnace slag. The slag supplied
by Kotou~, s. r. o. (CZ) was ground to a fineness of about
380 m2/kg (Blaine). It was a neutral slag with the basi-
city coefficient Mb = (CaO + MgO)/(SiO2 + Al2O3) equal
to 1.08, and its chemical composition in mass fractions
(w/%) was: SiO2 39.75, Al2O3 6.61, Fe2O3 0.46, CaO
39.03, MgO 10.45, Na2O 0.38, K2O 0.63, MnO 0.37,
SO3 0.71. A small amount of merwinite and a trace of
quartz were present as the only crystalline phases. Solid
sodium water glass having a SiO2/Na2O ratio of 1.95 was
used as an alkaline activator. Its chemical composition in
mass fractions (w/%) was: SiO2 50.75, Na2O 26.78, H2O
22.47.
To assess the influence of the aggregate on the sta-
bility of the AAS mortar at very high temperatures two
types of the aggregate were used. One of them was
quartz sand as a type of the commonly used standard ma-
terials and the second one was burnt clay, chamotte. The
main mineralogical phases found in the chamotte
aggregate with an XRD analysis were quartz, cristobalite
and mullite.
2.2 Sample preparation and heat treatment
Sodium silicate activator was suspended and partially
dissolved in water. Then, slag was added and the mixture
was stirred in a planetary mixer for about 3 min to pre-
pare homogeneous slurry. Finally, the aggregate was
added into the slurry and stirred to prepare fresh mortar.
The amount of the activator added was 20 % of the mass
of the slag, the aggregate/slag ratio was 3.0 and a water/
slag ratio of 0.4 was used to achieve an accurate con-
sistency.
The mixes were cast into prismatic moulds of the size
of 40 mm × 40 mm × 160 mm. After 24 h the hardened
specimens were immersed into a water bath at 20 °C for
another 27 d. After this period, they were allowed to air
dry for another 5 d before undergoing a high-temperature
treatment. The hardened mortars were heated in a Muffle
furnace to temperatures of (200, 400, 600, 800, 1000,
and 1200) °C at a constant heating rate of 5 °C min–1.
The specimens were kept at the given temperature for 1 h
and then allowed to cool down slowly to room tempera-
ture.
The heat-treated specimens were tested for their resi-
dual mechanical properties, which were compared with
those obtained for the unheated AAS mortar at the age of
28 d. Flexural strengths were determined using the stan-
dard three-point-bending test and compressive strengths
were measured on the far edge of each of the two
residual pieces obtained from the flexural test according
to the EN 196-1 standard.
The pore distribution was evaluated by means of a
mercury intrusion porosimetry analysis, conducted on
the samples using a Micromeritics Poresizer 9300 poro-
simeter that can generate the maximum pressure of 207
MPa and can evaluate a theoretical pore diameter of
0.006 μm. Micrographs of the alkali-activated slag mor-
tars were taken with a TESCAN MIRA3 XMU scanning
electron microscope in the SE mode. The experiments
were carried out on dry samples that were sputtered with
gold and an acceleration voltage of 25 kV was used.
X-ray diffraction analyses were carried out using a
Bruker D8 advance system equipped with a Cu tube
(	(K) = 0.154184 nm). The instrument is also equipped
with an Anton Paar HTK 16 temperature attachment
facilitating the measurements at up to 1600 °C.
3 RESULTS AND DISCUSSION
The mechanical properties and microstructures of the
AAS mortars subjected to high temperatures were
compared with the samples, treated only at ambient tem-
perature, marked as šRef’ (Table 1). The compressive
P. ROVNANÍK, Á. DUFKA: EFFECT OF THE AGGREGATE TYPE ON THE PROPERTIES OF ALKALI-ACTIVATED ...
710 Materiali in tehnologije / Materials and technology 49 (2015) 5, 709–713
Figure 1: Comparison of the relative residual compressive strengths
of AAS mortars after high-temperature treatment
Slika 1: Primerjava relativnih preostalih tla~nih trdnosti AAS-malt po
obdelavi na visoki temperaturi
Table 1: Compressive and flexural strengths of AAS mortars with
different aggregate types after exposure to high temperatures
Tabela 1: Tla~na in upogibna trdnost AAS-malt z razli~no vrsto
agregata po izpostavitvi visokim temperaturam
Tempera-
ture (°C)
Compressive strength
(MPa) Flexural Strength (MPa)
Quartz Chamotte Quartz Chamotte
Ref 73 92 9.6 10.4
200 55 77 4.4 7.5
400 46 74 3.8 6.2
600 31 42 2.6 4.8
800 10 31 1.7 4.2
1000 16 29 1.0 5.0
1200 26 49 2.9 12.5
strengths of the AAS mortars with chamotte reached 92
MPa which is very close to the value for the AAS matrix
itself;8 therefore, the aggregate seems not to be a limiting
factor for the mechanical properties of such composites.
On the contrary, quartz sand caused a strength decrease
by 19 MPa. This might be explained with weaker con-
tacts between the AAS matrix and the aggregate. After
the exposure to an elevated temperature, the compressive
strengths gradually decreased (Figure 1). The trend,
however, appeared to be the same for both composites;
therefore, it can be attributed to the dehydration and de-
composition of the AAS matrix. The strength decrease of
the AAS composite with quartz above 600 °C was partly
caused by the phase transition of quartz at 573 °C
(-quartz → -quartz), accompanied by a change in the
volume. The exposure of the AAS material to tempera-
tures above 800 °C caused an increase in the strength
that reached 53 % of the original strength for the AAS
with the chamotte aggregate and 36 % of that obtained
for the AAS with quartz.
The flexural strengths of both AAS composites
treated at an ambient temperature were quite similar.
After a high-temperature treatment the strengths de-
creased due to the decomposition processes in the
matrix; however, in the case of the quartz aggregate the
strength deterioration was more pronounced (Figure 2).
The minimum value for the chamotte composite was
achieved at 800 °C when the strength dropped to 40 % of
the reference value, while the mortar with quartz exhi-
bited the minimum strength at 1000 °C reaching only
10 % of the original value. An enormous difference in
the mechanical performance was observed after the ex-
posure to 1200 °C. The flexural strength of the chamotte
composite exceeded the reference value by 20 %, while
the strength achieved for the quartz sample remained
very low.
The pore distribution obtained by means of mercury
intrusion porosimetry up to 400 °C resembled the poro-
sity of the matrix itself.8 The main difference in the pore
distribution appeared after the exposure to 600 °C (Fig-
ure 3). There was an increase in the pore volume in the
range of large capillary pores (1–10 μm) for the quartz
composite, which can be clearly attributed to the rever-
sible phase transition to high quartz, followed by an
increase in the aggregate volume by approximately
4.5 %.10 After the exposure to 1200 °C, the major part of
the porosity can be attributed to the pores > 10 μm. This
is mainly caused by the shrinkage of the matrix during
the crystallization of the new phases. However, the total
porosity of the chamotte composite was much lower and
it also corresponded to the differences in the achieved
mechanical parameters.
P. ROVNANÍK, Á. DUFKA: EFFECT OF THE AGGREGATE TYPE ON THE PROPERTIES OF ALKALI-ACTIVATED ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 709–713 711
Figure 4: High temperature XRD of AAS mortar with chamotte
aggregate (pg – plagioclase)
Slika 4: Visokotemperaturna XRD in AAS-malta s {amotnim agrega-
tom (pg – plagioklas)
Figure 3: Differences in pore distribution for AAS mortars with
quartz and chamotte after exposure to 600 °C and 1200 °C, respec-
tively
Slika 3: Razlike v razporeditvi por v AAS-`lindrah s kremenom in
{amotom po izpostavitvi na temperaturo 600 °C in 1200 °C
Figure 2: Comparison of the relative residual flexural strengths of
AAS mortars after high temperature treatment
Slika 2: Primerjava relativne preostale upogibne trdnosti AAS-`linder
po obdelavi na visoki temperaturi
The changes in the mineralogical phases of the AAS
mortars were observed by means of a high-temperature
XRD analysis. Since quartz is quite an inert phase that
does not react with the AAS matrix even at a very high
temperature, there was practically no difference between
the XRD patterns of the quartz composite and the AAS
matrix itself.8 On the other hand, chamotte contains two
main crystalline phases, quartz and mullite. The latter is
not stable when such a high concentration of alkalis is
present in the material and reacts with the AAS matrix to
form two plagioclase phases, albite and anothite, at tem-
peratures exceeding 1000 °C (Figure 4). These two
phases were not observed in the XRD pattern of the pure
AAS matrix giving the proof of such an explanation.
Since the melting point of albite is in the range of
1100–1120 °C,11 it caused a partial fusing of the contact
zone between the aggregate grains and the matrix during
the heating to 1200 °C.
The microstructure and morphology of the AAS mor-
tars were investigated with scanning electron micro-
scopy. To explain the difference in the mechanical
behaviour of various AAS composites we focused on the
interstitial zone between the aggregate grains and the
matrix. Figure 5 shows the microstructure of the AAS
mortar with the quartz aggregate. The connection bet-
ween the quartz grains and the AAS matrix is very weak
and it is clearly evident that a small gap is present bet-
ween both phases as a result of the volume changes in
quartz at 573 °C. Therefore, the quartz grains can be
easily pulled out of the matrix, which predominantly
causes a deterioration of the flexural strength. Contrary
to quartz, the surfaces of the chamotte grains were sin-
tered with the matrix after the exposure to very high tem-
peratures (Figure 6). This finally resulted in the forma-
tion of ceramic bonds which were so strong that a
fracture of the aggregate grains occurred during the
mechanical testing.
4 CONCLUSIONS
This study aimed to estimate the influence of the
aggregate type on the behaviour of the alkali-activated
slag mortars subjected to very high temperatures. The
alkali-activated slag materials exhibit a much better per-
formance compared to the concrete made from ordinary
Portland cement when subjected to very high tempera-
tures or fire. A utilization of chamotte as an aggregate in
the AAS composites brought several benefits compared
to quartz sand. Chamotte is a material that is produced
from natural clay at 1350 °C and, therefore, it is more
thermally stable than pure quartz that undergoes a phase
transition at 573 °C. From the viewpoint of the mecha-
nical properties, the chamotte aggregate did not limit the
mechanical parameters of the AAS paste at temperatures
of up to 1000 °C and it even considerably improved the
flexural strength of the composite at 1200 °C due to the
sintering of the aggregate surface layer with the AAS
matrix. Although a degradation of the material occurred
upon heating, the compressive strength of the chamotte
composite was higher by 15 % and the flexural strength
was even four times higher compared to the quartz com-
posite after the exposure to 1200 °C. Therefore, the pro-
perties of the chamotte aggregate predestine its applica-
tion in the alkali-activated slag composites with an
improved high-temperature resistance.
Acknowledgement
This outcome was achieved with the financial support
of the Czech Science Foundation, project GA CR
14-25504S and the Ministry of Education, Youth and
P. ROVNANÍK, Á. DUFKA: EFFECT OF THE AGGREGATE TYPE ON THE PROPERTIES OF ALKALI-ACTIVATED ...
712 Materiali in tehnologije / Materials and technology 49 (2015) 5, 709–713
Figure 5: Microstructure of AAS mortar with quartz aggregate after
exposure to 1200 °C (SEM)
Slika 5: Mikrostruktura AAS-malt s kremenovim agregatom po
izpostavitvi na 1200 °C (SEM)
Figure 6: Microstructure of AAS mortar with chamotte aggregate
after exposure to 1200 °C (SEM)
Slika 6: Mikrostruktura AAS-malt s {amotnim agregatom po izposta-
vitvi na 1200 °C (SEM)
Sports of the Czech Republic under the "National Sus-
tainability Programme I" (project No. LO1408 AdMaS
UP), as an activity of the regional Centre AdMaS,
"Advanced Materials, Structures and Technologies". The
authors would like to thank Dr Patrik Bayer for
measuring some of the microstructural data.
5 REFERENCES
1 A. F. Bingöl, R. Gül, Fire Mater., 33 (2009), 79–88, doi:10.1002/
fam.987
2 G. F. Peng, Z. S. Huang, Constr. Build. Mater., 22 (2008) 4,
593–599, doi:10.1016/j.conbuildmat.2006.11.002
3 K. Y. Kim, T. S. Zun, K. P. Park, Cem. Concr. Res., 50 (2013),
34–40, doi:10.1016/j.cemconres.2013.03.020
4 P. Kalifa, F. D. Menneteau, D. Quenard, Cem. Concr. Res., 30
(2000), 1915–1927, doi:10.1016/S0008-8846(00)00384-7
5 K. D. Hertz, Fire Safety J., 38 (2003) 2, 103–116, doi:10.1016/
S0379-7112(02)00051-6
6 M. Zeiml, D. Leithner, M. Lackner, H. A. Mang, Cem. Concr. Res.,
36 (2006) 5, 929–942, doi:10.1016/j.cemconres.2005.12.018
7 M. Guerrieri, J. G. Sanjayan, F. Collins, Fire Mater., 33 (2009) 1,
51–62, doi:10.1002/fam.983
8 P. Rovnaník, P. Bayer, P. Rovnaníková, Constr. Build. Mater., 47
(2013), 1479–1487, doi:10.1016/j.conbuildmat.2013.06.070
9 R. Zhao, J. G. Sanjayan, Mag. Concr. Res., 63 (2011) 3, 163–173,
doi:10.1680/macr.9.00110
10 U. Raz, S. Girsperger, A. B. Thompson, Schweiz. Mineral. Petrogr.
Mitt., 82 (2002), 561–574
11 J. P. Greenwood, P. C. Hess, J. Geophys. Res., 103 (1998) B12,
29815–29828, doi:10.1029/98JB02300
P. ROVNANÍK, Á. DUFKA: EFFECT OF THE AGGREGATE TYPE ON THE PROPERTIES OF ALKALI-ACTIVATED ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 709–713 713

A. GRAJCAR: MICROSTRUCTURE EVOLUTION OF ADVANCED HIGH-STRENGTH TRIP-AIDED BAINITIC STEEL
715–720
MICROSTRUCTURE EVOLUTION OF ADVANCED
HIGH-STRENGTH TRIP-AIDED BAINITIC STEEL
RAZVOJ MIKROSTRUKTURE NAPREDNEGA
VISOKOTRDNOSTNEGA BAINITNEGA JEKLA Z UPORABO TRIP
Adam Grajcar
Silesian University of Technology, Institute of Engineering Materials and Biomaterials, Konarskiego Street 18a, 44-100 Gliwice, Poland
adam.grajcar@polsl.pl
Prejem rokopisa – received: 2014-08-01; sprejem za objavo – accepted for publication: 2014-10-08
doi:10.17222/mit.2014.154
The present work is focused on monitoring the microstructure evolution of the thermomechanically rolled medium-C steel
containing Si and Al to prevent the precipitation of carbides. The initial microstructure consists of carbide-free bainite,
polygonal ferrite and blocky-type and interlath retained austenite. To monitor the transformation behaviour of the retained
austenite transforming into strain-induced martensite, tensile tests were interrupted at particular strains (5 %, 10 %, 15 %, 18.4
%). The amount of the retained austenite determined with XRD and EBSD techniques decreased from 17 % to 6.8 % upon
straining, which was related to the transformation-induced plasticity effect. The identification of the retained austenite and
strain-induced martensite was carried out using an electron scanning microscope equipped with EBSD (electron backscatter
diffraction). It was found that sheet samples are strongly strengthened by the transformation of the blocky retained austenite,
whereas the layers and films of the 
 phase are over-stabilized and they do not transform into martensite.
Keywords: AHSS, bainitic steel, retained austenite, thermomechanical treatment, TRIP effect, strain-induced martensite
Delo obravnava pregled razvoja mikrostrukture pri termomehanskem valjanju srednje oglji~nega jekla s Si in Al, ki prepre~ita
izlo~anje karbidov. Za~etna mikrostruktura je bila iz bainita brez karbidov, poligonalnega ferita ter zaostalega avstenita kockaste
oblike in zaostalega avstenita med lamelami. Za spremljanje vedenja zaostalega avstenita pri pretvorbi v deformacijsko
induciran martenzit so bili natezni preizkusi prekinjeni pri posameznih raztezkih (5 %, 10 %, 15 %, 18,4 %). Dele` zaostalega
avstenita, dolo~en z XRD in EBSD, se je zmanj{al po natezni deformaciji iz 17 % na 6,8 %, kar je posledica vpliva z
deformacijo povzro~ene pretvorbe na plasti~nost. Identifikacija zaostalega avstenita in napetostno induciranega martenzita je
bila izvr{ena z vrsti~nim elektronskim mikroskopom, opremljenim z EBSD (difrakcija odbitih elektronov). Ugotovljeno je, da
se vzorci plo~evine bolj utrjujejo s pretvorbo kockastega zaostalega avstenita, medtem ko so plasti in sloji 
-faze bolj
stabilizirani in se ne pretvorijo v martenzit.
Klju~ne besede: AHSS, bainitno jeklo, zaostali avstenit, termomehanska obdelava, u~inek TRIP, napetostno induciran martenzit
1 INTRODUCTION
Advanced high-strength steels (AHSS) for the auto-
motive industry cover a wide group of multiphase steels
showing a superior balance between the strength and
plasticity. Dual-phase (DP), complex-phase (CP) and
transformation-induced plasticity (TRIP) steels are
representatives of the AHSSs containing ferrite as the
main structural constituent.1–3 The strain-induced mar-
tensitic transformation of retained austenite during cold
forming, resulting in additional plasticity, stimulated
increased research activities to assess the microstruc-
ture-property relationships in TRIP-aided steels.2–4 The
demand for ultra-high-strength steels can be satisfied
when a soft ferritic matrix is replaced by harder bainite.
Hence, a growing importance of the AHSSs containing
increased bainite amounts can be observed in4–9. The
desired ductility can be only obtained in advanced car-
bide-free bainitic steels, achieved with Si, Al, or Si-Al
alloying. Unfortunately, Al-alloyed steels have a lower
solid-solution strengthening effect. Thus, increased car-
bon contents are used or microalloying is employed to
improve the strength levels with grain refinement and
precipitation strengthening, in a similar way as in the
HSLA steels.9–12 Recently, medium-Mn bainitic-auste-
nitic steels have also been developed.8,13 Medium-C or
medium-alloyed sheet steels are used for the products
capable of mechanical binding and chemical adhesion14
or they can be also used for forgings and wires.4,6
Bainite-based steels containing large volume frac-
tions of retained austenite can be produced with cold
rolling or hot rolling. The knowledge concerning the
relationships between the manufacturing conditions, the
microstructure and the mechanical properties of cold-
rolled and annealed products is sufficiently documented
in literature studies.5,6,10,14 However, there are only few
reports on the microstructure-property relationships in
hot-rolled medium-C, carbide-free bainitic steels.4,7 It is
known that thermomechanical processing requires a
detailed time-temperature regime during a sheet’s cool-
ing.13,15 It is especially important for multiphase steels
because austenite retention requires one or two isother-
mal holding steps upon cooling. Isothermal holding of
steel at a bainitic transformation temperature is obliga-
tory for the enrichment of austenite in carbon from baini-
tic ferrite. This schedule is used for bainitic-austenitic
Materiali in tehnologije / Materials and technology 49 (2015) 5, 715–720 715
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Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)715(2015)
alloys.9 An additional enrichment of austenite in C can
be provided with an isothermal holding step at the inter-
critical region5,15 or using low finishing rolling tempera-
tures.7
Mechanical properties of multiphase steels are de-
pendent on the volume fractions and properties of bainite
and retained austenite. The final fraction of retained
austenite that can be stabilized depends, in turn, on the
chemical composition of steel and heat-treatment condi-
tions. However, the most important is the mechanical
stability of retained austenite, dependent on its carbon
content, size, dislocation density, morphology (blocky,
interlath), etc.3–5,10 An efficient monitoring of the evo-
lution of retained austenite and strain-induced martensite
as a function of strain is crucial for understanding
microstructure-property relationships. This approach can
be used with typical TRIP-aided ferrite-based steels.1–3,16
The aim of the present study is to monitor the evolution
of the microstructure of TRIP-aided medium-C bainitic
steel.
2 EXPERIMENTAL PROCEDURE
The work was focused on monitoring the micro-
structure evolution of the Si-Al type TRIP-aided steel
containing 0.43 % C, 1.45 % Mn, 0.98 % Si, 1.0 % Al,
0.033 % Nb, 0.01 % Ti, 0.004 % S, 0.01 % P and 0.0028 %
N. The steel was produced with vacuum-induction
melting followed by casting under an Ar atmosphere, hot
forging and rough rolling to a thickness of 4.5 mm. The
chemical composition was designed to obtain a large
fraction of retained austenite and carbide-free bainite
using the complex Si+Al approach.2,3 Nb and Ti micro-
additions were used to increase the strength level
through grain refinement and precipitation strengthening.
The thermomechanical rolling was conducted in
3 passes between 1100 °C and 750 °C to the final thick-
ness of about 2 mm. The cooling schedule consisted of
the finishing rolling at 750 °C followed by slow cooling
of the sheet samples to a temperature of 600 °C within
25 s. Immerse cooling of the specimens at a rate of about
50 °C/s to an isothermal holding temperature of 450 °C
was the next step. The isothermal holding time was 600 s.
Finally, the samples were cooled at a rate of about 0.5
°C/s to room temperature. Then, the samples with a
gauge length of 50 mm and a width of 12.5 mm were
prepared for the tensile tests carried out at a strain rate of
5 × 10–3 s–1. To monitor the transformation of the
retained austenite into martensite as a function of the
strain, interrupted tensile tests were carried out. The
tensile tests were interrupted at different tensile strains in
steps of 5 % followed by EBSD investigations to deter-
mine the volume fraction of the remaining austenite at a
particular strain.
Metallographic specimens were taken along the roll-
ing and tensile directions. Scanning electron microscopy
(SEM) and transmission electron microscopy (TEM)
were used to reveal morphological details of the micro-
structure. Additionally, orientation imaging microscopy
(OIM) using SEM was applied. Metallographic observa-
tions were carried out with the SUPRA 25 SEM using
BSE electrons. Observations were performed on nital-
etched samples at an accelerating voltage of 20 kV. The
EBSD technique was carried out using the Inspect F
SEM equipped with Schottky field emission. After
classical grinding and polishing, the specimens were
polished with Al2O3 with a granularity of 0.1 μm. The
final step of the sample preparation was their ion polish-
ing using the GATAN 682 PECS system. Thin-foil inve-
stigations were carried out with JEOL JEM 3010 at an
accelerating voltage of 200 kV using a mixture of 490
mL H3PO4 + 7 mL H2SO4 + 50 g CrO3 as the electrolyte
for the sample preparation.
A determination of the retained austenite fraction and
its carbon content after the thermomechanical rolling
was performed using the X-ray phase analysis based on
the Rietveld method. X-ray tests were done using a
Philips PW 1140 diffractometer with Co radiation. The
positions of the maxima of the diffraction lines of the
austenitic phase were used to determine the lattice con-
stant of austenite (a
). This parameter is necessary to
calculate the carbon content in the retained austenite (C
)
using the following formula: a
 = 0.3578 + 0.0033C
.17
The fraction of the retained austenite at a particular
strain ((5, 10, 15, 18.4) %) was determined with EBSD
as the average value of three measurements. A magni-
fication lower than 3000-times was applied to obtain
reliable quantitative results.
3 RESULTS AND DISCUSSION
The initial microstructure of the steel after the ther-
momechanical rolling and isothermal holding at a tempe-
rature of 450 °C is presented in Figure 1. It consists of
bainite, fine-grained ferrite (F) and retained austenite
(RA). The structural constituents are characterized by a
significant refinement related to a large population of
nucleation sites in the 
 transformation after the
finishing rolling at a temperature of 750 °C. The largest
amount of retained austenite occurs between the laths of
bainitic ferrite (BF), while blocky grains of the 
 phase,
of up to about 4 μm, are located in the boundary regions
of the bainitic-austenitic (BA) islands. The interlath
retained austenite takes the form of continuous or
interrupted laths or small regions located in the bainitic
islands (Figure 1b). The average size of ferrite grains is
approximately 5 μm.
The amount of the retained austenite in the initial
state determined with an X-ray examination is 17.6 %.
This result is in good agreement with the EBSD mea-
surements, showing that the retained-austenite fraction
reaches 17.1 % (Table 1). The stabilization of the large
fraction of the retained austenite is due to its high carbon
content w(C
 = 1.64 %). The enrichment of austenite in
A. GRAJCAR: MICROSTRUCTURE EVOLUTION OF ADVANCED HIGH-STRENGTH TRIP-AIDED BAINITIC STEEL
716 Materiali in tehnologije / Materials and technology 49 (2015) 5, 715–720
carbon was possible due to its partitioning from polygo-
nal and bainitic ferrite during the slow cooling of the
sheet samples during the 
 transformation and sub-
sequent isothermal holding at 450 °C for 600 s. This C
content is higher when compared to the other TRIP-
aided ferrite-based steels w(C
  1.0–1.4 %).2,3,7,16,17 Due
to the very low finishing rolling temperature (750 °C)
almost all the austenite grains are pancaked (not
recrystallized) during the bainitic transformation. Thus,
both the blocky grains and the layers of the retained
austenite remain heavily deformed and exhibit a high
dislocation density.
At a strain of 5 % the amount of the retained auste-
nite decreases to about 13.1 % (Table 1), which corres-
ponds to 23 % of the initial amount of the 
 phase
transformed into martensite. The evidence of the
strain-induced martensitic transformation at the strain of
5 % is found on the micrographs in Figure 2. It can be
seen that the strain-induced martensite forms in the
largest grains of the retained austenite, usually located
on the boundaries of the bainitic-austenitic islands. On
the other hand, the laths of the retained austenite located
between the bainitic-ferrite laths are mechanically stable.
Table 1: Retained-austenite fractions after thermomechanical rolling
(0 % strain) and at particular strains determined with the EBSD
technique
Tabela 1: Dele` zaostalega avstenita, dolo~enega z EBSD-tehniko, po
termomehanskem valjanju (0 % raztezka) in pri dolo~enih raztezkih
0 % 5 % 10 % 15 % 18.4 %
17.1 ± 1.2 13.1 ± 1.4 11.2 ± 1.3 9.0 ± 1.0 6.8 ± 1.1
Figure 3 shows EBSD maps registered at a strain of
10 %. Comparing the image quality (IQ) map (Figure
3a) and phase map (Figure 3b) it is possible to identify
particular structural constituents. In the grey-scale IQ
map the retained austenite regions correspond to the dark
areas of a poor diffraction quality. The IQ factor repre-
sents a quantitative description of the sharpness of the
EBSD pattern. A lattice distorted by crystalline defects
such as dislocations and grain/subgrain boundaries has a
distorted Kikuchi pattern, leading to lower IQ values.1,2
A. GRAJCAR: MICROSTRUCTURE EVOLUTION OF ADVANCED HIGH-STRENGTH TRIP-AIDED BAINITIC STEEL
Materiali in tehnologije / Materials and technology 49 (2015) 5, 715–720 717
Figure 1: SEM micrographs of the: a) bainitic-ferritic steel containing
blocky and interlath retained austenite after thermomechanical rolling,
b) morphological details of the bainitic-austenitic (BA) islands; F –
ferrite, BA – bainitic-austenitic islands, RA – retained austenite, BF –
bainitic ferrite
Slika 1: SEM-posnetka bainitno-feritnega jekla, ki vsebuje: a) kockast
in medlamelni zaostali avstenit, po termomehanskem valjanju, b)
detajli morfologije bainitno-avstenitnih (BA) oto~kov; F – ferit, BA –
bainitno-avstenitni oto~ki, RA – zaostali avstenit, BF – bainitni ferit
Figure 2: SEM micrographs showing strain-induced martensite in
blocky grains of retained austenite at 5 % strain; magnifications of: a)
10000-times and b) 20000-times; F – ferrite, BA – bainitic-austenitic
islands, RA – retained austenite, BF – bainitic ferrite, M –
strain-induced martensite
Slika 2: SEM-posnetka, ki prikazujeta napetostno induciran martenzit
v kockastih zrnih zaostalega avstenita pri raztezku 5 %; pove~ava: a)
10000-kratna in b) 20000-kratna; F – ferit, BA – bainitno-avstenitni
oto~ki, RA – zaostali avstenit, BF – bainitni ferit, M – deformacijsko
induciran martenzit
Hence, the strain-induced martensite can be distin-
guished from the other BCC phases (bainitic ferrite and
polygonal ferrite) because it has the most defected struc-
ture and, thus, it is characterized by the poorest diffrac-
tion quality. Additionally, Wasilkowska et al.2, Zaefferer
el al.3 and Petrov et al.16 reported that bainitic ferrite and
martensite can be indentified if we know special
crystallographic orientations between these BCC phases
and retained austenite. It is known that the Kurdjumov-
Sachs (KS) orientation shows a misorientation angle of
43 °, whereas the Nishiyama-Wasserman (NW) orienta-
tion shows a 46 ° angle. During deformation they are
mixed with each other. The way to distinguish them is to
analyse the rotation-axis distributions. They are around
<0.18 0.18 0.97> and <0.2 0.08 0.98> for the KS and
NW, respectively.
Fortunately, the characteristic misorientation angles
fulfilling the Kurdjumov-Sachs and Nishiyama-Wasser-
man relationships are sufficiently different to be
distinguished using the EBSD mapping. Hence, Figure
3c shows the IQ map with the angles corresponding to
the KS and NW relationships with a deviation of ± 1 °. It
is clear that these special angles are mainly located
between the retained austenite (RA) and strain-induced
martensite (M). The strongly defected strain-induced
martensite occurs as dark regions in the phase map. The
transformation causes a gradual fragmentation of the
retained-austenite regions. Thus, bainitic-martensitic-
austenitic constituents (BMA) are formed.
A further increase in the strain to 15 % causes the
amount of the retained austenite to decrease to about 9 %
(Table 1). The martensitic transformation takes place in
large blocky grains of the retained austenite. The
initiation of the transformation at smaller blocky grains
located in the BA regions can be also observed (Figure
4). However, a high content of C in the retained austenite
results in the lack of transformation of the 
 phase layers.
According to the literature data2,3,7,9 the carbon content in
the blocky grains is smaller when compared to
retained-austenite films. The reason is a longer diffusion
path of carbon in the relatively large blocky grains. Thus,
the center of a grain has a smaller amount of carbon
compared to the external regions of the grain. The same
applies to the films of the retained austenite. The
diffusion path of C from the bainitic-ferrite laths into the
retained-austenite films is much shorter. Hence, these are
mechanically stable upon the straining. Their stability is
additionally increased by hard laths of bainitic ferrite,
which form a protecting barrier against strain distribution
into the retained austenite.4,9
The strengthening caused by the complex effects of
the thermomechanical processing, strongly defected
ferrite and bainite, strain-induced martensite and over-
stabilized films of the retained austenite, lead to a prema-
ture breaking of the samples at an elongation of 18.4 %
and UTS of 930 MPa. The volume fraction of the
retained austenite present after the sample breaking is
about 6.8 %. It means that about 40 % of the initial
amount of the retained austenite is untransformed. This
A. GRAJCAR: MICROSTRUCTURE EVOLUTION OF ADVANCED HIGH-STRENGTH TRIP-AIDED BAINITIC STEEL
718 Materiali in tehnologije / Materials and technology 49 (2015) 5, 715–720
Figure 3: a) Image quality (IQ) map with misorientation angles, b) phase map showing retained austenite – RA regions, and c) IQ map with
misorientation angles corresponding to the Kurdjumov-Sachs (42–44 °) and Nishiyama-Wasserman (45–47 °) relationships at 10 % strain; F –
ferrite, BMA – bainitic-martensitic-austenitic constituents, RA – retained austenite, BM – bainitic-martensitic constituents, M – strain-induced
martensite
Slika 3: a) Posnetek podro~ja kvalitete (IQ) z razporeditvijo slabo orientiranih zrn, b) na~rt faz prikazuje zaostali avstenit – RA podro~ja in c)
na~rt slabo orientiranih zrn po odvisnostih Kurdjumov-Sachs (42–44 °) in Nishiyama-Wasserman (45–47 °) pri raztezku 10 %; F – ferit, BMA –
bainitno-martenzitno-avstenitne sestavine, RA – zaostali avstenit, BM – bainitno-martenzitne sestavine, M – deformacijsko induciran martenzit
phase remains mainly in the form of thin layers and films
located between the hard laths of bainitic ferrite (Figure
5). Firstly, the high mechanical stability of the retained
austenite results from its high C concentration of 1.64 %.
Secondly, it is due to the fragmentation of the 
-phase
particles and the compressive stresses caused by the hard
A. GRAJCAR: MICROSTRUCTURE EVOLUTION OF ADVANCED HIGH-STRENGTH TRIP-AIDED BAINITIC STEEL
Materiali in tehnologije / Materials and technology 49 (2015) 5, 715–720 719
Figure 4: SEM micrographs showing strain-induced martensite in
small blocky grains of retained austenite located in BA regions at
15 % strain: a) magnifications of 10000-times and b) 20000-times; F –
ferrite, BA – bainitic-austenitic islands, RA – retained austenite, BF –
bainitic ferrite, M – strain-induced martensite
Slika 4: SEM-posnetka deformacijsko induciranega martenzita v
kockastih zrnih zaostalega avstenita v obmo~jih BA pri raztezku 15 %:
a) pove~ava 10000-kratna in b) 20000-kratna; F – ferit, BA – bainit-
no-avstenitni oto~ki, RA – zaostali avstenit, BF – bainitni ferit, M –
deformacijsko induciran martenzit
Figure 6: a) Image quality (IQ) map with misorientation angles, b)
phase map showing retained austenite – RA regions, and c) inverse
pole-figure map at rupture strain of 18.4 %; BF – bainitic ferrite, RA –
retained austenite, MA – martensitic-austenitic constituents, M –
martensite
Slika 6: a) Posnetek razporeditve kvalitete (IQ) s podro~ij slabo orien-
tiranih zrn, b) razporeditev faz prikazuje zaostali avstenit – RA pod-
ro~ja in c) razporeditev inverznih polovih figur pri raztezku 18,4 % ob
poru{itvi; BF – bainitni ferit, RA – zaostali avstenit, MA – marten-
zitno-avstenitne sstavine, M – martenzit
Figure 5: SEM micrograph showing strain-induced martensite in
blocky grains of retained austenite and stable interlath retained
austenite at rupture strain of 18.4 %; RA – retained austenite, BF –
bainitic ferrite, M – strain-induced martensite
Slika 5: SEM-posnetek deformacijsko induciranega martenzita v
kockastih zrnih zaostalega avstenita in stabilen zaostali avstenit med
lamelami pri raztezku 18,4 % ob poru{itvi; RA – zaostali avstenit, BF
– bainitni ferit, M – deformacijsko induciran martenzit
laths of bainitic ferrite. Finally, the stabilization of the 

phase is favoured due to a high density of the disloca-
tions in this phase remaining after the thermomechanical
processing.7
At the fracture strain, the IQ map (Figure 6a) no
longer coincides exactly with the inverse pole-figure map
(Figure 6c) showing the orientations of individual
grains. It is a result of a strongly defected multiphase
microstructure. The brightest areas in Figure 6 corres-
pond to the occurrence of defected ferrite and bainitic
ferrite, while dark areas of a poor diffraction quality can
be identified as the retained austenite and martensite.
The particles of the strain-induced martensite are the
reason of a substantial fragmentation of the retained-aus-
tenite regions. Thus, martensitic-austenitic (MA) regions
are formed (Figure 6b). A transmission electron micro-
graph reveals that the strain-induced martensite has a
fine-plate morphology (Figure 7). It is formed from the
retained austenite containing about the mass fraction
w(C) = 1.64 %. This high carbon concentration deter-
mines the plate morphology of the stain-induced marten-
site.
4 CONCLUSIONS
The thermomechanically processed, medium-C
bainitic steel contains a large volume fraction of retained
austenite (about 17 %), appearing as layers between the
bainitic-ferrite laths and blocky grains located in the
boundary regions of bainitic-austenitic islands. The mor-
phology of the retained austenite affects the transfor-
mation behaviour of this phase into martensite. The
blocky grains of the retained austenite gradually trans-
form into strain-induced martensite as a function of the
increasing strain. On the other hand, the interlath
retained austenite, more enriched in carbon than the
blocky grains, is stable upon the straining not providing
a TRIP effect. The mechanical stability of the retained
austenite is too high especially due to its high average
carbon content (w(C
 = 1.64 %). The high dislocation
density, the progressive fragmentation of the retained
austenite (its size) and the compressive stresses caused
by the bainitic-ferrite laths and strain-induced martensite
are equally important for the mechanical stabilization of
the retained austenite. It can be concluded that polygonal
ferrite and the subsequent partitioning of C during the

 transformation are not necessary for fine-grained,
thermomechanically processed, medium-C steel. Over-
stabilization of the retained-austenite layers should be
avoided by producing homogeneous bainitic-austenitic
microstructures with a lower C content in the 
 phase.
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A. GRAJCAR: MICROSTRUCTURE EVOLUTION OF ADVANCED HIGH-STRENGTH TRIP-AIDED BAINITIC STEEL
720 Materiali in tehnologije / Materials and technology 49 (2015) 5, 715–720
Figure 7: TEM micrograph showing the plate morphology of strain-
induced martensite at 18.4 % strain; BF – bainitic ferrite, RA –
retained austenite, M – strain-induced martensite
Slika 7: TEM-posnetek prikazuje morfologijo plo{~ deformacijsko
induciranega martenzita pri raztezku 18,4 %; BF – bainitni ferit, RA –
zaostali avstenit, M – deformacijsko induciran martenzit
T. YENER et al.: EFFECT OF ELECTRIC CURRENT ON THE PRODUCTION OF NiTi ...
721–724
EFFECT OF ELECTRIC CURRENT ON THE PRODUCTION OF
NiTi INTERMETALLICS VIA ELECTRIC-CURRENT-ACTIVATED
SINTERING
VPLIV ELEKTRI^NEGA TOKA PRI IZDELAVI INTERMETALNE
ZLITINE NiTi S SINTRANJEM, AKTIVIRANIM Z ELEKTRI^NIM
TOKOM
Tuba Yener1, Shafaqat Siddique2, Frank Walther2, Sakin Zeytin1
1Sakarya University, Engineering Faculty, Department of Metallurgy and Materials Engineering, Esentepe Campus, 54187 Adapazari, Sakarya,
Turkey
2TU Dortmund University, Faculty of Mechanical Engineering, Department of Materials Test Engineering (WPT), Baroper Str. 303,
44227 Dortmund, Germany
tcerezci@sakarya.edu.tr
Prejem rokopisa – received: 2014-08-01; sprejem za objavo – accepted for publication: 2014-10-09
doi:10.17222/mit.2014.161
This study focuses on investigating the fabrication of in-situ intermetallic NiTi composites from a powder mixture containing
the mass fractions 50 % nickel powder and 50 % titanium powder. The elemental powders were mixed in the stoichiometric
ratio corresponding to the NiTi intermetallic molar proportion of 1 : 1, ball-milled and uniaxially compressed under a pressure
of 170 MPa. Sintering was then carried out for 15 min in a steel mold using the electric-current-activated sintering method.
Electric-current values of 1000 A, 1300 A and 2000 A were used for the sintering while keeping the voltage in the range of 0.9
V to 1.2 V. The phases in the samples were analyzed with XRD and their Vickers hardness was measured as (701 ± 166) HV0.05.
Energy dispersive X-ray spectroscopy carried out with a scanning electron microscope (SEM-EDS) showed that the micro-
structures of the samples consist of different phases such as Ti, Ni2Ti3, NiTi2, Ni3Ti and TiO2 as a function of electric current.
The XRD analysis also supported the SEM-EDS results. The nano-indentation technique was used to determine the elastic
modulus of different phases.
Keywords: NiTi intermetallics, electric-current-activated sintering (ECAS), nano-indentation
Ta {tudija je usmerjena v preiskavo in situ izdelave intermetalnega NiTi-kompozita iz me{anice prahov z masnima dele`ema 50
% niklja v prahu in 50 % titana v prahu. Obe vrsti prahu sta bili zme{ani v stehiometri~nem razmerju, ki ustreza molskemu
razmerju NiTi 1 : 1, zmleti v kroglastem mlinu in enoosno stisnjeni pri tlaku 170 MPa. Petnajstminutno sintranje je bilo
izvr{eno v jeklenem kalupu s sintranjem, aktiviranim z elektri~nim tokom. Pri sintranju so bili uporabljeni elektri~ni tokovi
1000 A, 1300 A in 2000 A, medtem ko je bila napetost v obmo~ju med 0,9 V in 1,2 V. Faze v vzorcih so bile dolo~ene z
rentgenom (XRD) in izmerjena je bila trdota po Vickersu (701 ± 166) HV0,05. Energijska disperzijska rentgenska spektroskopija
(EDS), izvr{ena na vrsti~nem elektronskem mikroskopu (SEM-EDS), je pokazala, da je mikrostruktura vzorcev sestavljena iz
razli~nih faz, kot so Ti, Ni2Ti3, NiTi2, Ni3Ti in TiO2, odvisno od elektri~nega toka. XRD-analiza je podprla rezultate, dobljene s
SEM-EDS. Elasti~ni modul razli~nih faz je bil dolo~en z nanovtiski.
Klju~ne besede: intermetalna zlitina NiTi, sintranje, aktivirano z elektri~nim tokom (ECAS), nanovtiski
1 INTRODUCTION
To meet the increasing demand of the aerospace and
automobile industries for lightweight structural materials
suitable for high-temperature applications, researchers
focus on lightweight and high-strength intermetallics.1
Intermetallic compounds are promising materials for
structural and non-structural high-temperature applica-
tions (heat resistance, corrosion resistance, electronic
devices, magnets, super conductors).2,3 Especially NiTi
alloys are some of the most technologically important
shape-memory alloys4,5 which find extensive applications
in the mechanical, medical, electronic, chemical and
aerospace industries due to their excellent shape-memory
effect, high erosion resistance, high damping capacity
and biocompatibility. These properties make them suit-
able candidates for various engineering applications.2–7
Nickel titanium is a near-equiatomic intermetallic exhi-
biting distinctive and desirable thermo-mechanical
properties, namely, the thermal shape-memory effect and
super elasticity.8 Up to now, a number of processes in-
cluding self-propagating high-temperature synthesis
(SHS), thermal explosion, laser-melting deposition, cast-
ing and mechanical alloying techniques2,5,9,10 have been
used for manufacturing ýntermetallics. In the electric-
current-activated/assisted sintering (ECAS) technique, a
cold-formed compact obtained with uniaxial compres-
sion is inserted into a container heated by the passing
electrical current. A sintering pressure of 50 MPa is
applied and maintained throughout the sintering.11 The
present study aimed to determine the effect of the current
on the production of a NiTi intermetallic from the Ni and
Ti elemental powders.
Materiali in tehnologije / Materials and technology 49 (2015) 5, 721–724 721
UDK 621.762:621.762.5:669.018.29 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)721(2015)
2 EXPERIMENTAL PROCEDURE
2.1 Materials and methods
Powder materials from titanium (99.5 % purity,
35–44 μm) and nickel (99.8 % purity, 3–7 μm) were used
as the starting materials to manufacture a NiTi inter-
metallic compound. The Ni and Ti powders were mixed
in the stoichiometric ratio corresponding to the Ni-Ti
phase diagram (Figure 1), in a molar proportion of 1 : 1.
The powder mixture was cold-pressed before the
sintering to form a cylindrical compact in a metallic die
under a uniaxial pressure of 170 MPa. Dimensions of a
compact were 15 mm in diameter and 5 mm in thickness.
The production of the NiTi intermetallic compounds was
performed with the electric-current-activated sintering
technique in an open atmosphere at 1000–2000 A for 15
min as shown in Figure 2. The process parameters are
listed in Table 1.
Table 1: Process parameters for the samples
Tabela 1: Procesni parametri pri vzorcih
Sample code x/% Current (A) Voltage (V)
NT1 50Ti-50Ni 1.000 0.9–1.2
NT2 50Ti-50Ni 1.300 0.9–1.2
NT3 50Ti-50Ni 2.000 0.9–1.2
After the sintering, the specimens were unloaded and
air-cooled to room temperature. After metallographic
preparations, the resulting microstructures and phase
constitutions were characterized.
2.2 Characterization
The morphologies of the samples were examined
with scanning electron microscopy (SEM-EDS) in terms
of the resulting phases. An X-ray diffraction (XRD)
analysis was carried out using Cu-K radiation with a
wavelength of 0.15418 nm over a 2 range of 10–80 °.
The micro-hardness of the investigated samples was
measured employing the Vickers-indentation technique
with a load of 50 g using a Struers Duramin hardness
instrument. For determining the elastic modulus of the
samples, nano-indentation device DUH-211S by Shi-
madzu was used at a load of 98 mN and a dwell time of
10 s. The ImageJ programme was also used to determine
the porosity of fractions of the samples.
3 RESULTS AND DISCUSSION
3.1 SEM-EDS analysis
The morphologies of the as-received Ni and Ti pow-
ders are shown in Figure 3. The metallic Ni-powder
T. YENER et al.: EFFECT OF ELECTRIC CURRENT ON THE PRODUCTION OF NiTi ...
722 Materiali in tehnologije / Materials and technology 49 (2015) 5, 721–724
Figure 3: SEM micrographs of: a) Ni powder, b) Ti powder
Slika 3: SEM-posnetka prahov: a) Ni, b) Ti
Figure 2: Schematic of electric-current-assisted sintering (ECAS)
process
Slika 2: Shematski prikaz sintranja z elektri~nim tokom (ECAS)
Figure 1: Ni-Ti phase diagram3
Slika 1: Fazni diagram Ni-Ti 3
particles were generally spherical with a diameter of 4–7
μm. The Ti-powder particles had sharp corners and were
less than 40 μm in size.
SEM-EDS analyses of the NT1 and NT2 interme-
tallic compounds are shown in Figure 4. The micro-
structure in Figure 4a shows that a low current intensity
results in separately formed Ni and Ti areas. When
increasing the values of the current for NT2 (Figure 4b),
new phases like NiTi2 (Figure 1) start to form, but these
microstructures are still far from the stoichiometric
composition of the main NiTi phase.
When increasing the current to 2000 A for the NT3
sample (Figure 5), the main phases in the microstructure
such as NiTi, NiTi2, Ni3Ti are evident. Besides, there are
also a small amount of the residual Ti phase and a small
oxidation problem due to the open atmosphere. It can be
inferred that a high current intensity leads to a high
amount of intermetallic phases. According to the ImageJ
image-analysis programme, a porosity of 0.2 % can be
found in the microstructures.
As can be seen from the SEM and EDS analyses pre-
sented in Figure 6, the reaction of the NiTi-compound
synthesis was not completed. It is assumed that the
applied voltage was insufficient for a complete transfor-
mation of the NiTi phases during the sintering.
3.2 XRD analysis
The XRD analysis (Figure 7) shows that the main
phase of the composite is NiTi2. The NiTi phase is also
seen as a small peak. Additionally, Ni3Ti, Ti and TiO are
the other phases existing in the NT3 composite. These
results support the observations from the SEM-EDS
analysis.
3.3 Hardness and the elastic modulus
The measured HV0.05 hardness values for the NT1,
NT2 and NT3 samples are (250 ± 27), (405 ± 71) and
T. YENER et al.: EFFECT OF ELECTRIC CURRENT ON THE PRODUCTION OF NiTi ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 721–724 723
Figure 6: SEM-EDS analyses of NT3 composite
Slika 6: SEM-EDS-analize kompozita NT3
Figure 4: SEM-EDS analyses of: a) NT1, b) NT2 composites
Slika 4: SEM-EDS-analizi kompozitov: a) NT1, b) NT2
Figure 5: SEM-EDS analyses of NT3 composite
Slika 5: SEM-EDS-analizi kompozita NT3
(701.6 ± 166), respectively. The hardness values are in
good agreement with the literature.1,12 The elastic-modu-
lus values for the NT1, NT2 and NT3 samples deter-
mined with the nano-indentation technique are (98.3 ± 9)
GPa, (120 ± 23) GPa and (132 ± 25) GPa, respectively.
Nano-indentation is an important technique for probing
the mechanical behavior of materials at small-length
scales.13 The results for the elastic modulus cannot be
compared to the literature because of the contrasting re-
ported values.14 However, the obtained value of 135 GPa
is between the elastic-modulus values for the Ni and Ti
materials.
4 CONCLUSIONS
The reaction of the NiTi-compound synthesis was not
completed. It is inferred that the applied voltage was
insufficient for a complete transformation of the NiTi
phases. When increasing the current intensity from 1000
A to 2000 A, the fractions of intermetallic phases NiTi,
NiTi2 and Ni3Ti increased as well. There was also a small
amount of retained metallic Ti in the microstructure. The
hardness of the intermetallic composites was increased
from 250 HV to 701 HV by increasing the current inten-
sity, due to the formation of higher amounts of the inter-
metallic phases. The elastic modulus of the composites
was between the elastic-modulus values for the Ni and Ti
materials.
Acknowledgements
This paper was completed within the framework of
the PhD work of Mrs. Tuba YENER during a research
visit at the TU Dortmund University, the Department of
Materials Test Engineering (WPT) between 06–09/2014.
The host institution and the home institution, the Sakarya
University, the Department of Metallurgy and Materials
Engineering, gratefully acknowledge the excellent
scientific exchange and the financial support of the PhD
student exchange.
5 REFERENCES
1 X. Songa et al., Microstructure and mechanical properties of Nb- and
Mo-modified NiTi-Al-based intermetallics processed by isothermal
forging, Materials Science & Engineering A, 594 (2014), 229–234,
doi:10.1016/j.msea.2013.11.070
2 N. Ergin et al., An investigation on TiNi intermetallic produced by
electric current activated sintering, Acta Physica Polonica A, 123
(2013), 248–249, doi:10.12693/APhysPolA.123.248
3 M. Farvizi et al., Microstructural characterization of HIP consoli-
dated NiTi-nano Al2O3 composites, Journal of Alloys and Com-
pounds, 606 (2014), 21–26, doi:10.1016/j.jallcom.2014.03.184
4 F. Neves et al., Mechanically activated reactive forging synthesis
(MARFOS) of NiTi, Intermetallics, 16 (2008), 889–895,
doi:10.1016/j.intermet.2008.04.002
5 P. Novak et al., Effect of SHS conditions on microstructure of NiTi
shape memory alloy, Intermetallics, 42 (2013), 85–91, doi:10.1016/
j.intermet.2013.05.015
6 T. Yener et al., Synthesis and characterization of metallic-interme-
tallic Ti-TiAl3, Nb-Ti-TiAl3 composite produced by electric current
activated sintering (ECAS), Mater. Tehnol., 48 (2014) 6, 847–850
7 G. Chan et al., In-situ observation and neutron diffraction of NiTi
powder sintering, Acta Materialia, 67 (2014), 32–44, doi:10.1016/
j.actamat.2013.12.013
8 Z. Huan et al., Porous NiTi surfaces for biomedical applications,
Applied Surface Science, 258 (2012), 5244–5249, doi:10.1016/
j.apsusc.2012.02.002
9 M. Krasnowski et al., Al3Ni2-Al composites with nanocrystalline
intermetallic matrix produced by consolidation of milled powders,
Advanced Powder Technology, 25 (2014), 1362–1368, doi:10.1016/
j.apt.2014.03.018
10 A. Molladavoudi et al., The production of nanocrystalline cobalt tita-
nide intermetallic compound via mechanical alloying, Intermetallics,
29 (2012), 104–109, doi:10.1016/j.intermet.2012.05.012
11 R. Orrù et al., Consolidation/synthesis of materials by electric cur-
rent activated/assisted sintering, Materials Science and Engineering
R: Reports, 63 (2009), 127–287, doi:10.1016/j.mser.2008.09.003
12 F. Gao et al., Abrasive wear property of laser melting/deposited
Ti2Ni/TiNi intermetallic alloy, Transactions of Nonferrous Metals
Society of China, 17 (2007), 1358–1362, doi:10.1016/S1003-6326
(07)60277-5
13 X. Deng et al., Deformation behavior of (Cu, Ag)-Sn intermetallics
by nanoindentation, Acta Materialia, 52 (2004), 4291–4303,
doi:10.1016/j.actamat.2004.05.046
14 Y. Liu et al., Apparent modulus of elasticity of near-equiatomic NiTi,
Journal of Alloys and Compounds, 270 (1998), 154–159,
doi:10.1016/S0925-8388(98)00500-3
T. YENER et al.: EFFECT OF ELECTRIC CURRENT ON THE PRODUCTION OF NiTi ...
724 Materiali in tehnologije / Materials and technology 49 (2015) 5, 721–724
Figure 7: XRD analysis of NT3 sample
Slika 7: Rentgenogram vzorca NT3
J. STETINA et al.: CONTROL OF SOFT REDUCTION OF CONTINUOUS SLAB CASTING WITH A THERMAL MODEL
725–729
CONTROL OF SOFT REDUCTION OF CONTINUOUS SLAB
CASTING WITH A THERMAL MODEL
KONTROLA MEHKE REDUKCIJE PRI KONTINUIRNEM LITJU
SLABOV S TERMI^NIM MODELOM
Josef Stetina, Pavel Ramik, Jaroslav Katolicky
Brno University of Technology, Technicka 2, 616 69 Brno, Czech Republic
stetina@fme.vutbr.cz
Prejem rokopisa – received: 2014-08-09; sprejem za objavo – accepted for publication: 2014-10-17
doi:10.17222/mit.2014.189
The internal quality of cast steel slabs during the radial continuous casting is significantly affected by the setting of support
rollers. During the passes through the support rolls the steel cools from 1400 °C to 600 °C, therefore the shrinkage of the
material must be controlled with the setting of the reduction profile. The setting of the reduction profile with fixed rollers is a
compromise used for all the cast steel. During the continuous casting the wear of the rolls occurs and this must also be
considered when setting the profile with rollers. A 3D thermal model is used for the optimum setting. As the reduction setting
cannot be optimized for older casters, internal defects are often present. New machines for continuous casting of steel are fitted
with soft reduction, i.e., the system for controlling the roller position and setting the reduction profile for any type of steel. The
control system is connected on-line with the 3D thermal model.
Keywords: concast slab, numerical models, contraction, soft reduction
Notranja kakovost ulitih jeklenih slabov pri radialnem kontinuirnem litju je odvisna od nastavitev podpornih valjev. Ko jeklo
prehaja preko podpornih valjev, se ohlaja iz 1400 °C na 600 °C. Pri tem nastaja kr~enje materiala, kar je treba upo{tevati pri
nastavitvi profila. Nastavitev profila redukcije z valji je kompromis za vse ulito jeklo. Med kontinuirnim litjem nastaja tudi
obraba valjev, ki jo je treba upo{tevati pri nastavitvi profila z valji. Za optimalno nastavitev podpornih valjev je uporabljen 3D
termi~ni model. Nastavitve redukcije ni mogo~e optimirati pri starih livnih napravah, zato so pogoste notranje napake. Nove
naprave za kontinuirno litje jekla so opremljene z mehko redukcijo, to je s sistemom upravljanja kontrole pozicije valjev, ki
omogo~a nastavitev profila redukcije za vsako jeklo. Ta sistem kontrole je on-line povezan s 3D termi~nim modelom.
Klju~ne besede: concast slab, numeri~ni modeli, kr~enje, mehka redukcija
1 INTRODUCTION
Materials with high utility parameters are increas-
ingly in demand and the traditional production is being
replaced with higher-quality steels. More and more so-
phisticated aggregates using more sophisticated techno-
logical procedures are being implemented. To maintain
the competitiveness and diversification of the production
and to expand to the other markets, it is necessary to
monitor the technological development.
In the case of concasting, it is not possible to fulfil
these requirements without the application of the models
of caster processes depending on thermo-mechanical
relationships. These models can be applied both off-line
and on-line. For an off-line model the calculation time is
longer than the actual casting process. An on-line model
runs in real time and uses the data directly from the
operation – and the time used for the calculation is equal
or shorter than the time of the actual process.
These models support the designs of new machines
and the redesigns of old machines. They facilitate the
identification and quantification of any potential defect
and the optimization of various operational conditions
with the aim to increase the productivity and minimize
the occurrence of defects1.
Due to the complexity of the investigations of the
influences of all the involved factors, it is not possible to
develop a mathematical model that would cover all of
them. It is better to group the influences according to the
three main parameters, the heat and the mass transfer and
the mechanical behaviour.
The influences of the heat and the mass transfer are
the greatest because the temperature field gives rise to
mechanical and structural effects. The development of a
model of the temperature field (of a slab) with an inter-
face providing the data for the mechanical stress and
strain models is therefore the top-priority task.
The new caster for casting steel must be designed to
cast different types of steel, including special types, with
the maximum quality. Figure 1 shows a diagram of a
modern caster with a straight first section, followed by a
curvature. In segments 1 to 12 the displacement of the
whole structure and of individual rollers with a small
radius can be controlled hydraulically. Another trend in
the modern caster is a split of the secondary-cooling area
into several zones, in this case up to 17 individually con-
trolled cooling zones equipped with air-mist nozzles.
More information about the optimization of the secon-
dary cooling is found in2,3.
Materiali in tehnologije / Materials and technology 49 (2015) 5, 725–729 725
UDK 519.61/.64:621.74.047 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)725(2015)
In continuous casting of steel there are two problems
that affect the quality of the resulting slabs. The first is
the contraction of steel during its cooling in the caster; to
prevent internal shrinkage the rollers that are in contact
with the blank must be set to reduce the profiles of
individual segments. A fixed reduction is optimal only
for one temperature profile of one blank, but a change in
the casting speed makes the reduction larger or smaller.
The second problem is bulging, i.e., the bulge of a soli-
dified shell due to the hydrostatic pressure of the steel.
This bulge requires additional rollers and thus leads to
the alternating stress, decreasing the possibility of crack-
ing. The solution of both problems is soft reduction (Fig-
ure 2). On one side of the roller track, the caster from
Figure 1 usually has a small radius with options in real
time to adjust the positions of the rollers and, thereby,
control the reduction profile. The reduction can be
carried out only in the place where the centre of the slab
is still stiff and the optimum point is the end of the soli-
dification (the metallurgical length). This paper aims to
find out when it is necessary to use soft reduction during
steel casting and when it is desirable to have linear or
other profiles. To answer these questions a model of the
temperature field4 is applied.
2 MODEL OF THE SLAB TEMPERATURE
FIELD
The 3D model was first designed as an off-line ver-
sion and later as an on-line version so that it worked in
real time. After being corrected and tested and thanks to
the universal nature of the code, it is possible to imple-
ment it on any caster. The numerical model takes into
account the temperature field of the entire slab (from the
meniscus of the level of the melt in the mould to the
cutting torch) using a 3D mesh containing more than a
million nodal points.
The solidification and cooling of a concast slab are
global problems of the 3D transient-heat and mass trans-
fer. If the heat conduction within the heat transfer in this
system is decisive, the process is described with the
Fourier-Kirchhoff equation. It describes the temperature
field of a solidifying slab in all three states: at the tempe-
ratures above the liquidus (i.e., the melt), within the
interval between the liquidus and solidus (i.e., in the
mushy zone) and at the temperatures below the solidus
(i.e., the solid state). For reliable solutions it is appro-
priate to use the explicit numerical method of control
volumes. A numerical simulation of the release of the
latent heats of the phase or structural changes is carried
out by introducing the enthalpy function dependent on
temperature T, preferably in the form of the enthalpy
related to unit volume Hv, containing the latent heats.
After an automated generation of the mesh (pre-process-
ing), the model uses the thermophysical material pro-
perties of the investigated system, including their depen-
dence on the temperature – in the form of tables or using
polynomials: heat conductivity k, specific heat capacity c
and density 
 of the cast metal. The temperature distri-
bution in the slabs is described with the enthalpy balance
equation.
A simplified Equation (1), suitable for application on
radial casters with a large radius, where only the speed
J. STETINA et al.: CONTROL OF SOFT REDUCTION OF CONTINUOUS SLAB CASTING WITH A THERMAL MODEL
726 Materiali in tehnologije / Materials and technology 49 (2015) 5, 725–729
Figure 2: Diagram of soft reduction
Slika 2: Shematski prikaz mehke redukcije
Figure 1: Slab caster
Slika 1: Naprava za ulivanje slabov
Figure 3: Enthalpy functions of steels
Slika 3: Funkcija entalpij jekel
(of the movement of the slab) component w in the
z-direction is considered, is:
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
H
t z
w H k
T
x
T
y
T
zz
v
v+ = + +
⎛
⎝
⎜
⎞
⎠
⎟( )
2
2
2
2
2
2 (1)
Volume enthalpy Hv as a thermodynamic function of
the temperature must be known for each specific steel. It
depends on the steel composition and the cooling rate.
The dependence of function Hv for three different steels
is shown in Figure 3.
The unknown enthalpy of the general nodal point of
the slab in the next time step ( + ) is expressed with
the following explicit formula:
H H Qz Qz Qy Qy
Q
i j k i j k i j i j i iv v, , , ,
( ) ( )
, ,(
 = + + + + +
+
1 1
x Qx
x y z
1+ ⋅
⋅ ⋅
)

  
(2)
The heat flow through the general nodal point (i, j, k)
in the z-direction is described with Equation (3):
Qz k
A
z
T T A w Hi j
z
i j k i j k z z i j k, , ,
( )
, ,
( ) ( )( )
, ,
= − −+Δ 1
  
v (3)
The initial condition for the solution is the setting of
the initial temperature in individual points of the mesh.
The appropriate value is the highest possible tempera-
ture, i.e., the casting temperature. The boundary con-
ditions in different places and different systems are
described with Equations (4a) to (4e):
1. T = Tcast at the meniscus (4a)
2. − =k
T
n
∂
∂
0 at the plane of symmetry (4b)
3. − = −k
T
n
a b
L
w z
∂
∂
mould
in the mould (4c)
4. [ ]− = + + + ⋅k T
n
h T T T T
∂
∂ tc surface
2
amb
2
surface amb 0 ( )( )
⋅ −( )T Tsurface amb in the secondary- and tertiary-cooling
zones (4d)
5. − = − −k
T
n
T w z
∂
∂
11513 7 20 76 0 2 0 16. ( ). . .surface  beneath the
support rollers5. (4e)
3 PARAMETRIC STUDY OF SLAB CASTING
For this parametric study we selected three steels
with different chemical compositions, where the prob-
lems of internal defects4 were solved. The chemical com-
positions are shown in Table 1, which also shows the
liquidus and solidus temperatures.
Solidification-analysis package IDS using the che-
mical compositions for calculating the thermophysical
properties is used to model the temperature field.6 As can
be seen from Figure 4 the density in the temperature
range between the liquidus temperature and 800 °C is
almost identical for all three investigated steel grades.
J. STETINA et al.: CONTROL OF SOFT REDUCTION OF CONTINUOUS SLAB CASTING WITH A THERMAL MODEL
Materiali in tehnologije / Materials and technology 49 (2015) 5, 725–729 727
Figure 4: a) Density and b) contraction of the selected steels
Slika 4: a) Gostota in b) kr~enje izbranih jekel
Table 1: Selected types of steel with their compositions used for the calculation
Tabela 1: Kemijska sestava izbranih jekel, uporabljenih pri izra~unu
Class Mark Group C Mn Si P S Cu Ni Cr Mo V Ti Al Nb Tsolidus Tliquidus
1.0045 S355G8 1 0.011 1.550 0.450 0.075 0.025 0.000 0.000 0.000 0.00 0.000 0.000 0.375 0.020 1443.6 1514.1
1.8978 L555MB 2 0.050 1.600 0.300 0.150 0.010 0.250 0.150 0.250 0.000 0.000 0.015 0.350 0.055 1466.0 1516.4
1.6580 30CRNIMO 3 0.280 0.450 0.250 0.125 0.025 0.100 1.900 1.900 0.350 0.000 0.000 0.300 0.000 1371.5 1493.8
An analogous conclusion can be made for the contrac-
tion (Figure 4). However, the cross-sectional reduction
of a cast blank has to be considered for the temperatures
lower than 800 °C, together with the influence of the
structural-phase changes in the solid state. In such cases,
the casting control has to take into account the
cross-sectional reduction due to the phase changes, or
the secondary-cooling requirements, to be properly up-
dated to keep the temperature of the blank above 800 °C.
The next step is a parametric study of the simulations
of casting the three steel classes and three casting speeds
((0.8, 1.0 and 1.2) m/min), representing the expected
caster operating range. The calculated waveforms of the
isolines of the liquidus and solidus are shown in Figure 5.
These calculations and other parameters, such as the
temperature superheat, the primary and secondary cool-
ing, are set in equal forms. A constant reduction of the
profile is considered, so the resulting slab thickness is
200 mm (Figure 6).
Figures 6 and 7 show the calculated contraction
profiles for the selected classes of steel. In both cases the
calculation considered is calculated for a casting speed
of 1 m/min and the conditions corresponding to Figure 5
as the basic requirements for the output profile on the
machine and a slab thickness of about 200 mm. The
thick dashed line marks the roller positions obtained with
the narrowing of the machine profile. To allow the use of
soft reduction the input size of a blank must be larger.
The difference in the size of the mould is 205 mm for the
caster without soft reduction and 210 mm for the caster
with soft reduction.
The temperature field in the form of mileage iso-
curves of the liquidus and solidus considering soft reduc-
tion is shown in Figure 8. With respect to the position of
the isolines of the solidus and liquidus temperature, the
temperature field is virtually the same, as shown in the
middle of Figure 5. However, it should be understood
that the initial thickness of the slab is larger by 5 mm
because soft reduction shortens the metallurgy length.
4 CONCLUSION
This paper presents a 3D numerical model of a tem-
perature field (for concasting steel) using the in-house
software and its applications for the initial study that
compares the waveforms of liquidus and solidus isolines
J. STETINA et al.: CONTROL OF SOFT REDUCTION OF CONTINUOUS SLAB CASTING WITH A THERMAL MODEL
728 Materiali in tehnologije / Materials and technology 49 (2015) 5, 725–729
Figure 5: Liquidus and solidus temperatures for casting speeds of
(0.8, 1.0 and 1.2) m/min
Slika 5: Likvidusna in solidusna temperatura za hitrosti ulivanja (0,8,
1,0 in 1,2) m/min
Figure 8: Course liquidus and solidus isolines when applying soft
reduction in segments 6 and 7
Slika 8: Likvidusne in solidusne izolinije pri uporabi mehke redukcije
v segmentih 6 in 7
Figure 6: Contraction due to the thickness and slab-temperature
decreases and the setting of the roll gap
Slika 6: Profil kr~enja zaradi zmanj{anja debeline in temperature
slaba ter nastavitev re`e med valji
Figure 7: Contraction profile that uses soft reduction in segments 6
and 7
Slika 7: Profil kr~enja pri uporabi mehke redukcije v segmentih 6 in 7
and strand shrinkage. Furthermore, we compared a caster
with a constant constriction, i.e., with a fixed adjustment
between the roller gap and the caster with adjustable
rollers used for controlling the narrowing, i.e., soft
reduction. The results of this analysis will be used in a
mathematical model designed for the management of
soft reduction. The mathematical model will be inte-
grated in the control systems at all the levels within the
automation pyramid that is being developed in the ASM
Automation Company (www.asmautomation.com). One
of the future tasks is a study of the stochastic behaviour
of the continuous casting process7.
Acknowledgments
This work is a result of the research and scientific
activities of the NETME Centre, the regional R&D cen-
tre built with the financial support from the Operational
Programme Research and Development for Innovations
within a NETME Centre project (New Technologies for
Mechanical Engineering), Reg. No. CZ.1.05/2.1.00/01.
002 and, in the follow-up sustainability stage, supported
through NETME CENTRE PLUS (LO1202) with the
financial means from the Ministry of Education, Youth
and Sports of the Czech Republic under the National
Sustainability Programme I.
Nomenclature
a constant of the mold (W/m2)
A area (m2)
b constant of the mold (W/(m2 s0.5))
c specific heat capacity (J/(kg K))
htc heat-transfer coefficient (W/(m2 K))
Hv volume enthalpy (J/m3)
k heat conductivity (W/(m K))
Lmould length of the mould (m)
T temperature (K)
Tamb ambient temperature (°C, K)
Tcast casting temperature (°C, K)
Tsurface temperature of the surface (°C, K)
q density of the heat flow (W/m2)
Qx, Qy, Qz heat flows (W)
x, y, z axes in a given direction (m)
wz casting speed in a given direction (m/s)

 density (kg/m3)
o Stefan-Bolzmann constant (W/(m2 K4))
 emissivity (–)
 angle of the contact (°)
 time (s)
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4 J. K. Brimacombe, The Challenge of Quality in Continuous Casting
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Materiali in tehnologije / Materials and technology 49 (2015) 5, 725–729 729

T. SKALAR et al.: OPTIMIZATION OF OPERATING CONDITIONS IN A LABORATORY SOFC TESTING DEVICE
731–738
OPTIMIZATION OF OPERATING CONDITIONS IN A
LABORATORY SOFC TESTING DEVICE
OPTIMIZACIJA OBRATOVALNIH RAZMER LABORATORIJSKE
GORIVNE CELICE SOFC
Tina Skalar, Martin Lubej, Marjan Marin{ek
University of Ljubljana, Faculty of Chemistry and Chemical Technology, Ve~na pot 113, 1000 Ljubljana, Slovenia
marjan.marinsek@fkkt.uni-lj.si
Prejem rokopisa – received: 2014-08-22; sprejem za objavo – accepted for publication: 2014-12-01
doi:10.17222/mit.2014.209
Solid-oxide fuel cells (SOFCs) are devices that convert chemical energy into electrical energy. They are assembled from three
solid components, which, in principle, differ in their thermal expansion coefficients. These differences may cause residual
stresses during an operation and consequently lead to physical damage of a cell. From this perspective, the reliability of an
operating SOFC is seriously dependent on the thermal stress built inside the multi-layer structure. In this work, some operating
conditions in the SOFCs based on an SDC electrolyte, an LSM cathode and a Ni-SDC anode were investigated both experi-
mentally and theoretically. Within the theoretical analysis, the residual stresses and temperature profiles inside an operating cell
were modelled with the finite-element method. The results of the mathematical modelling of the warm-up, steady-state or cool-
down periods were used to optimize the cell geometry and the thickness of individual cell layers, and to determine the most
appropriate operating cell conditions. In order to experimentally confirm some theoretical calculations, a new SOFC testing
system was developed, which enabled relatively easy assembling and dismantling as well as quick changes in the operating
conditions (temperature, atmospheres). Based on the modelling results, optimization of the operating conditions was proposed
in order to reduce the thermal stresses built in the materials.
Keywords: SOFC, modelling, temperature gradient, thermal stress, optimization
Visokotemperaturne gorivne celice pretvarjajo kemi~no energijo preko katalitskih reakcij v elektri~no energijo. Sestavljene so iz
treh trdnih (kerami~nih) materialov, ki imajo razli~ne temperaturne razteznostne koeficiente. Razli~ni raztezki materialov
povzro~ajo mehanske napetosti, kar ima lahko za posledico degradacijo (razpad) strukture ve~slojnega sistema. Zanesljivost
gorivne celice je tako odvisna od ustvarjenih mehanskih napetosti v materialih. V tem delu smo raziskovali razmere pri obrato-
vanju tako eksperimentalno kot teoreti~no za ve~plastno preizkusno tableto. Zapisali smo matemati~ni model za opis generi-
ranih mehanskih napetosti med segrevanjem, v stacionarnem stanju in med ohlajanjem ter z metodo kon~nih elementov pri{li do
numeri~ne re{itve modela. Dobljene rezultate smo uporabili pri optimizaciji geometrije celice in debeline posameznih plasti v
ve~plastni tableti. Za eksperimentalno preverjanje dobljenih rezultatov modeliranja smo razvili nov preizkusni sistem, ki
omogo~a relativno hitro sestavljanje in razstavljanje ter hitro spremembo obratovalnih razmer (temperature, atmosfere). Na
osnovi dobljenih rezultatov je bila predlagana optimizacija preizkusnega sistema za zmanj{anje mehanskih napetosti v
materialih.
Klju~ne besede: SOFC, modeliranje, temperaturni gradienti, termi~ne napetosti, optimizacija
1 INTRODUCTION
Solid-oxide fuel cells (SOFCs) are highly promising
and environmentally friendly energy converters with a high
energy-conversion efficiency, converting chemical energy
of reactants into electrical energy.1–4
SOFCs, as any other batteries, consist of three diffe-
rent layers, called SOFC membranes, i.e., an electrolyte,
an anode and a cathode, and operate at temperatures bet-
ween 600 °C and 1000 °C. As an electrolyte, samaria-
doped ceria (SDC), a well-known advanced ceramic
material, was recently proposed to reduce the SOFC
operating temperatures due to its high oxygen-ion con-
ductivity.5–9
A cermet of metallic nickel and ceramic, i.e., SDC
(Ni-SDC), is most commonly used as the anode material
in a SOFC. The main reasons for the selection of Ni-
SDC are its thermal-expansion coefficient that is com-
parable with the SDC electrolyte, a relatively good ionic
and electronic conductivity and a catalytic activity
acceptable for hydrogen or hydro-carbon electro-oxida-
tion. On the cathode side, doped lanthanum manganite
(i.e., lanthanum strontium manganite, LSM) is most
commonly used.10–13
Since SOFC membranes, in principle, are composed
of various materials, their thermal-expansion coefficients
(TEC) may differ. For instance, the TEC of an anode
cermet is normally somewhat higher than those of the
electrolyte and cathode. These differences in the TECs
may build stresses in the SOFC membrane during a tem-
perature alteration and, consequently, lead to a SOFC-
membrane delamination. From this perspective, the
reliability of an SOFC membrane is seriously dependent
on the SOFC operating conditions, especially during its
heating to the operating steady-state conditions or cool-
ing down to room temperature.14–18
One way to approach the problem of determining the
optimum operating conditions of an SOFC system is to
Materiali in tehnologije / Materials and technology 49 (2015) 5, 731–738 731
UDK 621.35/.36:621.352.6 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)731(2015)
employ mathematical modelling. In this way, several
parameters, i.e., temperature profiles and stress fields,
can be followed without performing complicated and
lengthy SOFC experimental tests. The finite-element-
method (FEM) modelling was already successfully used
by several authors to address similar engineering prob-
lems.19–22
With the FEM, a continuous problem can be trans-
formed into a discrete problem with a finite number of
degrees of freedom.23–25
In this paper, the optimum operating parameters in
which an SOFC multilayer membrane is not exposed to a
critical thermal stress caused by temperature gradients
are determined. Additionally, the problem of the opti-
mum cell geometry and the thicknesses of individual cell
layers during the warm-up or steady-state periods is
addressed.
2 EXPERIMENTAL PROCEDURE
2.1 Material and membrane preparation
An electrolyte made of samarium-doped ceria (SDC)
or an anode made of nickel oxide samarium-doped ceria
(NiO-SDC) and a lanthanum manganite (LSM) cathode
were synthesized using the modified Pechini method and
carbonate precipitation, respectively. The exact synthesis
procedures are described in26,27. Before a multi-layer
membrane was prepared, a supporting layer (electrolyte
or anode) was pressed into a tablet and pre-sintered at
1000 °C in air in order to ensure the mechanical strength
for the screen printing of other layers. The thickness of
the screen-printed layers was 20 μm. If an anode-
supported system was prepared, an SCD electrolyte was
first screen-printed on a NiO-SDC pre-sintered anode.
Such a multi-layer system was co-sintered at 1250 °C in
order to form a good contact between the layers. The
same procedure was used for electrolyte-supported
systems with an exception, in which an SDC was first
pre-sintered and a NiO-SDC was then screen printed and
co-sintered. After the co-sintering, an LSM cathode was
screen printed on the electrolyte, opposite to the anode
side. Such a multi-layer system was again co-sintered at
1200 °C.
2.2 Testing-system assembly
The testing system (Figure 1) was composed of a
multilayer tablet (anode/electrolyte/cathode), which was
sealed in steatite ceramics. The steatite ceramics was
placed between two quartz tubes to prevent the mixing of
the inner and outer atmospheres, or the mixing of the
cathode oxidative and anode reductive atmospheres
(Viton o-rings were used for sealing). The whole system
was then placed into another protective quartz tube.
Heaters were mounted outside the protective quartz tube.
The inner and outer quartz tubes were embedded into a
stainless-steel (1.4404) case and fixed to the construction
with four screws. The testing system was isolated with a
-alumina insulation 2 cm. The NiCr-Ni thermocouple
was pushed through a small hole in the stainless-steel
case and coupled with the display.
2.3 Used model
The gas flow and temperature distribution were simu-
lated using numerical methods. The transport of mass
and energy was calculated using the continuum theory.28
The heat was introduced as a homogeneous-volume heat
source at the heating element. The heat was dissipated at
the outer boundaries with a heat flux corresponding to
the wall temperature and the heat-transfer coefficient
obtained with the Churchill and Chu correlation.29 The
system of differential equations was solved utilizing the
finite-elements method in a 2D axisymmetric geometry,
and the stationary state was achieved at the 35th iteration.
The mechanical stress inside the tested membrane
was calculated using a linear elastic material model30 in
combination with the thermal expansion. The reference
temperature was defined as the temperature, at which
there is no stress in the material. The mechanical model
was solved utilizing the finite elements in a 2D
axisymmetric geometry.
For all the materials, the average emissivity () of 0.8
was used. Young’s modulus (E), Poisson’s ratio (), the
density (
), the thermal-expansion coefficient () were
200 · 109 Pa,23 0.32,23 7600 kg/m3, 11.5 · 10–6 K–1,31 55 ·
109 Pa,23 0.17,27 8900 kg/m3, 13 · 10–6 K–1 23 and 40 · 109
Pa,31 0.25,31 6500 kg/m3,31 11.4 · 10–6 K–1 31 for the SDC,
NiO-SDC and the LSM cathode, respectively. The den-
sities for the anode and electrolyte materials were cal-
culated as Archimedes’ densities from the dimensions of
the sintered bodies.
3 RESULTS AND DISCUSSION
Finding the optimum design and operating conditions
makes the SOFC testing rather time-consuming. More-
over, in order to track various parameters during an
SOFC continuous run, the testing systems are usually
rather complicated and difficult to assemble, making
visual observing of a tested multilayer SOFC membrane
practically impossible. For this reason, a new SOFC test-
ing system was developed, in which a multilayer tablet
(anode/electrolyte/cathode) is sealed in a steatite-cera-
mics holder and embedded into a double quartz-glass
tube, which is then put into a tube oven (as presented in
the experimental work, Figure 1).
Such a system has a rather simple design, is easy to
assemble/dismantle or manipulate, and enables quick
changes in the operating conditions (temperature, atmo-
spheres). The true advantage of the simplified SOFC
design, however, is the fact that it enables the start of
new experiments in short cycles, thereby overcoming the
time-consuming difficulty of more complex testing
732 Materiali in tehnologije / Materials and technology 49 (2015) 5, 731–738
T. SKALAR et al.: OPTIMIZATION OF OPERATING CONDITIONS IN A LABORATORY SOFC TESTING DEVICE
systems. However, short testing cycles are always related
to rapid temperature changes inside the testing system. It
is well known that the temperature increases or decreases
in a ceramic multilayer system with different tempera-
ture-expansion coefficients, thereby building stresses that
may initiate crack formation, which eventually leads to a
complete system delamination. In order to avoid such an
undesirable multilayer degradation, two approaches are
possible: i) numerous tests to determine the temperature
regime during a testing cycle, or ii) employing a mathe-
matical simulation model and predicting the temperature
gradients as well as the built stress fields in the tested
membrane, as was the case in this work.
The first addressed question is the position of the
heating elements in the testing system as the heating
elements may be placed inside or outside the tested
SOFC cell. Positioning the heating elements inside each
of the SOFC chambers (the cathode and anode sides) has
the advantage of a very fast temperature change, as
shown in our previous work.24 However, the heating of
the individual chambers and the tested membrane posi-
tioned between both chambers with two point sources
placed above and below the membrane also causes some
undesired temperature gradients in the membrane, where
the center of the membrane is always hotter than its outer
regions. These temperature gradients are the greatest
during the membrane’s rapid heating but cannot be
eliminated completely even in the steady operating state
due to the heat exchange between the SOFC testing cell
and its surroundings. For this reason, the heating ele-
ments were placed around the double quartz tube. In
such a modification, additional heat-insulating layers
between the heater and the membrane are introduced (the
quartz tube itself). However, due to the powerful heating
elements, rapid temperature changes are still possible
(Figure 2) even with the heaters placed outside the
quartz tubes.
The newly developed SOFC testing design enables
the work with multi-layer membrane systems of various
shapes, e.g., a cube or a cylinder. Altering the shape of a
multi-layer system also causes some changes in the tem-
perature distribution in the heated membrane. The tem-
perature gradients are much more pronounced alongside
the radius of the membrane due to the relative position of
the heating elements and practically negligible alongside
the z-axis. Furthermore, the heating mode of the multi-
layer membrane suggests that the biggest temperature
gradients appear during the membrane heating (shortly
after the heaters start to heat up the outer quartz tube)
and gradually declines when the temperature of the inner
chambers of the SOFC testing system approaches its
final value. According to the results of the mathematical
modelling, the maximum temperature gradients from the
membrane center to its edge T1 are  90 °C or  75 °C
for cube or cylinder membranes, respectively (Figure 3).
This rather noticeable difference in the maximum
temperature gradients is introduced by the design of the
SOFC testing system itself. Since the heating elements
are positioned cylindrically around the quartz tubes, the
outer positions of the cube membrane are unevenly
distant from the heaters and thus exposed to various heat
fluxes. At the same time, due to the relatively thin mul-
ti-layer membrane ( 3 mm), the temperature gradient
alongside the z-axis, T2, is only 2.2 °C for both mem-
brane designs. Since the temperature gradients inside the
T. SKALAR et al.: OPTIMIZATION OF OPERATING CONDITIONS IN A LABORATORY SOFC TESTING DEVICE
Materiali in tehnologije / Materials and technology 49 (2015) 5, 731–738 733
Figure 1: Scheme of a test cell: a) anode inlet/outlet atmosphere, b)
cathode inlet/outlet atmosphere and c) tube oven
Slika 1: Shema preizkusne celice: a) vpih in izpih anodnega plina, b)
vpih in izpih katodnega plina in c) obodni grelnik
Figure 2: T versus t relationship measured at the inner side of the
quartz tube during the SOFC testing-system start-up; the small insert
represents the measuring point
Slika 2: Odvisnost T od t, merjena na notranji strani kremenove cevi
pri zagonu sistema SOFC; izrez prikazuje to~ko merjenja temperature
Figure 3: Temperature distribution in the tested membranes: a) cubic
membrane and b) cylindrical membrane
Slika 3: Porazdelitev temperature v preizkusni membrani: a) v obliki
kvadra in b) v obliki valja
tested membrane are one of the prime causes for the
induced thermal stresses (as discussed later) a cubic
membrane is much more susceptible to potential crack-
ing. For this reason, such membranes are omitted from
further modelling.
Starting a new testing cycle and heating up the mem-
brane is a critical point that decisively influences the
survival chances of a multi-layer membrane. In order to
obtain a greater insight into the circumstances during the
system heating, the temperature distribution inside the
testing system (including the tested membrane) is mathe-
matically modelled at several pre-chosen times and pre-
sented in Figure 4. For mathematical modelling, a
multilayer electrolyte-supported membrane with thick-
nesses of the individual layers (anode/electrolyte/catho-
de) being 20 μm/2000 μm/20 μm, respectively, is
assumed. From the sequence of the pictures showing the
temperature distribution, it is evident that the tempe-
rature inside the testing system starts to rise due to the
hot quartz tube that heats the adjacent gas and causes its
circulation. The system design, where a stainless-steel
case is used to cover the quartz tubes at both ends, is the
reason for further temperature increases. More precisely,
stainless steel heats up relatively quickly, serving as an
additional heat source for the gases inside individual
chambers causing their increased circulation. The ver-
tical position of the testing system forces the hot gases to
flow alongside the z-axis. Due to this vertical flow of the
hot gas inside the bottom chamber, the bottom plane of
the membrane heats up more quickly than its top plane,
causing temperature gradients in the vertical cross-sec-
tion throughout the membrane (the temperature distri-
butions at various times throughout the membrane are
presented in Figure 4). At some critical point, the
temperature gradient T2 in the membrane is maximal
(Figures 4c and 4h). This critical time was determined
as 160 s or 300 s if the heating element worked with its
maximum power (powered with 220 V) or with just
46 % of its maximum power (powered with 150 V),
respectively. In the latter case, the temperature inside the
SOFC testing system typically rises more slowly, also
reducing the built temperature gradients T1 and T2
(Figure 5). Since the membrane is a multilayer system in
which individual layers exhibit different tempera-
ture-expansion coefficients, temperature gradient T2 is
particularly important. Dissimilar temperature-expansion
coefficients, together with temperature gradient T2,
may introduce a situation in which the expansion of an
individual layer during the membrane heating differs
from the expansion of the adjacent layer to a degree that
causes a membrane delamination. After the critical time,
temperature gradient T2 slowly diminishes and beco-
mes negligible in the steady operating state (Figures 4e
and 4j).
A verification of the modelled temperature fields was
conducted via additional temperature measurements
T. SKALAR et al.: OPTIMIZATION OF OPERATING CONDITIONS IN A LABORATORY SOFC TESTING DEVICE
734 Materiali in tehnologije / Materials and technology 49 (2015) 5, 731–738
Figure 4: Temperature distributions in both inner chambers of the SOFC testing system and throughout the membrane at various times presented
in a vertical plane: a) to e) the heating element is powered with 220 V, f) to j) the heating element is powered with 150 V
Slika 4: Porazdelitev temperature, prikazana v vertikalni ravnini v preizkusni celici po razli~nih ~asih segrevanja: a) do e) pri nazivni napetosti
grelnika 220 V, f) do j) pri nazivni napetosti grelnika 150 V
using NiCr-Ni thermocouples. The measuring points at
each specific time are marked in Figure 4. It is evident
that the mathematical model used describes the tempe-
rature field in the tested system rather well (Table 1).
Temperature mismatches between the modelled and
measured temperatures in either point and at each spe-
cific time do not exceed 8 °C. These slight mismatches
are ascribed to a somewhat difficult positioning of the
thermocouple at the predetermined point and in the
dynamic temperature field (the response time of the
NiCr-Ni thermocouple is  2 s).
Table 1: Comparison between modelled and measured temperatures
(marked point in Figure 4) in the tested system at various heating
times
Tabela 1: Primerjava modelirane in merjene temperature (to~ka
merjenja je ozna~ena na sliki 4) v preizkusnem sistemu pri razli~nih
~asih segrevanja
220 V 150 V
theating /s Tmodelled/K
Tmeasured
/K theating /s
Tmodelled
/K
Tmeasured
/K
40 306 304 40 304 303
160 882 877 300 756 750
7200 1379 1371 7200 1233 1227
Another important issue addressed by mathematical
modelling is the fluid-velocity simulation inside both
SOFC chambers (Figure 6). According to the simula-
tion, the causes of the gas turbulence in the SOFC cham-
bers are the fresh-gas inlet and the created temperature
gradients. The cylindrical design of the SOFC testing
system causes the gas velocity to be the highest close to
the chamber center ( 2.2 · 10–3 m s–1) and relatively low
near the surface of the tested membrane. The complex
movement path of gas molecules is indicated by their
trajectories. The fluid-velocity pictures inside the anode
and cathode chambers, however, are not mirror-inverted.
Due to the vertical rise of the hot gases, the region of
high-gas velocity in the bottom chamber is closer to the
membrane relative to the top chamber. Despite the
lowest gas velocity of  3 μm s–1 near the membrane
surface (10 μm from the surface), no gas pockets with a
stationary atmosphere could be observed, which is
important from the practical viewpoint of an operating
SOFC. Only dynamic atmosphere conditions at the
membrane surface may deliver fresh reactants to the
diffusion layer of the porous-membrane catalyst and
remove gaseous products. Additionally, it can be pre-
sented with the mathematical modelling that the intro-
duction of fresh cold gases negligibly influences the
temperature fields inside both of the SOFC chambers.
The gases entering the anode or cathode chambers are
practically instantly heated up to the temperature of the
chambers operating at a steady state. This fact is not sur-
prising since the gas inlets are relatively low ( 15 mL
min–1) while the chambers are relatively large and only
 72 W is needed to heat a gas flow of 15 mL min–1
from room temperature to the chamber temperature.
In an SOFC multi-layer system, in principle, three
different materials are combined in a single membrane.
Since the temperature-expansion coefficients of indivi-
dual layers are not identical, any temperature alteration
may build a stress alongside a phase border and, conse-
quently, cause a delamination of the multi-layer mem-
brane. The first critical point in the lifecycle of a multi-
layer membrane arises during its preparation. Specifi-
cally, the membrane is prepared by co-sintering all three
layers and, subsequently, reducing NiO to Ni in the
anode composite. During the sintering, a plastic defor-
mation of the material is allowed. From this perspective,
it can be assumed that the system is stress-free at the sin-
tering temperature. In the testing system, the stress-free
state is adopted at 1250 °C, at which the NiO/SDC
T. SKALAR et al.: OPTIMIZATION OF OPERATING CONDITIONS IN A LABORATORY SOFC TESTING DEVICE
Materiali in tehnologije / Materials and technology 49 (2015) 5, 731–738 735
Figure 6: Fluid-velocity simulation together with gas trajectories in
both inner chambers
Slika 6: Simulacija gibanja plinov skupaj s trajektorijami plina v no-
tranjih prostorih preizkusnega sistema
Figure 5: Relationships of induced temperature gradients T1 and
T2 with t, presented in a vertical plane through the multi-layer
membrane
Slika 5: Inducirani temperaturni gradienti T1 in T2 v odvisnosti od
~asa t, prikazani v vertikalni ravnini skozi ve~plastno membrano
anode, the SDC electrolyte, and the LSM cathode are
co-sintered. The subsequent reduction of NiO to Ni
should not induce any additional stress since Ni may
undergo a plastic deformation. Thermal stresses in the
multilayer system start to build up below the plastic-
deformation temperature, reaching their peak at room
temperature.32 Depending on the temperature-expansion
coefficients of the adjacent layers (13 · 10–6 K–1,
11.5 · 10–6 K–1 and 11.4 · 10–6 K–1 for Ni/SDC, SDC and
LSM, respectively) these built-up stresses may be as
high as several hundred MPa and may cause cracks to
form.
The residual stresses after the membrane preparation
are practically impossible to measure directly; however,
by adopting an appropriate mathematical model, they
can be calculated rather easily. Thus, if Ni/SDC-SDC-
LSM (20 μm-3000 μm-20 μm) is cooled down to room
temperature (after the reduction of NiO to Ni) the
residual thermal stresses are calculated as  130 MPa
for Ni/SDC of the anode layer, close to the anode-elec-
trolyte phase boundary,  9 MPa for the electrolyte
adjacent to the anode layer, and the lowest stress for
LSM (< 1 MPa), at the SDC-LSM phase boundary.
Together with the calculated number, the nature of the
built stresses must also be taken into consideration. Due
to a higher thermal-expansion coefficient, the Ni/SDC
layer tends to shrink more than the SDC electrolyte,
while LSM shrinks to approximately the same degree as
SDC. In a hypothetical situation in which two layers sub-
jected to uneven material shrinkage have approximately
the same thickness, the induced stress should bend the
bilayer system where the layer with the smaller shrink-
age forms a concave surface at the phase boundary. In
view of the chosen set of materials and the actual mem-
brane design, the degree of shrinkage suggests that the
stress in the anode layer in the cooled-down state is
always tensile, while the electrolyte layer is under
compression. Therefore, at room temperature, the anode
appears to be the potential weak point of the multi-layer
system. However, the maximum calculated tensile stress
in the anode layer does not exceed the Ni/SDC tensile
strength, which is  150 MPa.33 As the final result, the
calculated values suggest that the multi-layer system
should survive the first thermal cycle (sintering and
reduction), which is also in accordance with practical
observations.
During the repeated heating from room temperature
to operating temperatures (700–800 °C) it would be
expected that the stresses in the multi-layer membrane
slowly diminish since the increasing temperature brings
the system closer to a stress-free state. However, it turns
out that the picture of the stresses built during the heat-
ing is far more complex. As mentioned previously, the
membrane is not heated up evenly. Instead, the outer
layers of the membrane are heated up faster than the
inner layers due to the circulating hot gases in the top
(anode) and bottom (cathode) chambers. Due to these
temperature gradients and the fact that the layers have
different temperature-expansion coefficients, it appears
that the center of the membrane (relative to the membra-
ne outskirts) is always at higher loads regarding the built
stresses (Figure 7). The adopted mathematical model
predicts that the maximum stress is created at the critical
time when the temperature gradient T2 is maximal. In a
situation in which the stress in the membrane (a thick-
ness of 20 μm Ni/SDC, 2000 μm SDC, 20 μm LSM, a
diameter ø = 11 mm, the critical time t = 160 s) is mo-
delled, it can be calculated that any given temperature
above the cooled-down state causes some relaxation and
reduces the tensile stress in the anode layer. In contrast,
the relatively thick electrolyte adjacent to the anode
stretches unevenly in accordance with the created tem-
perature gradient. In addition, the stress in the electrolyte
layer near the anode/electrolyte phase boundary is also
built up due to anode stretching, meaning that the elec-
trolyte is under the tensile stress at critical times. The
absolute value of the induced tensile stress in the elec-
trolyte layer after  160 s of the heating at the maximum
power (220 V) is  70 MPa (Figure 7a). This value
exceeds the tensile strength of SDC, which is 53 MPa.34
When the maximum tensile strength is surpassed, the
material is prone to cracking. If the heating power is re-
duced (46 % of the heater’s maximum power; the heaters
are powered with 150 V), the induced tensile stress in the
electrolyte at the critical time ( 300 s) decreases to 27
MPa, which is acceptably low and demonstrates that a
somewhat slower heating rate significantly reduces the
stress load in the membrane.
Another interesting situation is demonstrated in
Figure 7b, in which the membrane diameter is increased
to ø = 18 mm. According to the mathematical modelling,
membranes with larger diameters will build higher
stresses under the same heating conditions. The absolute
value of the induced tensile stress in the electrolyte layer
at the critical time ( 160 s) thus increases to  83 MPa
T. SKALAR et al.: OPTIMIZATION OF OPERATING CONDITIONS IN A LABORATORY SOFC TESTING DEVICE
736 Materiali in tehnologije / Materials and technology 49 (2015) 5, 731–738
Figure 7: Induced stresses in the membrane (20 μm–2000 μm-20 μm,
t = 160 s) presented in a vertical plane: a) ø = 11 mm, b) ø = 18 mm
Slika 7: Inducirane napetosti v membrani (20 μm–2000 μm-20 μm, t =
160 s), predstavljene v vertikalni ravnini: a) ø = 11 mm, b) ø = 18 mm
(the heaters powered with 220 V) or  30 MPa (the
heaters powered with 150 V).
Since membranes in modern SOFC systems may be
designed as anode- or electrolyte-supported, the next
logical task addresses the thicknesses of individual layers
and the associated stress load in the membrane during
the heating. This question was approached by conducting
a series of mathematical modelling tasks, in which the
built stresses versus the thicknesses of individual layers
were calculated (Figure 8). The investigated scenario
includes the membrane discs (ø = 11 mm) in which the
thicknesses of the Ni/SDC anode or SDC electrolyte
layers vary from 20 μm to 3000 μm (the thickness of the
LSM cathode layer is constant, being 20 μm). Since the
most challenging conditions, with respect to the built-up
stress, occur at the critical time of the heating when the
temperature gradients are the highest (160 s or 300 s if
the heaters are powered with 220 V or 150 V, respec-
tively), only the maximum stresses built in the individual
layers are presented. According to the results, several
distinct features can be easily identified. Firstly, during
the heating, the stress in the anode is always compres-
sive, while the stresses in the electrolyte and cathode are
tensile. Secondly, the stress in one component decreases
with an increase in its thickness, when the thicknesses of
the other two components are fixed. In addition, de-
creases in the compressive stress in the anode layer cause
increases in the tensile stress in both the electrolyte and
the cathode and vice versa. Thirdly, the thicknesses of
the anode and the electrolyte have a pronounced influ-
ence on the stresses in the adjacent layers. Fourthly, due
to very similar temperature-expansion coefficients of
SDC and LSM, the tensile stress in the cathode layer
never exceeds its tensile strength, which is assumed as
138 MPa.34
The absolute values of the calculated stress also
reveal that the electrolyte layer is most likely to crack.
The maximum tensile stress in the electrolyte (adjacent
to the anode) reaches  180 MPa if the membrane is
composed of three 20 μm layers and heated with the
maximum heating power (Figure 8a). Even if the
electrolyte layer is broadened to 3000 μm, the maximum
tensile stress is still  50 MPa, which is very close to the
tensile strength of SDC. Under the same conditions, the
tensile stress in the cathode and the compressive stress in
the anode change from  40 MPa to less than 10 MPa
and, therefore, never reach the tensile or compressive
strength of LSC and Ni/SDC, respectively. In order to
improve the survival chances of the electrolyte, the
heating rate has to be reduced. If the heaters are powered
with 150 V, an electrolyte thickness of  2 mm or more
is enough to avoid the exceeding overcritical tensile
stress. Powering the heaters with 150 V indeed results in
a reduced heating rate. However, even with the reduced
heating rate, the SOFC operating temperature of 700 °C
is reached in 540 s, which favours the developed SOFC
testing system for quick test cycles.
If the thicknesses of the electrolyte and cathode
layers are kept constant, and the anode layer is broad-
ened from 20 μm to 3000 μm, the built maximum tensile
stresses in the electrolyte are rather high (Figure 8b). In
all the investigated membrane designs, they exceed the
tensile strength of SDC. It appears that the relatively
large difference in the thermal-expansion coefficients
between Ni/SDC and SDC and the thick anode layer
substantially reduce the chances of the multi-layer
system to survive the heating. In order to reduce the
probability of failure in the anode-supported membrane,
the design heating rate must be further decreased.
4 CONCLUSIONS
A new SOFC system enabling quick testing cycles
was developed and successfully described using the
finite-element modelling. It was shown that Ni/SDC-
SDC-LSM membranes were subjected to temperature
gradients during the heating. These gradients were pri-
T. SKALAR et al.: OPTIMIZATION OF OPERATING CONDITIONS IN A LABORATORY SOFC TESTING DEVICE
Materiali in tehnologije / Materials and technology 49 (2015) 5, 731–738 737
Figure 8: Maximum stresses in the multi-layer SOFC membrane as a
function of: a) electrolyte thickness (danode = dcathode = 20 μm,
ømembrane = 11 mm, theating = 160 s at 220 V or theating = 300 s at 150
V) and b) anode thickness (delectrolyte = dcathode = 20 μm, ømembrane =
11 mm, theating = 160 s at 220 V or theating = 300 s at 150 V)
Slika 8: Maksimalna napetost v ve~slojni SOFC-membrani kot
funkcija: a) debeline elektrolita (danoda = dkatoda = 20 μm, ømembrana =
11 mm, tsegr. = 160 s pri 220 V oz. tsegr. = 300 s pri 150 V), b) debeline
anode (delektrolit = dkatoda = 20 μm, ømembrana = 11 mm, tsegr. = 160 s pri
220 V oz. tsegr. = 300 s pri 150 V)
marily caused by the heating mode in the testing system,
where hot gases circulated inside the anode and cathode
compartments and caused non-uniform membrane heat-
ing. The critical point with respect to a membrane failure
was recognized during the early stage of the system
heating when the temperature gradients in the membrane
were the largest. During the heating, the electrolyte and
cathode were always under a tensile stress, while the
anode was under a compressive stress. The maximum
thermally induced stress during the heating was observed
in the electrolyte layer adjacent to the anode/electrolyte
phase boundary. The absolute values of the induced
maximum stress increased with the increasing diameter
of the membrane. Additionally, the electrolyte-supported
membranes had greater chances of surviving rapid
temperature increases. In an electrolyte layer of  2 mm
or more, the SDC tensile stress was never exceeded even
if the system was heated with the maximum power. The
anode-supported membranes, in contrast, were much
more prone to cracking due to a relatively large diffe-
rence between the thermal-expansion coefficients of the
Ni/SDC anode and SDC electrolyte. Even a somewhat
reduced heating rate still induced the occurrence of over-
critical tensile stresses in the electrolyte. To reduce the
probability of an anode-supported membrane failure, the
heating power in the testing system should be further
lowered, at least during the early stage of the membrane
heating.
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T. SKALAR et al.: OPTIMIZATION OF OPERATING CONDITIONS IN A LABORATORY SOFC TESTING DEVICE
738 Materiali in tehnologije / Materials and technology 49 (2015) 5, 731–738
I. VOREL et al.: PRODUCTION OF SHAPED SEMI-PRODUCTS FROM AHS STEELS BY INTERNAL PRESSURE
739–744
PRODUCTION OF SHAPED SEMI-PRODUCTS FROM AHS
STEELS BY INTERNAL PRESSURE
IZDELAVA POLPROIZVODOV IZ AHS-JEKEL, OBLIKOVANIH Z
NOTRANJIM TLAKOM
Ivan Vorel1, Hana Jirková1, Bohuslav Ma{ek1, Petr Kurka2
1University of West Bohemia in Pilsen, Research Centre of Forming Technology, Univerzitni 22, 306 14 Pilsen, Czech Republic
2Fraunhofer Institute for Machine Tools and Forming Technology IWU, Reichenhainer Straße 88, 09126 Chemnitz, Germany
frost@vctt.zcu.cz
Prejem rokopisa – received: 2014-09-04; sprejem za objavo – accepted for publication: 2014-11-05
doi:10.17222/mit.2014.218
The constant evolution of industry brings demands for components of ever-higher quality, complex shapes and excellent
mechanical properties, with reduced operating costs and with improved reliability in service. Such demands can be met by
processing high-strength, low-alloyed steels using novel heat-treatment methods. One of such methods is the quenching and
partitioning (Q&P) process, which can lead to high strengths of about 2000 MPa at elongation levels of more than 10 %. The
Q&P process can be combined with unconventional procedures for making complex-shaped parts to manufacture components
with excellent mechanical properties and an intricate shape. This paper describes the sequence of internal high-pressure
forming, hot stamping and the Q&P process used for making a functional product with good mechanical properties and a
complex shape. The purpose of the experimental programme was to employ a production chain for processing hollow,
thin-walled stock. Using a step-by-step optimization, commercially usable products were obtained. The microstructure of the
products consisted primarily of martensite and a small fraction of bainite. Such a microstructure guarantees a high ultimate
strength. The strength level throughout the part was not less than 1950 MPa, combined with an excellent elongation of
approximately 15 %.
Keywords: Q&P process, hot stamping, internal high pressure forming
Stalen razvoj industrije zahteva komponente vedno bolj{ih kvalitet, kompleksnih oblik, odli~nih mehanskih lastnosti pri ni`jih
proizvodnih stro{kih in z ve~jo zanesljivostjo obratovanja. Take zahteve se lahko dose`ejo z izdelavo visokotrdnostnih
malolegiranih jekel z uporabo novih metod toplotne obdelave. Ena od takih metod je postopek kaljenja in delitve (Q&P), ki
omogo~a velike trdnosti okrog 2000 MPa pri ve~ kot 10-odstotnem raztezku. Postopek Q&P se lahko kombinira z
nekonvencionalnimi metodami izdelave delov s kompleksno obliko in za izdelavo komponent z odli~nimi mehanskimi
lastnostmi in komplicirano obliko. Predstavljeni ~lanek obravnava zaporedne operacije notranjega visokotla~nega
preoblikovanja, vro~ega stiskanja in Q&P-postopka za izdelavo funkcionalnih proizvodov z dobrimi mehanskimi lastnostmi in
zapleteno obliko. Namen eksperimentalnega dela je bil postaviti proizvodno verigo za izdelavo votlih tankostenskih kosov. Z
optimizacijo korak za korakom so bili izdelani komercialno uporabni proizvodi. Mikrostruktura proizvodov je bila sestavljena
predvsem iz martenzita in majhnega dele`a bainita. Taka mikrostruktura zagotavlja veliko natezno trdnost. Nivo trdnosti po
vsem kosu ni bil manj{i od 1950 MPa v kombinaciji z odli~nim raztezkom okrog 15 %.
Klju~ne besede: Q&P-proces, vro~e stiskanje, preoblikovanje z velikim notranjim tlakom
1 INTRODUCTION
Advanced high-strength steels (AHSS) find an ever-
increasing number of uses in a wide range of industrial
sectors. Thanks to their excellent mechanical properties,
they become the materials of choice for structural parts,
which can thus be more lightweight than, but equally
reliable and safe as, those made of conventional steels.
Due to the continuously evolving needs of industry, new
processing methods must always be sought to be com-
bined in a comprehensive production chain.
2 Q&P PROCESSING
The Q&P process is an advanced method of heat
treatment for high-strength steels.1 It relies on the
processes that take place in the steel between the Ms and
Mf temperatures and are related to carbon diffusion.
Carbon migrates from the super-saturated martensite to
the retained austenite, making the latter more stable. The
Q&P process comprises heating to the austenitizing
temperature and quenching to the quenching temperature
(QT), which lies between the Ms and Mf temperatures.
This produces a mixture of martensite and retained
austenite.2,3 The workpiece is then heated to, and held at,
the partitioning temperature (PT).4
With regard to the austenite stabilization, carbon
must not precipitate in the form of carbides at the PT. To
ensure this, aluminium, silicon or phosphorus should be
used as alloying additions. In order to strengthen the
solid solution, the steel can be alloyed with manganese
and chromium. These elements shift the pearlitic and
bainitic transformation curves in the transformation
diagrams towards longer times, thus reducing the critical
cooling rate. As silicon is present in the steel in order to
retard carbide precipitation, the optimum level of chro-
mium must be controlled, as chromium in the presence
Materiali in tehnologije / Materials and technology 49 (2015) 5, 739–744 739
UDK 621.77:621.777.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)739(2015)
of carbon acts as a carbide-former. The addition of man-
ganese as an austenite-forming element effectively pre-
vents any free ferrite from forming. Free ferrite can
impair the mechanical properties of the steel.5–7
2.1 Internal High Pressure Forming
This method can be used for manufacturing com-
plex-shaped, hollow products. It is based on forming a
hollow feedstock in a closed die using pressurized gas.
The final temperature of the product can be controlled by
changing the die-opening time. At the end of this
die-opening time, the product is removed from the die
and cooled in air or heat treated further.
3 EXPERIMENTAL PROGRAMME
The goal of the experimental programme was to form
hollow feedstock using internal, high-pressure gas and
then apply the Q&P process. The experimental material
was high-strength 42SiCr steel (Table 1). The initial mi-
crostructure of this steel consisted of ferrite and pearlite.
The ultimate strength was 981 MPa, the elongation
reached 30 % and the hardness was 295 HV10. Using the
JMatPro software, the important Ms and Mf temperatures
were found: 289 °C and 178 °C, respectively (Figure 1).
3.1 Experimental Production of Shaped Products
The hollow feedstock used for making the shaped
products was 380 mm in length and 48 mm in diameter.
Its wall thickness was 4 mm (Figure 2).
The sequence comprised heating to 915 °C and
holding for 25 min to guarantee full austenitizing. Then
the feedstock was transferred to a die at the ambient
temperature. After the die was closed, the feedstock was
formed by pressurized gas at an overpressure of 700 bar.
The working gas was nitrogen. The contact between the
workpiece and the die wall caused the workpiece to cool
down rapidly to a pre-defined temperature. The die wall
was at room temperature and was cooled by water. This
was achieved by selecting the appropriate die-opening
time between 5 s and 20 s. After the die-opening time,
the workpiece is removed. In order to explore the effect
of Q&P processing, the final step of the proposed
sequence, on the microstructure evolution and the related
mechanical properties, the workpieces were further
processed in three different ways (Figure 3). Some
workpieces were cooled in air after hot stamping in the
I. VOREL et al.: PRODUCTION OF SHAPED SEMI-PRODUCTS FROM AHS STEELS BY INTERNAL PRESSURE
740 Materiali in tehnologije / Materials and technology 49 (2015) 5, 739–744
Table 1: Chemical composition of steel 42SiCr, w/%
Tabela 1: Kemijska sestava jekla 42SiCr, w/%
C Si Mn Cr Mo Nb P S Ms/°C Mf /°C
0.42 2.6 0.59 1.33 0.03 0.03 0.01 0.01 289 178
Figure 2: Experimental stock and the resulting shape after hot
stamping
Slika 2: Preizkusni kos in dobljena oblika po vro~em stiskanju
Figure 1: TTT diagram of 42SiCr steel calculated using JMatPro soft-
ware
Slika 1: TTT-diagram jekla 42SiCr, izra~unan s programsko opremo
JMatPro
Figure 3: Experimental modes of heat treatment
Slika 3: Eksperimentalni na~ini toplotne obdelave
die. The product from this group was hot stamped in the
die for 15 s and was then removed and cooled in air (Q210
°C – Cair). The temperature at the moment of removal
from the die was 210 °C. Another group comprised
workpieces that were cooled in the die, then cooled in air
to the ambient temperature and were then reheated in a
furnace to 280 °C, where they were held for 25 min, and
then removed and cooled in air. The product from the
second group was cooled in the die for 15 s to a
temperature of 205 °C (Q150 °C – Cair – A280 °C). Yet another
group included workpieces that were removed from the
die and immediately placed in a furnace. The significant
product from this third group was cooled in the die for
15 s as well. Then it was placed in a furnace and
annealed at 260 °C for a time of 25 min (Q200 °C – P200 °C).
The temperature after removal from the die reached
200 °C (Table 2). They were therefore treated using the
Q&P process.
Table 2: Parameters of the experimental heat treatment
Tabela 2: Parametri eksperimentalne toplotne obdelave
Austeni-
tizing
°C
QT-tem-
perature
°C
PT-tem-
perature
°C
Temper-
ing time
min
(Q210 °C – Cair) 915 210 – –
(Q205 °C – Cair – A280 °C ) 915 205 280 25
(Q200 °C – P260 °C) 915 200 260 25
4 RESULTS AND DISCUSSION
Representative products were selected from these
groups of workpieces for metallographic examination
and for mechanical property measurements. The speci-
mens for the mechanical property measurement were
taken from the products representing the different pro-
cessing sequences. The specimen locations were chosen
to allow an evaluation of the effects of the production
sequence and of the position within the die upon the
resulting product properties. Representative specimens
were taken from the parts with which the workpiece was
held in the die, from the area of the steepest change in
cross-section, which is the point of the largest plastic
strain, and from the transitions between the different
diameters. A total of 14 test specimens were taken from
each product (Figure 4)
Due to the shape of the workpieces, miniature tension
test specimens with a gauge length of 5 mm and a cross-
section of 2 mm × 1.5 mm were used for measuring the
mechanical properties (Figure 5).
4.1 Mechanical Properties
The product that was only cooled in air after removal
from the die (Q210 °C – Cair) showed a very high ultimate
strength (Table 3). None of the 14 specimens exhibited
ultimate strengths of less than 2260 MPa. The elongation
level was higher than 10 % in all cases. The specimen
with the lowest ultimate strength, 2263 MPa, and an
elongation of 12 %, was T3/5, which had been taken
from the transition to the largest diameter. It also showed
the lowest hardness of all the specimens: 667 HV10. It is
likely that this part of the workpiece was not in full con-
tact with the die wall. No similar effect of the specimen
position was found in any other specimen from this pro-
duct.
Table 3: Results of tensile testing (proof stress Rp0.2, ultimate tensile
strength Rm and elongation A5 mm) and hardness testing (HV10) of the
semi-product (Q210 °C – Cair) upon air cooling
Tabela 3: Rezultati nateznih preizkusov (meja te~enja Rp0,2, natezna
trdnost Rm in raztezek A5 mm) in merjenja trdote (HV10) polproiz-
vodov (Q210 °C – Czrak) po ohlajanju na zraku
Rp0.2/MPa Rm/MPa A5 mm/% HV10
T3/2 1550 2324 14 709
T3/3 1534 2314 11 685
T3/5 1518 2263 12 667
T3/8 1641 2333 10 701
T3/12 1681 2345 10 738
T3/14 1652 2310 16 705
All the specimens from the product (Q205 °C – Cair –
A280 °C) exhibited ultimate strengths of more than 2110
MPa and A5 mm elongation levels of more than 14 %
(Table 4). A low ultimate strength was found in the T6/5
specimen: 2133 MPa. The T6/5 specimen also showed
A5 mm elongation of 18 %. This situation is similar to that
in the product that was removed from the die and cooled
I. VOREL et al.: PRODUCTION OF SHAPED SEMI-PRODUCTS FROM AHS STEELS BY INTERNAL PRESSURE
Materiali in tehnologije / Materials and technology 49 (2015) 5, 739–744 741
Figure 5: Shape of specimen for mechanical testing
Slika 5: Oblika vzorca za mehanske preizkuse
Figure 4: Selected locations of the T7 (Q200 °C – P260 °C) workpiece,
upon the Q&P process
Slika 4: Izbrana podro~ja kosa T7 (Q200 °C – P260 °C) pri Q&P-procesu
Table 4: Results of tensile testing of the semi-product (Q150 °C – Cair –
A280 °C)
Tabela 4: Rezultati nateznih preizkusov polproizvodov (Q150 °C –
Czrak – A280 °C)
Rp0.2/MPa Rm/MPa A5 mm/% HV10
T6/2 1822 2117 18 656
T6/3 1862 2160 17 653
T6/5 1838 2133 18 672
T6/8 1886 2189 18 637
T6/12 1771 2144 14 623
T6/14 1818 2196 19 671
in air. By contrast, the hardness in the corresponding
location of the latter product was the highest of all the
locations tested: 672 HV10.
The ultimate strength of the product (Q200 °C – P260 °C)
was approximately 250 MPa to 300 MPa lower than that
of the product that had been removed from the die and
merely cooled in air (Table 5). In addition, it is approxi-
mately 150 MPa to 200 MPa lower than in the product
that was removed from the die, cooled in air and then
reheated to 280 °C and held for 25 min. As in the pre-
vious cases, the lowest ultimate strength was found in
the T7/5 location. It was 1914 MPa. The corresponding
A5 mm elongation was 20 % and the hardness reached
600 HV10. The specimens in the remaining locations
exhibited strengths of approximately 2000 MPa. The
Q&P processing of the product improved its elongation.
It was between 17 % and 21 %.
Table 5: Results of tensile testing of the product (Q200 °C – P260 °C)
prepared using the Q&P process
Tabela 5: Rezultati nateznih preizkusov proizvodov (Q200 °C – P260
°C), izdelanih po Q&P-postopku
Rp0.2/MPa Rm/MPa A5 mm/% HV10
T7/2 1576 1978 17 623
T7/3 1490 1956 21 615
T7/5 1379 1914 20 600
T7/8 1600 2006 18 596
T7/12 1584 1952 18 597
T7/14 1626 1978 17 608
4.2 Metallographic and Fractographic Characteriza-
tion
With reference to the results of the tension test, the
fracture surfaces were examined. In all cases, ductile
fractures with dimples were found (Figure 6). No signs
of brittle fracture were detected, even in the product that
was cooled in air after quenching in the die.
The microstructure was examined in locations corres-
ponding to those of the samples taken for mechanical
testing. In almost all the specimens taken from the
product (Q200 °C – Cair) the microstructure consisted of
martensite and a small proportion of bainite (Figure 7).
Only those areas where the workpiece was not in full
I. VOREL et al.: PRODUCTION OF SHAPED SEMI-PRODUCTS FROM AHS STEELS BY INTERNAL PRESSURE
742 Materiali in tehnologije / Materials and technology 49 (2015) 5, 739–744
Figure 7: T3 (Q200 °C – Cair) product upon air cooling: a) martensite-
bainite microstructure (point T3/2), b) martesite-bainite with a small
fraction of ferrite (point T3/12)
Slika 7: T3 (Q200 °C – Czrak)-proizvod po ohlajanju na zraku: a) mar-
tenzitno-bainitna mikrostruktura (to~ka T3/2), b) martenzit-bainit z
majhnim dele`em ferita (to~ka T3/12)
Figure 6: Fracture surfaces: a) T3 (Q200 °C – Cair) product that was cooled in air (T3/3), b) T6 (Q205 °C – Cair – A280 °C) that was cooled in air and
then annealed in a furnace (T6/3), c) T7 (Q200 °C – P260 °C) product after Q&P process (T7/3)
Slika 6: Povr{ine prelomov: a) T3 (Q200 °C – Czrak)-proizvod, ohlajen na zraku (T3/3), b) T6 (Q205 °C – Czrak – A280 °C), ohlajen na zraku in potem
`arjen v pe~i (T6/3), c) T7 (Q200 °C – P260 °C) proizvod po Q&P-postopku (T7/3)
contact with the die wall contain some free ferrite, which
is probably due to the resulting insufficient cooling rate
An examination of the microstructure of the product
(Q205 °C – Cair – A280 °C) revealed tempered martensite and
bainite (Figure 8). The locations where the workpiece
was not in full contact with the die wall contained a
small amount of free ferrite.
The Q&P-processed product (Q200 °C – P260 °C) con-
tained a mixture of martensite and bainite (Figure 9). As
in the products mentioned above, the locations of partial
contact between the workpiece and the die wall con-
tained a small amount of free ferrite.
5 CONCLUSION
Using internal high-pressure forming, products with
complex shapes were obtained. A detailed inspection of
their surfaces revealed no distinct discontinuities. In the
product that was cooled in air after removal from the die,
the 15 s in the die caused quenching to 210 °C. Its mi-
crostructure was a mixture of martensite, bainite and a
small proportion of free ferrite. No location of the pro-
duct showed an ultimate strength of less than 2260 MPa.
The elongation was more than 10 %.
The product that was cooled in air after removal from
the die and then annealed in a furnace had a temperature
of 205 °C after the 15 s die-opening time. The annealing
of the product at 280 °C for 25 min produced a micro-
structure of tempered martensite, bainite and a small
amount of free ferrite. No location of the product, where
the mechanical properties were measured, showed an
ultimate strength of less than 2110 MPa. Annealing after
removal of the product from the die and subsequent cool-
ing in air caused the elongation level to increase from 10
% to 14 %.
Cooling in the die for 15 s, which is used as part of
the Q&P process, led to the quenching of the product to
200 °C. The subsequent reheating to the carbon-parti-
tioning temperature of 260 °C and holding for 25 min
resulted in a microstructure that consisted of tempered
martensite, bainite and free ferrite. The ultimate strength
of the product did not decrease below 1900 MPa at any
I. VOREL et al.: PRODUCTION OF SHAPED SEMI-PRODUCTS FROM AHS STEELS BY INTERNAL PRESSURE
Materiali in tehnologije / Materials and technology 49 (2015) 5, 739–744 743
Figure 9: T7 (Q200 °C – P260 °C) product upon Q&P process: a) mar-
tensite-bainite microstructure with a small fraction of ferrite (point
T7/2), b) martensite-bainite with a small fraction of ferrite (point
T7/12)
Slika 9: T7 (Q200 °C – P260 °C) proizvod po Q&P postopku: a) mar-
tenzitno-bainitna mikrostruktura z majhnim dele`em ferita (to~ka
T7/2), b) martenzit-bainit z majhnim dele`em ferita (to~ka T7/12)
Figure 8: T6 (Q205 °C – Cair – A280 °C) product upon air cooling and
annealing: a) tempered martensite-bainite microstructure (point T6/2),
b) tempered martensite-bainite microstructure (point T6/12)
Slika 8: T6 (Q205 °C – Czrak – A280 °C)-proizvod po ohlajanju na zraku
in popu{~anju: mikrostruktura: a) popu{~eni martenzit-bainit (to~ka
T6/2), b) mikrostruktura: popu{~eni martenzit-bainit (to~ka T6/12)
location. The Q&P processing led to an increase in the
elongation from 10 % to 18 %.
The measurement of mechanical properties by means
of tension testing did not identify any substantial effect
of the location, and thus no substantial effect of the strain
magnitude on the ultimate strength.
Acknowledgment
This paper includes results achieved in the project
CZ.1.05/3.1.00/14.0297 Technological Verification of
R&D Results II, individual activity Hollow Shafts for
Passenger Cars Produced by Heat Treatment with the
Integration of Q-P Process and the project GA^R
P107/12/P960 Influence of a Structure Modification on
Mechanical Properties of AHS Steel. The paper also
includes results from the project SGS-2013-028 Support
of Students Research Activities in Materials Engineering
Field.
6 REFERENCES
1 L. Kucerova et al., Comparison of Microstructures and Properties
Obtained After Different Heat Treatment Strategies of High Strength
Low Alloyed Steel, Journal of Iron and Steel Research International,
18 (2011), 427–431
2 J. Speer et al., Carbon Partitioning into Austenite after Martensite
Transformation, Acta Materialia, 51 (2003), 2611–2622,
doi:10.1016/S1359-6454(03)00059-4
3 J. Speer et al., The Quenching and Partitioning Process: Background
and Recent Progress, Materials Research, 8 (2005) 4, 417–423,
doi:10.1590/S1516-14392005000400010
4 T. Tsuchiyma et al., Quenching and partitioning treatment of a
low-carbon martensitic stainless steel, Materials Science and
Engineering A, 532 (2012), 585–592, doi:10.1016/j.msea.2011.
10.125
5 H. Jirkova, L. Kucerova, B. Masek, Effect of Quenching and Parti-
tioning Temperatures in the Q-P Process on the Properties of AHSS
with Various Amounts of Manganese and Silicon, Materials Science
Forum, 706–709 (2012), 2734–2739, doi:10.4028/www.scientific.
net/MSF.706-709.2734
6 L. Kucerova et al., Optimization of Q-P Process Parameters with Re-
gard to Final Microstructures and Properties, Annals of Daaam For
2009 & Proceedings of the 20th International Daaam Symposium, 20
(2009), 1035–1036
7 B. Masek et al., Improvement of Mechanical Properties of automo-
tive Components Using Hot Stamping with Integrated Q-P Process,
Journal of Iron and Steel Research International, 18 (2011) 1–2,
730–734
I. VOREL et al.: PRODUCTION OF SHAPED SEMI-PRODUCTS FROM AHS STEELS BY INTERNAL PRESSURE
744 Materiali in tehnologije / Materials and technology 49 (2015) 5, 739–744
T. ALAM et al.: COMPOSITE-MATERIAL PRINTED ANTENNA FOR A MULTI-STANDARD WIRELESS APPLICATION
745–749
COMPOSITE-MATERIAL PRINTED ANTENNA FOR A
MULTI-STANDARD WIRELESS APPLICATION
TISKANA ANTENA IZ KOMPOZITNEGA MATERIALA ZA
VE^STANDARDNO BREZ@I^NO UPORABO
Touhidul Alam1, Mohammad Rashed Iqbal Faruque1, Mohammad Tariqul Islam2,
Norbahiah Misran2
1Space Science Center (ANGKASA), Universiti Kebangsaan Malaysia, 43600UKM Bangi, Selangor, Malaysia
2Department of Electrical, Electronic and Systems Engineering, Universiti Kebangsan Malaysia, 43600UKM Bangi, Selangor, Malaysia
touhid13@yahoo.com
Prejem rokopisa – received: 2014-09-14; sprejem za objavo – accepted for publication: 2014-10-08
doi:10.17222/mit.2014.232
This paper presents a printed multi-standard wireless antenna fabricated from a cost effective composite material to cover GPS
(1575 MHz), GSM 1800, GSM 1900, WLAN (2400 MHz), LTE band 40 (2.3–2.4 GHz), WiMAX (3600 MHz) and WLAN
(5.1–5.35 GHz) frequency bands. The reported antenna is incorporated with two distinct monopole radiators with a meander-
line-type ground plane. The wireless mobile antenna can be conveniently simulated with the commercially available EM
simulation software (CST Microwave Studio) using the finite-difference time-domain (FDTD) method. A parametric analysis of
the antenna geometry is demonstrated and the specific absorption rate (SAR) is also analyzed with a human-head model.
Keywords: antenna, material, multiband, meander line, wireless communication, SAR
Ta ~lanek predstavlja tiskano ve~standardno brez`i~no anteno, izdelano iz cenovno ugodnega kompozitnega materiala, ki obsega
frekven~ne pasove GPS (1575 MHz), GSM 1800, GSM 1900, WLAN (2400 MHz), LTE pas 40 (2,3–2,4 GHz), WiMAX (3600
MHz) in WLAN (5,1–5,35 GHz). V predstavljeno anteno sta vgrajena dva razli~na monopolna sevalnika v obliki zanke na
osnovni ravnini. Brez`i~no mobilno anteno se lahko prikladno simulira s komercialno razpolo`ljivo EM simulacijsko program-
sko opremo (CST Microwave Studio) z uporabo metode domene s kon~no diferenco ~asa (FDTD). Prikazana je parametri~na
analiza geometrije antene in analizirana je bila hitrost specifi~ne absorpcije (SAR) na modelu ~love{ke glave.
Klju~ne besede: antena, material, ve~pasovno, linija zanke, brez`i~na komunikacija, SAR
1 INTRODUCTION
Multiband antenna design with a low-profile and
stable performance has recently been a significant issue
to the researchers. Therefore, research has been focused
on minimizing the physical size of individual parts of a
modern wireless system. In the most recent years,
printed planar antennas have been thought to be most
suitable for multiband wireless applications in view of
their unique features, for example, light weight, low costs,
simple fabrication, multi-frequency mode and stable per-
formances.1,2 Several types of multiband antenna have
been studied for GPS, GSM, WLAN, LTE band 40,
WiMAX and WLAN applications.
A compact dual-arm-structure mobile handset an-
tenna was designed for a multi-band wireless mobile
operation, covering DCS, PCS, UMTS and WLAN (2.4
GHz) frequency bands.3 The dimensions of the antenna
were 119 mm × 50 mm. An inverted L-shaped antenna
was presented for wireless communication, covering
DCS, PCS and IMT tri-bands.4 Chen et al.5 introduced a
modified T-shaped planar monopole antenna for DCS
1800, PCS 1900, UMTS and WLAN applications.
However, its dimensions were 65 mm × 40 mm.
This paper presents a multiband printed monopole
antenna for wireless communication, which can operate
within the existing wireless standards: GPS (1.565–1.585
GHz), GSM (1800, 1900), WiMAX (3.6 GHz), WLAN,
LTE band 40 (2.3–2.4 GHz) and WLAN (5.47–5.9
GHz). The overall volume of the proposed antenna is 30
mm × 60 mm, which is at least 26.5 % less than4, 50 %
less than3 and 40 % less than5, considering its length and
width. However, according to the IEEE and ICNIRP
guidelines the specific absorption rate of the proposed
antenna must be confirmed and the value should be less
than 1.6 W/kg in a 1 g averaging mass and 2 W/kg in a
10 g averaging mass of biological tissues.6,7 To comply
with this requirement, the SAR value of the proposed
antenna was analyzed and compared.
2 PROPOSED ANTENNA CONFIGURATION
The geometric configuration and the physical dimen-
sions of the proposed antenna are illustrated in Figure 1.
The proposed antenna consists of a meander-line radiator
with a defected ground plane. The antenna is printed on
an FR4-material (a relative permittivity of 4.6, a loss tan-
gent of 0.02) substrate with dimension of 30 mm × 60
mm × 1.6 mm. A 50  microstrip feeding line is con-
nected with an inverted S-shaped radiator. The specifi-
cations of the proposed antenna are listed in Table 1.
Materiali in tehnologije / Materials and technology 49 (2015) 5, 745–749 745
UDK 669.018.95:621.396.67 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)745(2015)
Table 1: Antenna-design specifications
Tabela 1: Specifikacije antene
Parameter code Value (mm) Parameter code Value (mm)
L 60 L9 10
W 30 L10 4.5
L1 17.5 L11 5
L2 7 Lf 22
L3 20 W1 2
L4 10 W2 5.87
L5 22 W3 6
L6 20 W4 13.125
L7 15 W5 25
L8 10
3 ANTENNA PERFORMANCE WITH AN
EPOXY-RESIN-POLYMER SUBSTRATE
The proposed planar microstrip patch antenna was
designed and analyzed using a finite-difference time
domain (FDTD) based CST Microwave Studio. The
designed antenna was fabricated on a recently available
1.6 mm thick, low-cost, durable, polymer-resin substrate
using an in-house printed-circuit-board (PCB) prototyp-
ing machine. The substrate material consists of an epoxy
matrix reinforced with woven glass. The composition of
epoxy resin and fiber glass varies in the thickness and is
direction dependent. One of the attractive properties of
polymer-resin composites is that they can be shaped and
reshaped repeatedly without losing their material pro-
perties.8 Due to the lower manufacturing cost, ease of
fabrication, design flexibility and market availability of
the proposed material, it has become popular for its use
as the substrate in patch-antenna designs. The material is
composed of 60 % of fiber glass and 40 % of epoxy
resin. Figure 2 shows the steps needed to construct an
epoxy-resin-polymer substrate (FR4).9
Moreover, a parametric study was performed for
several substrate materials, illustrated in Figure 3. The
dielectric properties of these materials are listed in Table
2. It is seen from Figure 3 that the FR4 substrate mate-
rial shows a better performance in terms of reflection
coefficient than the other materials listed in Table 2.
Table 2: Dielectric properties of substrate materials
Tabela 2: Dielektri~ne lastnosti materiala podlage
Substrate material Relative permittivity(r)
Dielectric loss
tangent
Glass (Pyrex) 4.82 0.0012
FR4 4.60 0.02
Taconic CER-10 10.00 0.0023
Teflon (PTFE) 2.10 0.01
4 RESULTS AND DISCUSSION
The prototype of the proposed antenna was fabricated
using the FR4 substrate material with a relative per-
mittivity of 4.4 and a loss tangent of 0.02. The simulated
and measured reflection coefficients of the proposed
T. ALAM et al.: COMPOSITE-MATERIAL PRINTED ANTENNA FOR A MULTI-STANDARD WIRELESS APPLICATION
746 Materiali in tehnologije / Materials and technology 49 (2015) 5, 745–749
Figure 3: Reflection coefficient of the proposed antenna for the mate-
rials listed in Table 2
Slika 3: Koeficient refleksije predlagane antene pri materialih iz
tabele 2
Figure 1: Geometry of the proposed antenna: a) top view, b) bottom
view
Slika 1: Geometrija predlagane antene: a) pogled od zgoraj, b) pogled
od spodaj
Figure 2: Flow chart of FR4-material construction9
Slika 2: Shema poteka priprave FR4-materiala9
antenna are presented in Figure 4. The measured and
simulated reflection coefficients are identical. It is clear-
ly seen that four operating bandwidths for the multi-band
operation are obtained. The measured bandwidths, deter-
mined with the reflection coefficient –10 dB, are 755
MHz (1.255–1.98 GHz), 674 MHz (2.16–2.89 GHz), 44
MHz (3.26–3.7 GHz) and 34 MHz (5.14–5.48 GHz)
which cover GPS (1575 MHz), GSM 1800, GSM 1900,
WLAN (2400 MHz), LTE band 40 (2.3–2.4 GHz),
WiMAX (3600 MHz) and WLAN (5.1–5.35 GHz).
The surface-current distributions at 1.8 GHz, 2.4
GHz and 3.6 GHz are demonstrated in Figure 5. From
this figure, it is seen that with the increasing frequency
the current flow increased in the S-shaped region. The
radiation patterns at 1.8 GHz, 2.4 GHz and 3.6 GHz are
shown in Figure 6. It is clear from this figure that the
radiation patterns for both the E-plane and H-plane are
omnidirectional at 3.6 GHz. But some directivity was
T. ALAM et al.: COMPOSITE-MATERIAL PRINTED ANTENNA FOR A MULTI-STANDARD WIRELESS APPLICATION
Materiali in tehnologije / Materials and technology 49 (2015) 5, 745–749 747
Figure 7: Radiation efficiency and gain of the proposed antenna
Slika 7: U~inkovitost sevanja in pridobitev predlagane antene
Figure 5: Surface-current distribution of the proposed antenna: a) 1.8
GHz, b) 2.4 GHz and c) 3.6 GHz
Slika 5: Razporeditev tokov na povr{ini predlagane antene: a) 1,8
GHz, b) 2,4 GHz in c) 3,6 GHz
Figure 6: Radiation patterns of the proposed antenna for the
frequencies of: a) 1.8 GHz, b) 2.4 GHz, c) 3.6 GHz
Slika 6: Vzorec sevanja predlagane antene pri frekvencah: a) 1,8 GHz,
b) 2,4 GHz in c) 3,6 GHz
Figure 4: Simulated and measured reflection coefficients of the
proposed antenna
Slika 4: Simuliran in izmerjen koeficient refleksije predlagane antene
shown for the E-plane at 1.8 GHz and 3.6 GHz. More-
over, the simulated radiation efficiency and peak gain of
the proposed antenna are presented in Figure 7. This
figure shows that the maximum radiation efficiency of
81.25 % and the minimum of 67.5 % were achieved. In
addition, the maximum peak gain of 3.72 dB was also
obtained with the proposed antenna. A brief comparison
of antenna performances is presented in Table 3.
Table 3: Comparison of antenna performances
Tabela 3: Primerjava zmogljivosti antene
Antenna Dimensions(mm × mm)
Resonances
(GHz)
Bandwidth
(MHz)
Max. gain
(dB)
3 119 × 50 1.71–2.48 770 5.4 at 2.15GHz
5 65 × 40 1.66–2.59,4.48–5.89 930, 1410 Not given
Proposed
antenna 60 × 30
1.255–1.98,
2.16–2.89,
3.26–3.7,
5.14–5.48
755, 674,
44, 34
3.72 at
4.10 GHz
5 SPECIFIC ABSORPTION RATE (SAR)
The analysis of the health risk of the electromagnetic
radiation of wireless devices is extensively in progress.
These devices paved the way for an extensive utilization
of mobile phones in modern society resulting in in-
creased concerns about the inimical radiation.10–12 There
are several factors that affect the electromagnetic inter-
action; a close proximity between the human head and a
wireless device is one of them. The specific absorption
rate is defined with the power absorbed per mass of bio-
logical tissue and it is expressed with the units of watts
per kilogram (W/kg). Currently, two international bodies
have developed guidelines for limiting the effects of the
electromagnetic radiation on human health. The EM
absorption limit specified in IEEE C95.1:20056 is 1.6
W/kg in a 1 g averaging mass and 2 W/kg in a 10 g ave-
raging mass of tissue, which is similar to the limit stated
in the International Commission on Non-Ionizing
Radiation Protection (ICNIRP) guideline.
Table 4: SAR values of the proposed antenna
Tabela 4: Vrednosti SAR za predlagano anteno
Frequency
(GHz)
SAR for 1 g
(W/kg)
SAR for
10 g (W/kg)
Absorbed
power
(rms) W
S11 with
human head
model (dB)
1.8 1.10 0.763 0.1179 –10.5
2.4 1.04 0.715 0.0957 –12.76
In designing antennas for wireless communication, it
is very important to analyze the SAR values of the
proposed antenna. In this research, a SAR analysis was
performed, with the reference power of the wireless de-
vice set to 500 mW. The distance between the head and
the handset was 4.5 mm. Moreover, the commercially
available CST Microwave Studio software head-phantom
model was adopted for this study. The head phantom is
made of two layers, one is the shell and the other is fluid.
The shell material specifications are:  = 5, μ = 1, tan  =
0.05; and the specifications of the fluid inside the shell
are:  = 42, μ = 1, el. conductivity of 0.99 S/m, fluid den-
sity of 1000 kg/m3. In addition, the SAR values at 1.8
GHz and 2.4 GHz are presented in Figure 8 and listed in
Table 4. The obtained 1 g SAR for the proposed antenna
at 1.8 GHz is 1.10 W/kg, which is about 27 % less than
the reference value.3
6 CONCLUSION
This paper presents a new printed planar antenna for
GPS, GSM, WLAN, LTE band 40, WiMAX and WLAN
wireless applications with a better antenna performance,
including impedance bandwidth, antenna gain, radiation
pattern and radiation efficiency obtained over operating
bands. The experimental results validated the simulated
ones. Moreover, the proposed antenna satisfies the re-
quirements relating to the specific absorption rate. The-
refore, the overall performance of the proposed antenna
makes it suitable for the wireless mobile application.
7 REFERENCES
1 W. C. Liu, C. M. Wu, Y. Dai, Design of triple-frequency micro-
strip-fed monopole antenna using defected ground structure, IEEE
Transactions on Antennas and Propagation, 59 (2011), 2457–2463,
doi:10.1109/TAP.2011.2152315
2 T. Alam, M. R. I. Faruque, M. T. Islam, Printed Circular Patch Wide-
band Antenna for Wireless Communication, Informacije MIDEM, 44
(2014), 212–217
3 D. Zhou, R. A. Abd-Alhameed, C. H. See, A. G. Alhaddad, P. S.
Excell, Compact wideband balanced antenna for mobile handsets,
IET Microwaves, Antennas & Propagation, 4 (2010), 600–608,
doi:10.1049/iet-map.2009.0153
4 Q. Rao, T. A. Denidni, New broadband dual-printed inverted
L-shaped monopole antenna for tri-band wireless applications,
Microwave and optical technology letters, 49 (2007), 278–280,
doi:10.1002/mop.22107
T. ALAM et al.: COMPOSITE-MATERIAL PRINTED ANTENNA FOR A MULTI-STANDARD WIRELESS APPLICATION
748 Materiali in tehnologije / Materials and technology 49 (2015) 5, 745–749
Figure 8: a) 1 g SAR at 1.8 GHz, b) 10 g SAR at 1.8 GHz, c) 1 g SAR
at 2.4 GHz and d) 10 g SAR at 2.4 GHz
Slika 8: a) 1 g SAR pri 1,8 GHz, b) 10 g SAR pri 1,8 GHz, c) 1 g
SAR pri 2,4 GHz in d) 10 g SAR pri 2,4 GHz
5 S. B. Chen, Y. C. Jiao, W. Wang, F. S. Zhang, Modified T-shaped
planar monopole antennas for multiband operation, IEEE Transac-
tions on Microwave Theory and Techniques, 54 (2006), 3267–3270,
doi:10.1109/TMTT.2006.877811
6 IEEE Standard for Safety Levels with Respect to Human Exposure to
Radio Frequency Electromagnetic Fields, 3 kHz to 300 GHz, IEEE
Std C95.1-2005 (Revision of IEEE Std C95.1-1991), 2006, pp.
0_1-238, doi:10.1109/IEEESTD.2006.99501
7 International Non-Ionizing Radiation Committee of the International
Radiation Protection Association, Guidelines on limits on exposure
to radio frequency electromagnetic fields in the frequency range from
100 kHz to 300 GHz, Health Physics, 54 (1988), 9
8 I. Yarovsky, E. Evans, Computer simulation of structure and proper-
ties of crosslinked polymers: application to epoxy resins, Polymer,
43 (2002), 963–969, doi:10.1016/S0032-3861(01)00634-6
9 M. Samsuzzaman, M. Islam, J. Mandeep, N. Misran, Printed wide-
slot antenna design with bandwidth and gain enhancement on
low-cost substrate, The Scientific World Journal, 2014 (2014), article
ID 804068, doi:10.1155/2014/804068
10 M. R. I. Faruque, M. T. Islam, N. Misran, Evaluation of specific ab-
sorption rate (SAR) reduction for PIFA antenna using metamaterials,
Frequenz, 64 (2010), 144–149, doi:10.1515/FREQ.2010.64.7-8.144
11 M. R. I. Faruque, M. T. Islam, N. Misran, Influence of SAR Reduc-
tion in Muscle Cube with Metamaterial Attachment, Informacije
Midem-Journal of Microelectronics, Electronic Components and
Materials, 41 (2011), 233–237
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for Electromagnetic Absorption Reduction in Human Head, Progress
in Electromagnetics Research, 141 (2013), 463–478, doi:10.2528/
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T. ALAM et al.: COMPOSITE-MATERIAL PRINTED ANTENNA FOR A MULTI-STANDARD WIRELESS APPLICATION
Materiali in tehnologije / Materials and technology 49 (2015) 5, 745–749 749

Ý. SUGÖZÜ et al.: FRICTION AND WEAR BEHAVIOUR OF ULEXITE AND CASHEW IN AUTOMOTIVE BRAKE PADS
751–758
FRICTION AND WEAR BEHAVIOUR OF ULEXITE AND CASHEW
IN AUTOMOTIVE BRAKE PADS
ODPORNOST PROTI TRENJU IN OBRABI AVTOMOBILSKIH
ZAVORNIH OBLOG Z ULEKSITOM IN PRAHOM IZ INDIJSKEGA
OREHA
Ýlker Sugözü1, Ýbrahim Mutlu2, Ahmet Keskin3
1Mersin University, Tarsus Technology Faculty, Mersin, Turkey
2Afyon Kocatepe University, Technology Faculty, Afyonkarahisar, Turkey
3Abant Izzet Baysal University, Bolu Vocational School, Bolu, Turkey
ibrahimmutlu@aku.edu.tr
Prejem rokopisa – received: 2014-09-14; sprejem za objavo – accepted for publication: 2014-10-15
doi:10.17222/mit.2014.228
In the experimental studies, ulexite and cashew were investigated as new materials in brake pads. A newly formulated brake-pad
material with five different ingredients was produced using ulexite. Tribological properties of the friction materials were
obtained using the brake-test equipment. The friction and wear characteristics of the samples in contact with a disk made of cast
iron were studied. The change in the friction coefficient, the temperature of the friction surface and the amount of the wear were
examined to assess the performance of these samples. In addition, microstructural characterizations of the braking pads were
carried out using scanning electron microscopy (SEM). The results showed that the friction materials containing ulexite and
cashew have an important effect on the friction stability and fade resistance. The strategy proposed in this paper can be
considered for the alternative friction materials where ulexite and cashew can be used as friction materials in the brake pads.
Keywords: brake pad, composite materials, friction coefficient, ulexite, cashew
Preiskovan je bil vpliv uleksita in prahu iz indijskega oreha kot novega materiala v zavornih oblogah. Novo zasnovani material
za zavorne obloge iz petih sestavin je bil izdelan z uporabo uleksita. Tribolo{ke lastnosti tornega materiala so bile preizku{ene
na napravi za preizkus zavor. Preu~evano je bilo trenje in zna~ilnosti obrabe vzorcev v stiku z diskom iz sive litine. Za oceno
zmogljivosti vzorcev so bile preiskane spremembe koeficienta trenja, izmerjena temperatura na povr{ini stika in koli~ina obrabe.
Z vrsti~nim elektronskim mikroskopom (SEM) so bile dodatno preiskane mikrostrukturne zna~ilnosti zavornih oblog. Rezultati
so pokazali, da imajo torni materiali, ki vsebujejo uleksit in prah indijskega oreha, pomemben vpliv na stabilnost trenja in
odpornost proti odpovedi zavor pri vi{jih temperaturah. V tem ~lanku predlagana usmeritev se lahko upo{teva kot nadomestni
torni material za zavorne obloge na osnovi uleksita in prahu iz indijskega oreha.
Klju~ne besede: zavorne obloge, kompozitni material, koeficient trenja, uleksit, prah iz indijskega oreha
1 INTRODUCTION
The friction element of an automotive brake system
is one of the most important composite materials and
generally consists of more than ten ingredients. This is
because the friction materials have to provide a steady
friction force, a reliable strength, and a good wear resis-
tance in a broad range of braking circumstances.1
Recently, some studies have shown that the most fatal
accidents on the roads happen because of failed brake
systems.2,3 The performance of a brake system in a
vehicle is mainly determined by the tribological charac-
teristics of the friction couple, composed of a gray-iron
disk (or drum) and friction materials.1,4
The most important property of the friction materials
is a high friction coefficient so that they remain stable
under high forces and, especially, at high temperatures.2,3
The function of the brakes is to change the kinetic
energy into the heat energy by absorbing it and releasing
it into the atmosphere. If the generated heat exceeds the
capacity of a brake, there is a decrease in the friction
coefficient of the brake pads. When brake pads are ex-
posed to high temperatures for long time, they are
damaged. This damage results in a decrease in the brake
performance, a high brake-pad wear or noise.3,5 A great
deal of effort has been made to develop friction materials
with optimized tribological characteristics regardless of
the type of braking conditions.3,4
Currently, more than 800 raw materials cited in the
literature are used to produce friction materials for the
commercial brakes.6,7 Most automotive friction materials
contain a phenolic-resin binder with additions of mineral
fibers, fillers, friction-modifying compounds, abrasives
and metallic particles to modify the heat-flow characte-
ristics.8,9 In a simplistic sense, fibers are included for
their friction properties, heat resistance and thermal
conductivity. The physical and chemical properties of a
resin affect the wear process and friction characteristics
of the friction material. In addition, a straight phenolic
resin was used for brake friction materials and various
modified resins are available to improve the compressi-
bility, thermal stability, damping capacity and mecha-
nical strength.10 They play an important function of
Materiali in tehnologije / Materials and technology 49 (2015) 5, 751–758 751
UDK 539.92:531.43 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)751(2015)
toughening and strengthening the binder, which is quite
brittle in its pure form.11 Numerous compositions of
friction materials have been developed, largely by em-
pirical testing.8,9 Several projects were conducted to
investigate the friction materials that could improve the
braking effectiveness, while also contributing to the
energy efficiency.12 In the literature, there are a lot of
studies about using new materials for the brake pads in
order to increase the braking performance.13–18
Raw and refined borates are used in ceramic glazes;
thin coatings are applied onto ceramics. Borates are
consumed mostly in the glass and fiberglass industry.
Borates help the glass formation, reduce the thermal
expansion and give resistance and durability to the
ceramics. An addition of borates reduces the material
stresses caused by the temperature gradients, thus
making it more resistant to breaking. Borates are also
used in borosilicate glasses. Borosilicate glass is not only
highly resistant to a chemical attack, but it also has a
very low coefficient of thermal expansion and, as a
result, a high resistance to thermal shock. This thermal-
shock resistance exceeds that of the ordinary glass by a
factor of three.19
Ulexite is a mineral that combines calcium, sodium,
water molecules and boron in a complicated arrangement
with the following formula: NaCaB5O9·8H2O. It consists
of thin crystals that act like optical fibers. On the surface,
ulexite takes the shape of soft-looking masses and is
often called "cotton ball". This form also occurs beneath
the surface in veins similar to chrysotile, in which the
crystal fibers run across the width of the veins.20
Cashew friction dust is made of CNSL (cashew
nut-shell liquid). Due to its friction particles it is used as
a stability agent in brake products; it has a resilient
nature, acting as a cushioning agent of the engaging pro-
perty of the lining. Further, it is easily decomposed on
the surface of a brake lining at various elevated tempe-
ratures controlling the wear and acting as a protective
element by prohibiting the generation of excessive tem-
perature. It also easily absorbs the heat and disperses it
accross the whole area of the friction material. It is
mainly used as one of the prime raw materials for
heavy-automobile non-asbestos and asbestos brake
linings. It is also used for clutch facings and disc
pads.21,22
In order to test the performance of ulexite in auto-
motive friction materials, ulexite is sifted after grinding
to obtain dust as raw boron products. In the case of
ulexite and cashew being used together, when the
amount of ulexite is increased, the amount of cashew is
decreased. Five kinds of samples with different ingre-
dients were designed. The tribological properties of these
samples were determined using brake-test equipment.
The friction tests were performed up to 400 °C. The
change in the friction coefficient, the temperature of
friction surface and the amount of wear were examined
to identify the performance of these samples. In addition,
microstructural characterizations of the braking pads
were carried out using scanning electron microscopy
(SEM). The results showed that the friction materials
containing ulexite have an important effect on the
friction stability and fade resistance. Due to the heat and
abrasion resistance and the ability to reduce the thermal
expansion of ulexite, the use of this material in brake
pads can be attractive.
2 MATERIALS AND METHODS
In this study, a new automotive-brake friction mate-
rial was developed using an addition of ulexite. The
influence of ulexite on the brake’s friction characteristics
was especially examined. The friction materials investi-
gated in this study were variations of a NAO (non-asbes-
tos organic)-type material containing different ingre-
dients including ulexite. Ulexite was obtained from
Balikesir, the Bigadiç mine of Etibank in Turkey. The
composition of the ulexite studied in this work is shown
in Table 1.
Table 1: Composition of ulexite studied in this work19
Tabela 1: Sestava uleksita, uporabljenega v tej {tudiji19
Element B2O3 CaO Na2O SiO2
w/% 24–38 18–24 2–7 4–13
Five different samples were produced. These samples
contained ulexite, phenolic resin, steel fiber, copper,
aluminium oxide, graphite, brass particles, cashew and
barite. The friction-coefficient and temperature values
were stored in a databank. Friction coefficient/tempe-
rature/time graphs and the mean coefficients of friction
were obtained to identify the friction characteristics. An
analytical balance was used to weigh the ingredients.
The friction-material samples were produced in the
conventional procedure for a dry formulation following
dry mixing, pre-forming and hot pressing. These ingre-
dients were mixed for 10 min using a commercial
blender. The final mixture was loaded into a one-inch
square (small samples) mold for pre-forming under a
pressure of 9.8 MPa. The pre-formed samples were put
Ý. SUGÖZÜ et al.: FRICTION AND WEAR BEHAVIOUR OF ULEXITE AND CASHEW IN AUTOMOTIVE BRAKE PADS
752 Materiali in tehnologije / Materials and technology 49 (2015) 5, 751–758
Table 2: Ingredients of the samples (w/%)
Tabela 2: Vsebnosti v vzorcih (w/%)
Sample code UC-4UCH-4
UC-8
UCH-8
UC-12
UCH-12
UC-16
UCH-16
UC-20
UCH-20
Phenolic resin 22 22 22 22 22
Steel fibers 15 15 15 15 15
Al2O3 3 3 3 3 3
Brass particles 5 5 5 5 5
Graphite 3 3 3 3 3
Barite 20 20 20 20 20
Cu particles 8 8 8 8 8
Cashew 4 8 12 16 20
Ulexite 20 16 12 8 4
Total 100 100 100 100 100
into a hot-pressing mold and pressed at a pressure of
14.7 MPa and at 180 °C for 15 min. During the hot-
-pressing process, pressure was released several times to
release the gases that evolved from the cross-linking
reaction (polycondensation) of the phenolic resin.
Detailed conditions for each manufacturing step can be
found in the author’s other study.23 The compositions of
the friction materials studied in this work are shown in
Table 2.
In order to define the friction coefficients of the
automotive brake pad under different temperatures, a test
device was designed and manufactured. Figure 1 shows
a schematic view of the brake tester used in this study.
Using a real brake-disc-type tester, the friction-
coefficient characteristics of a pad next to the disc made
of cast iron were investigated by changing the pad. The
test sample was mounted on the hydraulic pressure and
pressed against the flat surface of the rotating disc.
Before performing the friction-coefficient tests, the
surfaces of the test samples and the cast iron discs were
ground with 320-grit sandpaper. The normal load was
varied to achieve a constant friction force. The braking
tests were carried out at a pressure of 1.05 MPa, a
velocity of 6 ms–1 and at temperatures from 50 °C to 400
°C for 500 s. An electrical heater was used in order to
achieve the friction-surface temperature of 400 °C. The
temperature and friction-coefficient values were stored in
the databank. The tests were repeated three times for
each sample.
Friction coefficient/temperature/time graphs were
obtained to identify the effects of these variables. The
friction coefficient of the surface-material couple needs
to be high and stable. The friction coefficient was cal-
culated by measuring the normal and tangential
pressures throughout the test 500 s. It was expressed as
the mean value of the entire braking dependence during
the friction-coefficient test. The specific wear rate was
determined with the mass method following the Turkish
Standard (TS 555) and British Standard (BS AU142) and
calculated with the following Equation (1):
V = (m1 – m2)/(L fm 
) (1)
where V is the specific wear (mm3/MJ), m1 is the mass
of the brake pad before the testing (kg), m2 is the mass
of the brake pad after the testing (kg), L is the friction
distance calculated using the number of revolutions and
the radius of the disc (m), fm is the average friction force
(N) and 
 is the density of the brake pad (kg/mm3).24,25
3 RESULTS AND DISCUSSION
3.1 Effect of the temperature on the friction perfor-
mance
In the present study, 10 samples were used. These
samples contained copper particles, phenolic resin,
Al2O3, steel fibers, brass particles, graphite, barite,
ulexite and cashew (Table 2). Only half of the samples
(with the UCH indices) were heat treated for 4 h at a
temperature of 180 °C. The remaining 5 samples were
untreated (with the UC indices). These samples included
4–20 % ulexite and 20–4 % cashew, respectively, and the
mean friction coefficient ranged from 0.394 to 0.454.
When the coefficient of friction (μ) was studied it
varied significantly during the initial stage of the testing.
This can be attributed to the fact that the size of the
contact area increased and the friction layer was
developed on the surface. In order to determine the
variations in the friction coefficient and the temperature
of the friction surface with the testing time of the
samples, the tests were performed at applied tempera-
tures from 50 °C to 400 °C (Figures 2 to 6). As seen
from these figures, the friction coefficients show
different features depending on the content.
Generally, the friction-coefficient value gradually
increases for all the samples until the 100 s and then
gradually decreases after the 400 s (Figures 2 to 6). The
reason of the increasing friction coefficient is the contact
of the metallic material with the disc surface. The wear
of the ingredients of the metallic materials is tested.
Ý. SUGÖZÜ et al.: FRICTION AND WEAR BEHAVIOUR OF ULEXITE AND CASHEW IN AUTOMOTIVE BRAKE PADS
Materiali in tehnologije / Materials and technology 49 (2015) 5, 751–758 753
Figure 2: Change in the friction coefficient versus the temperature as
a function of time for samples UC-4 and UCH-4
Slika 2: Spreminjanje koeficienta trenja v odvisnosti od temperature
in ~asa pri vzorcih UC-4 in UCH-4
Figure 1: Schematic view of the brake tester
Slika 1: Shematski prikaz preizku{evalnika zavor
Therefore, these metallic materials detach from the
surface of the brake pad and the friction coefficient
decreases. This phenomenon continues until new friction
surfaces appear. A rapid increase in the coefficient of
friction may lead to a rapid increase in the temperature
of the friction surface.
It was found that the friction coefficient decreases
with the increasing testing temperature. Generally, the
friction coefficient decreased between 350 °C and 400
°C due to the softening of the phenolic resin (Figures 2
to 6). As a result, fading occurred during the braking
action. Furthermore, with the increasing temperatures,
the ingredients in the brake pads affected each other due
to a faster diffusion. This phenomenon is called the
thermal fade.4 Therefore, one can say that the samples
remained unchanged only up to the temperature of 400
°C. There was an increase in μ before the 200 s, followed
by a decrease just after the 350 s when μ was almost
constant for the UCH-4, UCH-16 and UCH-20 samples
(Figures 2, 5 and 6). This degradation is somewhat slow,
having slight fluctuations. By the 300 s, as a result of the
friction, a temperature of 300–350 °C was achieved on
the friction surface.
As seen from these figures, it is understood that all
the heat-treated samples (UCH) have higher friction-
coefficient values and more stable friction coefficients
than the others. The differences between the friction
coefficients of the UC samples and the UCH samples are
obvious especially at the beginning and after the 400 s.
The difference at the initial state shows a good fit of the
heat-treated samples to the surface. At the states after the
400 s the heat-treated samples have a lower temperature
then the others at this time (Figures 3 to 6). Thus, the
friction coefficients of the UC samples decreased due to
the increase in the temperature.26 It is observed that the
friction-coefficient and temperature values of the
non-heat-treated samples quickly change and fluctuate.
As a result, it can be said that heat treatment makes the
friction coefficients stable and increases the friction-
coefficient values.
Nonetheless, it should be noted that a good stability
of μ is achieved using the samples under the working
condition considered. These results are consistent with
the behavior of the friction coefficients of all the sam-
ples. Therefore, if a μ value of 0.39–0.45 is desired,
additions of both ulexite and cashew can be used in the
brake pads in the amounts of 4–20 %. Furthermore, if the
stability of μ is desired and μ has higher values, the
UCH-20 sample is suggested as the best material for the
brake pads when compared to the others. Some middle
vibrations and noise were observed during the testing
with the friction assessment and screening test (FAST).
This vibration was typically observed at the beginning of
the test, before a stable friction layer was developed.
Ý. SUGÖZÜ et al.: FRICTION AND WEAR BEHAVIOUR OF ULEXITE AND CASHEW IN AUTOMOTIVE BRAKE PADS
754 Materiali in tehnologije / Materials and technology 49 (2015) 5, 751–758
Figure 6: Change in the friction coefficient versus the temperature as
a function of time for samples UC-20 and UCH-20
Slika 6: Spreminjanje koeficienta trenja v odvisnosti od temperature
in ~asa pri vzorcih UC-20 in UCH-20
Figure 5: Change in the friction coefficient versus the temperature as
a function of time for samples UC-16 and UCH-16
Slika 5: Spreminjanje koeficienta trenja v odvisnosti od temperature
in ~asa pri vzorcih UC-16 in UCH-16
Figure 4: Change in the friction coefficient versus the temperature as
a function of time for samples UC-12 and UCH-12
Slika 4: Spreminjanje koeficienta trenja v odvisnosti od temperature
in ~asa pri vzorcih UC-12 in UCH-12
Figure 3: Change in the friction coefficient versus the temperature as
a function of time for samples UC-8 and UCH-8
Slika 3: Spreminjanje koeficienta trenja v odvisnosti od temperature
in ~asa pri vzorcih UC-8 in UCH-8
3.2 Microstructural characterization of friction surfa-
ces
Apparently, friction layers are formed by the wear
particles generated during the friction. The chemistry
and structure of a friction layer depend on the bulk
materials (lining and disc), the testing conditions and the
environment. The role of the friction layer may vary
depending on its characteristics.27,28 The SEM micro-
graphs of the braking-pad surfaces after the braking test
are shown in Figures 7 and 8. The friction surfaces of
the samples were characterized using SEM (LEO 1430
VP). The sample surfaces for the SEM observations were
always coated with carbon. There are micro-voids on the
surfaces of almost all the samples. Micro-voids consist
of the metallic particles detached during the friction.
It is seen from Figures 7 and 8 (UC-8 and UCH-4)
that larger micro-voids occurred in the samples due to
detached metallic particles. As seen in these figures,
some particles are detached from the body causing
micro-voids. The micro-voids on the surfaces of the
samples can be classified as smaller or bigger in size.
The bigger micro-voids were formed due to the pitting of
the metallic particles during the friction. The worn
metallic particles imply that they actively participated in
the friction during the braking test. It is known that if the
coherent surface of a metal component is bigger, the
friction and wear will be increased. In addition to
micro-voids, there are some micro-cracks on the surfa-
ces. It is also observed that Al2O3 particles are distri-
buted homogeneously, therefore, contributing to the
effectiveness of the friction surfaces.
Several characteristic features can be observed on the
friction surfaces of the pads. White spots are seen on
Figure 7 (UC-8 and UC-20). Dark areas can also be seen
in Figures 7 and 8 (UC-8 and UCH-4). A more loose
contact on the trailing edge facilitates the access of air
and an uneven wear related to a higher oxidation (a
burn-off) of the phenolic resin. This is due to the dis-
tribution of the friction force over a pad surface.28 When
the pads heat up during the braking, the resin tends to
expand at very high temperatures and turn into glassy
carbon. Carbonized resins weaken the matrix and acce-
lerate the pad wear (Figure 7, UC-8 and UC-16).28,29 The
glassy phase loses its support and is torn off from the
surface by the shear force.30 Figures 7 and 8 relating to
UC-12, UC-16, UC-20, UCH-8 and UCH-12 (from the
SEM) show thick friction layers developed on the
surfaces of the pads.
In this particular case, the friction layer covering a
friction surface diminishes the abrasive effect of the
glassy phase by eliminating the sharp edge of the glass
and smoothing the friction surface. Hard glassy particles
typically act as an abrasive element and scratch off the
Ý. SUGÖZÜ et al.: FRICTION AND WEAR BEHAVIOUR OF ULEXITE AND CASHEW IN AUTOMOTIVE BRAKE PADS
Materiali in tehnologije / Materials and technology 49 (2015) 5, 751–758 755
Figure 8: SEM micrographs of brake-pad samples with the UCH code
Slika 8: SEM-posnetki vzorcev zavornih oblog z oznako UCH
Figure 7: SEM micrographs of brake-pad samples with the UC code
Slika 7: SEM-posnetki vzorcev zavornih oblog z oznako UC
cast-iron-disc counter face and the material adhering to
it.28 Apparently, the carbonaceous matrix was formed of
graphite, coke and degraded phenolic resin. The ingre-
dient with the plastic-deformation capability developed a
flake-like feature after the friction experiment (Figures 7
and 8; UC-16, UC-20, UCH-8 and UCH-12).
All matters were homogeneously distributed in the
matrix and, therefore, very few micro-voids were ob-
served in the structures (Figures 7 and 8; UC-4, UC-8,
UC-12, UCH-8 and UCH-20). The friction process is
characterized by the development of friction debris. Such
debris adheres to the friction surface and forms a friction
layer easily visible when examining a sample surface
after the testing (FAST) (Figures 7 and 8; UC-16,
UC-20, UCH-4 and UCH-8). A systematic analysis of
the surfaces of the composite materials indicated that the
friction process dominantly occurred on the friction
layer, which eventually covered the top of the bulk.
Well-developed friction layers on the friction surfaces as
well as their morphologies are easily visible. Detailed
views of the friction surfaces including the information
about the friction layers are shown in Figures 7 and 8.
Diffusion occurred in the Cu particles located in the
friction layers. Brighter areas marked as Cu in Figures 7
and 8 (UC-4, UC-8, UC-16 and UCH-12) represent the
regions where Cu interacted with the friction layers.31
3.3 Wear behaviour
Table 3 gives the mean coefficients of friction, the
standard deviations of coefficients of friction, densities,
hardness values and the specific wear rates of the tested
samples. As can be seen from Table 3, the friction
coefficients are in the appropriate category according to
the Turkish Standard TS 555 and British Standard BS
AU142. Also, the standard deviations are very small,
which means that the materials have stable friction cha-
racteristics.
Table 3: Typical characteristics of the brake pads used in this study
Tabela 3: Zna~ilnosti zavornih oblog, uporabljenih v tej {tudiji
Sample
code
Mean
coeffi-
cient of
friction
Standard
deviation
Density
(g cm–3)
Hardness
(Brinell)
Specific wear
(g mm–2)
UC–4 0.394 0.0182 1.935 18.2 0.26 × 10–6
UC–8 0.401 0.0172 1.842 16.2 0.27 × 10–6
UC–12 0.420 0.0192 1.765 15.9 0.34 × 10–6
UC–16 0.427 0.0216 1.704 14.7 0.37 × 10–6
UC–20 0.437 0.0223 1.680 13.9 0.38 × 10–6
UCH–4 0.410 0.0091 1.834 23.2 0.21 × 10–6
UCH–8 0.418 0.0127 1.782 20.3 0.22 × 10–6
UCH–12 0.430 0.0127 1.717 18.9 0.29 × 10–6
UCH–16 0.442 0.0140 1.652 18.2 0.33 × 10–6
UCH–20 0.454 0.0134 1.588 15.3 0.35 × 10–6
In the present study, no direct proportionality was
found among the density, the hardness and the wear
resistance due to the complexity of the composite
structures. However, the highest friction coefficient was
obtained for the UCH-20 sample. This sample includes
4 % of ulexite and 20 % of cashew and it was heat
treated. The lowest friction coefficient was obtained for
the UC-4 sample. This sample includes 20 % of ulexite
and 4 % of cashew and it was heat untreated. In addition,
less wear was observed on the heat-treated samples,
which included 4–20 % of ulexite and 20–4 % of
cashew, and their mean values of the friction coefficients
ranged from 0.41–0.45 (Table 3). Besides, the minimum
wear was obtained for the UC-4 and UCH-4 samples,
including the heat-treated and untreated ones. It is seen
from Table 3 that the heat-treated samples generally
have a better wear resistance than the untreated samples.
It is also seen from Table 3 that heat treatment
facilitated a more homogeneous braking-pad structure.
In addition, a more stabilised friction coefficient was
obtained due to heat treatment. It is assumed that some
detached materials left the samples and then structural
variations were obtained. Furthermore, with the in-
creasing temperature the ingredients in the braking pad
affected each other due to a faster diffusion. Therefore, it
can be said that heat treatment is essential for braking
pads. It is also known that the hardness of a sample in-
creases and the density decreases due to heat treatment,
while the specific wear ratio changes (Table 3). In the
present work, heat treatment facilitated a better perfor-
mance of all the samples compared with the untreated
ones. These results are consistent with the earlier
works.32–35
As can be seen from Table 3, the friction coefficients
achieved for samples UCH-20 and UCH-16 are approxi-
mately 0.45, which is considered to be very good when
compared to the coefficients of friction achieved with the
current brake pads. The UC-4 sample has the highest
density and hardness among the untreated samples. It
includes 20 % of ulexite and 4 % of cashew; the specific
gravity of ulexite is bigger than that of cashew. A higher
amount of ulexite causes a higher density. But the heat-
treated UCH-4 sample with the same ingredients has the
highest hardness among the all samples.
It is well known that the friction coefficient is usually
associated with an increase in the wear. In the present
work, all the samples confirmed this assumption (Table
3). But the UCH-16 and UCH-20 samples have better
values than the other samples. Therefore, one can say
that only samples UCH-16 and UCH-20 are preferred for
this kind of brake pads as they provide better mechanical
and microstructural properties than the other samples.
It can be seen from the results that there is no direct
correlation between the wear resistance and the hardness.
However, this is an unavoidable reality of these features
that affect each other. As can be seen from the table, a
high-hardness pad has a high wear resistance and lower
wear due to the amounts of the constituent components
of the samples. The components made of the pad mate-
Ý. SUGÖZÜ et al.: FRICTION AND WEAR BEHAVIOUR OF ULEXITE AND CASHEW IN AUTOMOTIVE BRAKE PADS
756 Materiali in tehnologije / Materials and technology 49 (2015) 5, 751–758
rials comprised of compounds of different contents can
have a hard structure and, consequently, a higher hard-
ness. But the amount of the resin component included in
the materials is not sufficient as the particle adhesion
surface is decreased and there is a quick separation of the
particles that make up the main structure at low strains,
while the friction can cause a wear that is higher than
expected.
The heating process generally affects the microstruc-
ture of a brake pad and, accordingly, the hardness
increases. It can be seen from Table 3 that the friction
coefficient of the UC-4 sample is 0.394 and its wear rate
is 0.26 · 10–6 and, after the heat treatment, these values
are 0.410 and 0.21 · 10–6, respectively. The situation is
similar for the other samples. Thus, after the heat treat-
ment, the coefficients of friction of the samples of the
same ingredients increase and their wear rates decrease.
Heat treatment has a positive impact on both the friction
coefficient and specific wear.
The mean coefficient of friction, hardness, and wear
rates of the heat-treated (UCH) and untreated (UC)
samples showing parallelism to each other exhibit the
characteristics of the material. The sample with a high
mean coefficient of friction also has a high wear rate. In
this case, the rubbing of the ruptures of large particles
from the brake pad due to the strain has a significant
impact. Also, since the pad’s hardness is lower than the
hardness of the disk, the wear naturally increases during
the process of creating a high coefficient of friction.
After the heat treatment, the densities of the samples
decrease. During the heat treatment at a low evaporation
temperature, volatile materials move away from the body
and form micro-scale porosity in the body. This situation
causes a decrease in the density, and the heat produced
during the friction (rubbing) is removed from the body
by the micro-pores, leading to a lower friction-surface
temperature and a stable friction performance.36
4 CONCLUSIONS
In this study, the effect of ulexite and cashew
amounts on the friction and wear behavior of a brake pad
used in the automotive industry is experimentally ana-
lyzed. As a result of the experiments, the structure and
chemical composition of the friction layer generated on
the friction surface differ significantly from those of the
bulk. It is apparent that no simple relationship exists
between the composition of the friction layer and the
bulk-material formulation. Heat treatment facilitated a
more homogeneous structure and, hence, microstructural
variations were minimized during the braking action. On
the other hand, heat treatment increased the hardness of
the samples and also decreased the density.
The highest friction coefficient was obtained for the
heat-treated samples. Smaller values were obtained for
the untreated samples. The UCH-4 sample had a more
stable friction coefficient during the FAST than the other
samples. As the UCH-4 and UCH-8 samples had lower
friction coefficients, their wear ratios and standard devi-
ations were also considerably lower. Out of 10 samples,
UC-20, UCH-12, UCH-16 and UCH-20 exhibited better
friction and wear properties. Therefore, these samples
can be suggested as brake-pad materials.
In the present work, the standard deviation was
within the acceptable range for all the samples. No direct
proportionality among the mean coefficients of friction,
the standard deviation and the wear resistance was found
due to the complexity of the composite structures. Some
micro-voids and micro-cracks were observed on the
worn surfaces.
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2000, 181–187
3 A. Kurt, M. Boz, Mater. Design, 26 (2005), 717–721, doi:10.1016/
j.matdes.2004.09.006
4 M. G. Jacko, S. K. Lee, Kirk–Othmer Encyclopedia of Chemical
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758 Materiali in tehnologije / Materials and technology 49 (2015) 5, 751–758
I. GUNES, M. OZCATAL: DIFFUSION KINETICS AND CHARACTERIZATION OF BORIDED AISI H10 STEEL
759–763
DIFFUSION KINETICS AND CHARACTERIZATION OF BORIDED
AISI H10 STEEL
KINETIKA DIFUZIJE IN KARAKTERIZACIJA BORIRANEGA
JEKLA AISI H10
Ibrahim Gunes, Melih Ozcatal
Department of Metallurgical and Materials Engineering, Faculty of Technology, Afyon Kocatepe University, 03200 Afyonkarahisar, Turkey
igunes@aku.edu.tr
Prejem rokopisa – received: 2014-09-22; sprejem za objavo – accepted for publication: 2014-10-17
doi:10.17222/mit.2014.238
In this study, case properties and diffusion kinetics of the AISI H10 steel borided in Ekabor-II powder were investigated by
conducting a series of experiments at temperatures of (1123, 1173 and 1223) K for (2, 4 and 6) h. The boride layer was
characterized with light microscopy, X-ray diffraction technique and micro-Vickers hardness tester. The X-ray diffraction
analysis of the boride layers on the surfaces of the steels revealed the existence of FeB, Fe2B, CrB, Cr2B and MoB compounds.
Depending on the chemical compositions of the substrates and boriding time, the boride-layer thickness on the surface of the
steel ranged from 12.86 μm to 63.72 μm. The hardness of the boride compounds formed on the surfaces of the steels ranged
from 1648 HV0.05 to 1964 HV0.05, whereas the Vickers-hardness value of the untreated steel was 306 HV0.05. The activation
energy (Q) of the borided steel was 160.594 kJ/mol. The growth kinetics of the boride layer formed on the AISI H10 steel and
its thickness were also investigated.
Keywords: AISI H10, boride layer, microhardness, kinetics, activation energy
V tej {tudiji so bile preiskovane lastnosti in kinetika difuzije v jeklu AISI H10, boriranem v prahu Ekabor-II (2, 4 in 6) h na
temperaturah (1123, 1173 in 1223) K. Boridni sloj je bil okarakteriziran s svetlobno mikroskopijo, z rentgensko difrakcijo in
merjenjem mikrotrdote po Vickersu. Rentgenska difrakcija boridnega sloja je odkrila prisotnost naslednjih spojin: FeB, Fe2B,
CrB, Cr2B in MoB.
Odvisno od kemijske sestave podlage in ~asa boriranja je bila debelina sloja borida na povr{ini jekla med 12,86 μm in 63,72 μm.
Trdota boridov, nastalih na povr{ini jekla, je bila med 1648 HV0,05 in 1964 HV0,05, medtem ko je bila Vickersova trdota
neobdelanega jekla 306 HV0,05. Aktivacijska energija (Q) pri boriranju jekla je bila 160,594 kJ/mol. Preiskovana je bila tudi
kinetika rasti in debelina boridnega sloja na jeklu AISI H10.
Klju~ne besede: AISI H10, sloj borida, mikrotrdota, kinetika, aktivacijska energija
1 INTRODUCTION
Boriding, or boronizing, is a thermochemical sur-
face-hardening process that can be applied to a wide
variety of ferrous, nonferrous and cermet materials. The
process involves heating a well-cleaned material in the
range of 973 K to 1373 K, preferably for 1 h to 12 h, in
contact with a boronaceous solid powder (boronizing
compound), paste, liquid, plasma, gaseous or electro-
chemical medium.1–7
Boron atoms, due to their relatively small size (an
atomic radius of 0.09 nm) and very mobile nature can
diffuse easily into ferrous alloys (an atomic radius of
0.124 nm) forming FeB and Fe2B intermetallic, non-
oxide, ceramic borides. The diffusion of B into steel
results in the formation of iron borides (FeB and Fe2B)
and the thickness of the boride layer is determined by the
temperature and time of the treatment.8–12
The basic advantage of boriding is that boride has a
high melting point and high hardness at elevated tem-
peratures and, consequently, researches of the boriding
of transition metals have been accelerated in the recent
decade, particularly for the applications in the production
of cutting tools and heavy gears, and in the automotive,
casting, textile, food-processing, packaging and ceramic
industries where huge friction-dependent energy losses
and intensive corrosion and wear occur.13–15 In this study,
the AISI H10 steel was borided considering these
advantages of boriding. The characterization and growth
diffusion of the obtained boride layers were calculated.
The main objective of this study was to investigate the
diffusion kinetics and the effects of the processing
parameters such as the temperature, the time and the
chemical composition on the boride layers formed on the
AISI H10 steel after powder-pack boriding at different
processing temperatures and times.
2 EXPERIMENTAL DETAILS
2.1 Boriding and characterization
The AISI H10 steel essentially contained mass
fractions 0.32 % C, 3.15 % Cr, 2.90 % Mo, 0.65 % V
and 0.40 % Mn. The test specimens were cut into
dimensions of Ø 25 mm × 8 mm, ground up to 1200 G
and polished using a diamond solution. The boriding
heat treatment was carried out in a solid medium con-
Materiali in tehnologije / Materials and technology 49 (2015) 5, 759–763 759
UDK 669.058:669.781:544.034 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)759(2015)
taining an Ekabor-II powder mixture placed in an
electrical-resistance furnace operating at temperatures of
(1123, 1173 and 1223) K for (2, 4 and 6) h under
atmospheric pressure. The microstructures of the po-
lished and etched cross-sections of the specimens were
observed under a Nikon MA100 light microscope. The
presence of borides formed in the coating layer was
confirmed with X-ray diffraction equipment (Shimadzu
XRD 6000) using Cu-K radiation. The thickness of
borides was measured with a digital thickness-measuring
instrument attached to the light microscope (Nikon
MA100). The hardness measurements of the boride layer
on each steel and of the untreated steel substrate were
made on the cross-sections using a Shimadzu HMV-2
Vickers indenter with a 50 g load.
2.2 Evaluation of the activation energy of boron
diffusion
In order to study the diffusion mechanism, borided
AISI H10 steel was used for this purpose. It is assumed
that boride layers grow parabolically in the direction of
the diffusion flux and perpendicular to the substrate
surface. So, the time dependence of the boride-layer
thickness can be described with Equation (1):
x Dt2 = (1)
where x is the depth of the boride layer (mm), t is the
boriding time (s) and D is the boron diffusion coeffi-
cient through the boride layer. It is a well-known fact
that the main factor limiting the growth of a layer is the
diffusion of boron into the substrate. It is possible to
argue that the relationship between the growth-rate con-
stant D, the activation energy Q, and the temperature T
in Kelvin, can be expressed as an Arrhenius equation
(Equation (2)):
D D
Q
RT
= −⎛⎝
⎜ ⎞
⎠
⎟
0 exp (2)
where D0 is a pre-exponential constant, Q is the activa-
tion energy (J/mol), T is the absolute temperature in
Kelvin and R is the ideal gas constant (J/(mol K)).
The activation energy for the boron diffusion in a
boride layer is determined with the slope obtained in the
plot of ln D vs. 1/T, using Equation (3):
ln lnD D
Q
RT
= −0 (3)
3 RESULTS AND DISCUSSION
3.1 Characterization of boride coatings
Light micrographs of the cross-sections of the
borided AISI H10 steel at the temperatures of 1123 K
and 1223 K for 2 h and 6 h are shown in Figure 1. As
can be seen the borides formed on the AISI H10 sub-
strate have a saw-tooth morphology. It was found that the
coating/matrix interface and the matrix can be signifi-
cantly distinguished and the boride layer has a columnar
structure. Depending on the chemical compositions of
the substrates, the boriding time and temperature, the
boride-layer thickness on the surface of the AISI H10
steel ranged from 12.86 μm to 63.72 μm in Figure 2.
Figure 3 gives the XRD patterns obtained at the
surface of the borided AISI H10 steel at 1123 K and
1223 K for the treatment times of 2 h and 6 h. The XRD
patterns show that the boride layer consists of borides
such as MB and M2B (M = metal: Fe, Cr). The XRD
results showed that the boride layers formed on the H10
steel contained FeB, Fe2B, CrB, Cr2B and MoB phases
(Figures 3a to 3d.)
Microhardness measurements were carried out along
a line from the surface to the interior in order to see the
variations in the hardness of the boride layer, the tran-
sition zone and the matrix, respectively. The microhard-
ness of the boride layers was measured at 10 different
locations at the same distance from the surface and the
I. GUNES, M. OZCATAL: DIFFUSION KINETICS AND CHARACTERIZATION OF BORIDED AISI H10 STEEL
760 Materiali in tehnologije / Materials and technology 49 (2015) 5, 759–763
Figure 2: Thickness values of boride layers with respect to boriding
time and temperature
Slika 2: Debelina borirane plasti glede na ~as in temperaturo boriranja
Figure 1: Cross-sections of borided AISI H10 steel: a) 1123 K – 2 h,
b) 1123 K – 6 h, c) 1223 K – 2 h, d) 1223 K – 6 h
Slika 1: Pre~ni prerez boriranega jekla AISI H10: a) 1123 K – 2 h, b)
1123 K – 6 h, c) 1223 K – 2 h, d) 1223 K – 6 h
average value was taken as the hardness. The results of
the microhardness measurements carried out on the
cross-sections, along the line from the surface to the
interior are presented in Figure 4. The hardness of the
boride layer formed on the AISI H10 steel varied bet-
ween 1648 HV0.05 and 1964 HV0.05. On the other hand,
the Vickers hardness value for the untreated AISI H10
steel was 306 HV0.05. When the hardness of the boride
layer is compared with the matrix, the boride-layer
hardness is approximately five times larger than that of
the matrix.
3.2 Kinetics
In this study, the effects of the processing tempe-
rature and boriding time on the growth kinetics of a
boride layer were also investigated. The kinetic para-
meters such as the processing temperature and time must
be known for the control of the boriding treatment.
Figure 5 shows the time dependence of the squared
value of the boride-layer thickness at increasing tempe-
ratures. This evolution followed the parabolic growth law
where the diffusion of boron atoms is a thermally
activated phenomenon. The growth-rate constant D at
each boriding temperature can be easily calculated with
Equation (1).
I. GUNES, M. OZCATAL: DIFFUSION KINETICS AND CHARACTERIZATION OF BORIDED AISI H10 STEEL
Materiali in tehnologije / Materials and technology 49 (2015) 5, 759–763 761
Figure 3: X-ray diffraction patterns of borided AISI H10 steel: a) 1123 K – 2 h, b) 1123 K – 6 h, c) 1223 K – 2 h, d) 1223 K – 6 h
Slika 3: Rentgenogrami boriranega jekla AISI H10: a) 1123 K – 2 h, b) 1123 K – 6 h, c) 1223 K – 2 h, d) 1223 K – 6 h
Figure 5: Time dependence of squared boride-layer thickness at
increasing temperatures
Slika 5: ^asovna odvisnost kvadrata debeline boridne plasti pri na-
ra{~ajo~i temperaturi
Figure 4: Hardness variation with the depth for the borided AISI H10
steel
Slika 4: Spreminjanje trdote po globini boriranega jekla AISI H10
As a result, the calculated growth-rate constants at
three temperatures, (1123, 1173 and 1223) K, are (5.58 ×
10–10, 9.67 × 10–10, and 1.18 × 10–9) cm2 s–1 for the bo-
rided AISI H10 steel. Table 1 lists the calculated values
of the growth constant for each boriding temperature.
Figure 6 describes the temperature dependence of
the growth-rate constant. The plot of ln D as a function
of the reciprocal temperature exhibits a linear relation-
ship according to the Arrhenius equation. The boron
activation energy can be easily obtained from the slope
of the straight line presented in Figure 6. The value of
the boron activation energy was then determined as
160.594 kJ/mol for the borided AISI H10 steel.
Table 2 compares the obtained value of energy
(160.594 kJ/mol) with the data found in the literature. It
is seen that the reported values of the boron activation
energy depended on the chemical composition of the
substrate and the used boriding method. The calculated
value in this study is comparable with the values re-
ported in the literature as seen in Table 2.16–21
Table 2: Comparison of the activation-energy values for diffusion of
boron with respect to different boriding media and substrates
Tabela 2: Primerjava aktivacijske energije za difuzijo bora glede na
razli~ne medije in snovi pri boriranju
Steel Temperaturerange (K)
Boriding
medium
Activation
energy
(kJ/mol)
Refe-
rences
AISI 8620 973–1073 Plasma paste 99.77 16
AISI W1
AISI 52100
1123–1323
1123–1223
Pack
Pack
177.8
269
17
18
AISI 1035 1073–1273 Salt bath 227.5 19
AISI H13 1073–1223 Powder 186.2 20
AISI H13 1073–1223 Salt bath 244 21
AISI H10 1123–1223 Pack 160.594 Presentstudy
A contour diagram describing the evolution of the
boride-layer thickness as a function of the boriding para-
meters (the time and the temperature) is shown in Figure
7. This contour diagram can be used for two purposes:
(1) to predict the coating-layer thickness with respect to
the processing parameters, namely, the time and tempe-
rature; (2) to determine the value of the processing time
and temperature for obtaining a predetermined coating-
layer thickness.22 The boride layer increased with an
increase in the boriding time and temperature for the
borided AISI H10 steel.
4 CONCLUSIONS
The following conclusions may be derived from the
present study.
• The boride types formed on the surface of the hot-
work tool steel have columnar structures.
• The boride-layer thickness obtained on the surface of
the AISI H10 steel was 12.86–63.72 μm, depending
on the chemical compositions of the substrates.
• The multiphase boride coatings that were thermo-
chemically grown on the AISI H10 steel consisted of
the FeB, Fe2B, CrB, Cr2B and MoB phases.
• The surface hardness of the borided steel was in the
range of 1648–1964 HV0.05, while for the untreated
steel it was 306 HV0.05.
I. GUNES, M. OZCATAL: DIFFUSION KINETICS AND CHARACTERIZATION OF BORIDED AISI H10 STEEL
762 Materiali in tehnologije / Materials and technology 49 (2015) 5, 759–763
Figure 7: Contour diagram describing the evolution of boride-layer
thickness as a function of boriding parameters
Slika 7: Konturni diagram, ki opisuje debelino plasti borida v odvis-
nosti od parametrov pri boriranju
Figure 6: Temperature dependence of the growth-rate constant
according to the Arrhenius equation
Slika 6: Temperaturna odvisnost konstante rasti, skladna z Arrheni-
usovo ena~bo
Table 1: Growth-rate constant (D) as the function of boriding
temperature
Tabela 1: Konstanta hitrosti rasti (D) v odvisnosti od temperature
boriranja
Material
Growth-rate constant (cm2 s–1)
Temperature
1123 K 1173 K 1223 K
AISI H10 5.58 × 10–10 9.67 × 10–10 1.18 × 10–9
• The boron activation energy was estimated to be
160.594 kJ/mol for the borided AISI H10 steel.
• A contour diagram relating the boride-layer thickness
to the boriding parameters (the time and the tempe-
rature) was proposed. It can be used as a simple tool
to select the optimum boride layer for a practical
utilization of this kind of material.
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I. GUNES, M. OZCATAL: DIFFUSION KINETICS AND CHARACTERIZATION OF BORIDED AISI H10 STEEL
Materiali in tehnologije / Materials and technology 49 (2015) 5, 759–763 763

M. BOY et al.: APPLICATION OF THE TAGUCHI METHOD TO OPTIMIZE THE CUTTING CONDITIONS ...
765–772
APPLICATION OF THE TAGUCHI METHOD TO OPTIMIZE THE
CUTTING CONDITIONS IN HARD TURNING OF A RING BORE
UPORABA TAGUCHI-JEVE METODE ZA OPTIMIZACIJO TRDEGA
STRU@ENJA ROBA IZVRTINE
Mehmet Boy1, Ibrahim Ciftci2, Mustafa Gunay3, Feridun Ozhan4
1Karabuk Vocational College, Karabuk University, Karabuk, Turkey
2Department of Manufacturing Engineering, Karabuk University, Karabuk, Turkey
3Department of Mechanical Engineering, Karabuk University, Karabuk, Turkey
4ORS Bearing, Ankara, Turkey
mboy@karabuk.edu.tr
Prejem rokopisa – received: 2014-09-29; sprejem za objavo – accepted for publication: 2014-11-28
doi:10.17222/mit.2014.246
This paper is focused on optimizing the cutting conditions for the surface roughness, inner-diameter error and roundness
obtained in hard turning of an inner ring bore. The hard-turning experiments were conducted on hardened and tempered AISI
52100 bearing rings using the L9 orthogonal array on a CNC lathe. The cutting speed, feed rate and number of the machined
part were selected as control factors. The optimum cutting conditions were determined using the signal-to-noise (S/N) ratio. S/N
ratios were calculated using the lower-the-better approach. An analysis of variance (ANOVA) was also employed to determine
the level of the effect of the control factors for the surface roughness, inner-diameter error and roundness. The statistical
analysis showed that the feed rate was the most significant factor for the surface roughness while the cutting speed was the most
significant factor for the roundness and inner-diameter error. Finally, the optimum cutting conditions were further confirmed
with confirmation tests.
Keywords: diameter error, hard turning, roundness, surface roughness, Taguchi method
^lanek je osredinjen na optimiranje razmer pri rezanju glede na hrapavost povr{ine, napake notranjega premera in okroglosti pri
trdem stru`enju notranjega roba izvrtine. Preizkusi trdega stru`enja so bili izvr{eni pri kaljenem in popu{~enem AISI 52100
obro~u le`aja z uporabo L9 ortogonalne matrike na CNC-stru`nici. Hitrost rezanja, hitrost podajanja in {tevilo obdelanih kosov
so bili izbrani kot kontrolni faktorji. Optimalne razmere rezanja so bile izbrane z uporabo razmerja signal-hrup (S/N).
S/N-razmerja so bila izra~unana s pribli`kom ~im manj tem bolj{e. Analiza variance (ANOVA) je bila tudi uporabljena za
dolo~itev vpliva kontrolnih faktorjev na hrapavost povr{ine, napak notranjega premera in okroglosti. Kon~no so bili optimalni
parametri potrjeni s potrditvenimi preizkusi.
Klju~ne besede: napaka premera, trdo stru`enje, okroglost, hrapavost povr{ine, Taguchi-jeva metoda
1 INTRODUCTION
Bearings are the most important components of
rotating mechanical systems. The highly precise surface
finish, material properties, cleanliness, dimensions and
tolerances of today’s bearings contribute significantly to
the product performance. Traditionally, bearing rings are
manufactured from the soft AISI 52100 steel through
forging, cold ring rolling and machining processes to
their near-net shapes. Then, they are hardened and tem-
pered to obtain the required hardness. The final geometry
and surface quality are achieved through grinding and
super-finishing processes. The hard-turning process has
been increasingly used as an alternative to grinding
operations. The hard-turning process has the capacity of
machining the hardened components such as roller
bearings, gears, shafts, cams, axles requiring a low
surface roughness and close dimensional and form tole-
rances. Hard turning, compared to grinding operations,
has many advantages such as a high material-removal
rate, lower production costs, a shorter cycle time, an
improved overall product quality, and a lower environ-
mental impact. However, in the area of precision hard
turning, due to the demands for a geometric accuracy of
a few micrometers, its application is limited by the un-
certainties with respect to the part quality and process
reliability.1–4
For a turned part, the diameter error, the surface
roughness and the roundness are the three most import-
ant quality characteristics.There are many causes for
dimensional and geometric accuracy errors in hard
turning. The roundness error is considered to be one of
the important geometrical errors of cylindrical compo-
nents because it has a negative effect on the accuracy and
other important factors such as the mating mechanical
components and wear of rotating elements.5,6
The machine-tool rigidity and clamping system
influence the form accuracy and dimensional tolerance in
precision hard turning. Possible machining errors may be
caused due to different types of clamping, the number of
clamping jaws, clamping forces, the hydraulic pressure
of a clamping system, the clamping rigidity and accu-
racy. The effects of clamping techniques on the defor-
mations of thin-walled rings have been investigated by
many researchers.7–13
Materiali in tehnologije / Materials and technology 49 (2015) 5, 765–772 765
UDK 621.941:519.233.4 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)765(2015)
Geometric errors mostly result from the inaccuracy
of tooling parts, such as inserts, tool holders and clamp-
ing devices.12 An investigation suggested that after
replacing an insert, the repeatability errors at the tip of
the insert can reach up to several microns, and the dis-
placement of a tool tip under a cutting load can also
reach several microns. Precision hard turning requires
highly precise machine tools to eliminate the form and
dimensional errors induced by the moving components
of a machine tool such as the spindle, the slide bed and
the tail stock.7–13
The surface roughness is one of the most important
requirements in a machining process, as it is considered
to be an indicator of the product quality. It indicates the
irregularities of a surface texture. Achieving a desired
surface quality is critical for the functional behavior of a
part. The surface roughness influences the performance
of mechanical parts and their production costs because it
affects the factors such as friction, ease of holding the
lubricant, electrical and thermal conductivities. A relati-
vely better surface finish may involve a higher cost of
manufacturing. The surface roughness and the roundness
error are affected by several factors including the cutt-
ing-tool geometry, the cutting speed, the feed rate, the
microstructure of a workpiece and the rigidity of a ma-
chine tool.14,15
In recent years, much work has been performed using
various statistical and experimental techniques. These
studies were mostly based on the design and analysis of
experimental methods to determine the effects of the
cutting parameters on the diameter error, the roundness
and the surface roughness of cylindrical parts. Islam5
investigated the effects of the cooling method, the blank
size and the work material on the dimensional accuracy
and surface finish of various turning parts using the tra-
ditional analysis, Pareto ANOVA and the Taguchi
method. The results showed that the work material had
the greatest effect on the diameter error and surface
roughness, while the major contributor to the circularity
was the blank size. Rafai and Islam6 experimentally and
analytically investigated the effects of the cutting para-
meters on the dimensional accuracy and surface finish in
dry turning. The Taguchi method and the Pareto ANOVA
analysis were used to determine the effects of the major
controllable machining parameters and, subsequently, to
find their optimum combination. They reported that
while the surface roughness could be optimized through
a proper selection of the feed rate, optimization of the
diameter error and circularity was difficult due to the
complex interactions between the input parameters.
Brinksmeier et al.8 analyzed the basic mechanisms caus-
ing a ring distortion in soft machining in order to derive
the strategies for its minimization. They concluded that
distortion was significantly influenced by the turning
before the heat treatment in the manufacture of bearing
rings and that the elastic ring deformation under the
clamping force led to variations in the depth of cut and
polygonal form. Brinksmeier and Sölter10 developed a
new method to predict the shape deviation of machined
workpieces with a complex geometry. This new method
combined the experimental results for the machining
workpieces with simple geometries using finite-element
simulations. This was achieved by making use of the
known source stresses in simple parts for which the
approach was validated. Finally, the method was applied
to predict the shape deviation of a ground linear rail
guide. Zhou et al.12 analyzed the possible error-driver
factors and error sources in the precision hard turning
and a strategy was proposed for an on-line compensation
of dimensional errors, based on the tool-wear monitoring
and a thermal-expansion prediction. They found that
within a certain range of flank wear, the developed me-
thod could significantly improve the geometric accuracy.
Sölter et al.13 analyzed the strategies for reducing the
roundness deviations of turned rings using a three-jaw
chuck. The simulations and measurements agreed reas-
onably well and showed that the minimum out-of-round-
ness strongly depends on the interaction of the angular
shift between the external and internal clamping with the
deviation between the segment-jaw diameter and the
inner-ring diameter.
The purpose of this study was to obtain the optimum
cutting conditions (the cutting speed, the feed rate, the
number of the machined part) for minimizing the surface
roughness, diameter error and roundness when hard
turning an inner-bearing ring bore. An L9 orthogonal
array was used in the design of the experiment. The
Taguchi method and an analysis of variance (ANOVA)
were also used to achieve this purpose. Furthermore,
ANOVA was used to determine the statistical signifi-
cance of the cutting conditions. In addition, the optimum
cutting conditions were aimed to be further confirmed by
confirmation tests.
2 MATERIAL AND METHOD
2.1 Material
The hard turning tests were performed on the inner
bearing ring bores of AISI 52100 steel, the composition
of which is given in Table 1. The geometry of the
inner-bearing ring is given in Figure 1. The rings were
manufactured from soft AISI 52100 through forging,
spheroidising, cold ring rolling, soft machining,
hardening and, finally, tempering processes to their
near-net shapes. The rings were hardened and tempered
to 58–62 HRc.
Table 1: Chemical composition of AISI 52100 steel (w/%)
Tabela 1: Kemijska sestava jekla AISI 52100 (w/%)
C Si Mn P Ni Cr Mo Cu Al Fe
0.99 0.24 0.36 0.016 0.06 1.43 0.02 0.10 0.017 Balance
2.2 Cutting inserts and the tool holder
The cutting tools used were commercial-grade
TiN-coated low-content CBN inserts with the geometry
of DCGW 11T304 produced by Sandvik Coromant.
766 Materiali in tehnologije / Materials and technology 49 (2015) 5, 765–772
M. BOY et al.: APPLICATION OF THE TAGUCHI METHOD TO OPTIMIZE THE CUTTING CONDITIONS ...
These inserts were recommended for machining hard-
ened steel by Sandvik and had the 7015 Sandvik desig-
nation. The inserts were rigidly mounted on a right-
hand-style tool holder designated by ISO as A25T
SDUCR 11 for hard turning inner ring bores.
2.3 Cutting conditions and measurements
The hard-turning tests were carried out on a twin-
spindle CNC chucker equipped with a pair of highly
precise collet chucks. Three cutting speeds (120, 140 and
160) m/min and three feed rates (0.04, 0.06 and 0.08)
mm/r with a fixed depth of cut of 50 μm were tested. At
each condition, a total of 12 rings were hard turned. Sur-
face-roughness and roundness measurements were
carried out on the hard-turned bore surfaces using a
Taylor Hobson precision instrument. The inner diameter
of an inner ring bore was measured with Diatest UD-1.
The surface-roughness, roundness and diameter-error
measurements were carried out on the 4th, 8th and 12th
hard-turned rings and the readings were averaged.
2.4 Design and analysis of the experiments
Optimization of machining processes is essential for
achieving a high responsiveness to production, which
provides a preliminary basis for the survival in today’s
dynamic market conditions. The Taguchi method is
widely used in engineering analyses and it is a powerful
design tool. This method dramatically reduces the
number of tests using orthogonal arrays and minimizes
the effects of the factors that cannot be controlled. Fur-
thermore, it provides a simple, efficient and systematic
approach to specifying the optimum cutting parameters
for a manufacturing process.16,17
The Taguchi method converts objective function
values to the signal-to-noise (S/N) ratio to measure the
performance characteristics at the levels of the control
factors against these factors. The S/N ratio is defined as
the desired signal ratio for an undesired random-noise
value and shows the quality characteristics of the expe-
rimental data.16,17 Usually, there are three categories of
the performance characteristic in an analysis of the S/N
ratio, i.e., the lower-the-better, the higher-the-better and
the nominal-the-better characteristics. The goal of this
study was to minimize the surface roughness, the inner-
diameter error and the roundness. Therefore, the lower-
the-better quality characteristic was used as shown in
Equation (1):
S
N
lg
n
Yi
i( )
= − ⋅ ⋅
⎡
⎣⎢
⎤
⎦⎥=
∑10 1 2
1
0
(1)
where Yi is the observed data during the experiment and
n is the number of experiments.16,17
In this study, the cutting speed (V), the feed rate (f)
and the number of the machined part (MP) were selected
as the control factors and their levels were determined as
shown in Table 2. The L9 orthogonal array of the Ta-
guchi method allowed us to determine the optimum cutt-
ing conditions and analyze the effects of the machining
parameters.
Table 2: Control factors and their levels
Tabela 2: Kontrolni faktorji in njihovi nivoji
Control factors Level 1 Level 2 Level 3
Cutting speed, V/(m/min) 100 120 140
Feed rate, f/(mm/r) 0.04 0.06 0.08
Machined part number, MP 4 8 12
ANOVA was used to find the relative contribution of
the cutting conditions in controlling the response of a
turning operation. The optimum combination of the
M. BOY et al.: APPLICATION OF THE TAGUCHI METHOD TO OPTIMIZE THE CUTTING CONDITIONS ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 765–772 767
Table 3: Experimental results and corresponding S/N ratios
Tabela 3: Rezultati preizkusov in odgovarjajo~a razmerja S/N
Run
Control factors Experimental results Signal-to-noise ratio (S/N)
V f MP IDe Ra R R Ra IDe
1 120 0.04 4 0.016 0.164 3.33 –10.448 15.703 35.917
2 120 0.06 8 0.012 0.293 3.04 –9.657 10.662 38.416
3 120 0.08 12 0.011 0.459 3.21 –10.130 6.763 39.172
4 140 0.04 12 0.010 0.176 2.63 –8.389 15.089 40.000
5 140 0.06 4 0.008 0.178 3.02 –9.600 14.991 41.938
6 140 0.08 8 0.005 0.363 2.35 –7.421 8.801 46.020
7 160 0.04 8 0.010 0.271 4.23 –12.526 11.340 40.000
8 160 0.06 12 0.013 0.342 4.6 -13.255 9.319 37.721
9 160 0.08 4 0.007 0.363 4.51 -13.083 8.801 43.098
Figure 1: Geometry of the inner ring used for hard-turning tests
Slika 1: Geometrija notranjega obro~a, uporabljenega pri preizkusih
trdega stru`enja
control factors for the surface roughness, diameter error
and roundness was determined with the ANOVA table
and S/N ratios. Lastly, confirmation tests were done
using the optimum cutting conditions found with the
Taguchi optimization method and thereby the validity of
the optimization was tested.
3 RESULTS AND DISCUSSION
The surface roughness, the inner-diameter error and
the roundness achieved during the hard turning of inner-
bearing rings were measured after the experiments
performed according to the L9 orthogonal array. The
lowest values of the surface roughness, the inner-dia-
meter error and the roundness significantly improve the
quality of the bearing assembly. For this reason, the
lower-the-better quality characteristic was used for the
calculation of the S/N ratio. The experimental results and
S/N ratios are given in Table 3.
3.1 Evaluation of the surface roughness for an inner
ring bore
The interaction plots for the surface-roughness para-
meters are given in Figure 2. The minimum value of the
surface roughness (Ra) was obtained for V = 120 m/min,
f = 0.04 mm/r and the 4th machined part. It can be seen
from Figure 2 that the Ra value increases as the feed rate
is increased from 0.04 mm/r to 0.08 mm/r when hard
turning the inner-bearing ring bores. With the raise in the
feed-rate value from 0.04 mm/r to 0.08 mm/r, a signifi-
cant increase is observed in the Ra value. No significant
relation between the surface roughness and the number
of the machined ring bore is seen.
The S/N ratios of the Ra data obtained from the
experimental results, used to determine the optimum
level of each variable, were calculated according to
Equation (1). Figure 3 illustrates the plots of the S/N
ratios that were calculated for Ra in the hard turning of
the inner ring bores. The S/N ratios of the factors for
each level are shown in Table 4. Different values () of
the S/N ratio between the maximum and the minimum
are also shown in Table 4. Therefore, by considering the
S/N ratios in Table 4 and Figure 3, the optimum cutting
conditions for the surface roughness were V2 (V = 140
m/min), f1 (f = 0.04 mm/r) and MP1 (the 4th machined
part). The smallest surface roughness and its S/N ratio
that can be obtained with these levels were calculated
using Equations (2) and (3). The Ra value and its S/N
ratio were determined as 0.131 μm and 17.621 dB, res-
pectively:
    G G G G G= + − + − + −( ) ( ) ( )V f MP0 0 0 (2)
Ra cal
G
=
−
10

 (3)
In these equations, G is the S/N ratio calculated at
the optimum levels (dB), G is the average S/N ratio of
all the variables (dB), V0 , f0 , MP0 are the mean S/N
ratios when factors V, f and MP are at the optimum levels
(dB), and Ra cal is the calculated surface-roughness (Ra)
value.
Table 4: Response table of S/N ratios for surface roughness
Tabela 4: Tabela odzivov razmerja S/N glede na hrapavost povr{ine
Factor
S/N ratio
Rank
Level 1 Level 2 Level 3 
V 11.043 12.961 9.821 3.14 2
f 14.044 11.658 8.122 5.922 1
MP 13.166 10.268 10.391 2.897 3
The analysis of variance (ANOVA) was used to find
which design parameters significantly affect the surface
roughness. This analysis was carried out at the 95 % con-
fidence level. The ANOVA results for the surface rough-
ness (Ra) are shown in Table 5. This table also shows the
degree of freedom (DF), the sum of squares (SS), the
mean square (MS), the F-values (F), the probability (P)
and the percentage-contribution ratio (PCR) of each
factor. A low P-value of 0.05 shows a statistically signi-
ficant level of the source for the corresponding response.
The F-ratios and their PCR were taken into consideration
to identify the significance levels of the variables. Table
5 indicates that the most effective variable for the Ra
M. BOY et al.: APPLICATION OF THE TAGUCHI METHOD TO OPTIMIZE THE CUTTING CONDITIONS ...
768 Materiali in tehnologije / Materials and technology 49 (2015) 5, 765–772
Figure 2: Interaction plots for surface roughness
Slika 2: Interakcijski diagrami za hrapavost povr{ine
Figure 3: Main-effect plot for surface roughness
Slika 3: Diagram glavnega u~inka na hrapavost povr{ine
value is the feed rate with the PCR of 62.84 %. It is well
known that the theoretical geometrical-surface roughness
is primarily a function of the feed for a given nose radius
and that it changes with the square of the feed-rate value.
It is once again shown that the feed rate has an important
effect on the surface roughness in hard turning.The other
variables having an effect on Ra are the number of the
machined part with a PCR of 19.00 % and the cutting
speed with a PCR of 17.74%.
Table 5: ANOVA for S/N ratios for surface roughness
Tabela 5: ANOVA za S/N-razmerje glede na hrapavost povr{ine
Source DF SS MS F P PCR
V 2 15.0351 7.5176 42.42 0.023 17.74
f 2 53.265 26.6325 150.28 0.007 62.84
MP 2 16.1066 8.0533 45.44 0.022 19.00
Error 2 0.3544 0.1772 0.42
Total 8 84.7612 100
3.2 Evaluation of the roundness for an inner ring bore
The interaction plots for the roundness and the
cutting conditions are given in Figure 4. The minimum
value of the roundness was obtained at V = 140 m/min
and f = 0.08 mm/r for the 8th machined part. The round-
ness is generally affected by the cutting parameters such
as the cutting speed, the feed rate and the depth of cut. In
Figure 4 the roundness decreases for all the feed rates
with the cutting speed increasing from 120 m/min to 140
m/min. This can be attributed to the cutting force de-
crease with the increasing cutting speed. This decrease is
explained with a thermal softening of the machined
material in the flow zone due to the increasing cutting
temperature. The highest roundness values for all the
feed rates were obtained at a cutting speed of 160 m/min.
Although a moderate increase in the cutting speed de-
creases the cutting force, a further increase in the cutting
speed was thought to increase the tool wear. This, in
turn, increases the roundness. In addition to the cutting
parameters, the roundness is also affected by other
factors such as the clamping system and the machine-
tool rigidity.
The experimental results and the calculated S/N ratios
for the roundness are given Table 3. The S/N ratios of
the factors for each level are shown in Table 6. The
values in Table 6 are given as the plots from Figure 5. It
is seen from Figure 5 and Table 6 that the control
factors can be used to reach the smallest roundness in the
machining of inner ring bores. The smallest roundness is
obtained in the following cutting conditions: V2, f3 and
MP2. The optimum cutting conditions for the roundness
were V2 (V = 140 m/min), f3 (f = 0.08 mm/r) and MP2
(the 8th machined part). The S/N ratios for the roundness
of Rcal were calculated with Equation (3). Consequently,
G and Rcal calculated for the optimum cutting conditions
were found to be –7.5487 dB and 2.38 μm, respectively.
Table 6: Response table of S/N ratios for roundness
Tabela 6: Tabela odgovorov S/N-razmerje glede na okroglost
Symbol
S/N ratio
Rank
Level 1 Level 2 Level 3 
V –10.079 –8.474 –12.955 4.482 1
f –10.458 –10.838 –10.212 0.626 3
MP –11.044 –9.869 –10.595 1.176 2
Table 7: ANOVA for S/N ratios for roundness
Tabela 7: ANOVA za S/N-razmerje glede na okroglost
Source DF SS MS F P PCR
V 2 30.9353 15.4676 110.12 0.009 91.18
f 2 0.5965 0.2982 2.12 0.320 1.76
MP 2 2.1115 1.0558 7.52 0.117 6.23
Error 2 0.2809 0.1405 0.83
Total 8 33.9242 100
The ANOVA results for the inner-ring-bore round-
ness are given in Table 7. The F-ratios and their PCRs
were taken into consideration to identify the significance
level of the variables. Table 7 shows that the most effec-
tive variable for the roundness is the cutting speed with a
PCR of 91.18 %. The contributions of the number of the
machined part and the feed rate were found to be 6.23 %
and 1.76 %, respectively. Consequently, the feed rate and
the number of the machined part do not have significant
M. BOY et al.: APPLICATION OF THE TAGUCHI METHOD TO OPTIMIZE THE CUTTING CONDITIONS ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 765–772 769
Figure 5: Main-effect plots for roundness
Slika 5: Diagrami glavnih vplivov na okroglost
Figure 4: Interaction plots for roundness
Slika 4: Diagrami medsebojnih vplivov na okroglost
effects on the roundness (P < 0.05). The error ratio was
calculated as 0.83 % and it is the smallest ratio.
3.3 Evaluation of the inner-ring-bore-diameter error
The interaction plots for the inner-diameter error and
the cutting conditions are given in Figure 6. The mini-
mum value of the inner-diameter error was obtained at
V = 140 m/min and f = 0.08 mm/r for the 8th machined
part. The cutting speed might have influenced the result-
ing cutting forces and this interaction might have con-
sequently influenced the diameter error in a number of
ways, i.e., by changing the elastic deformation of the
workpiece, by altering the tool wear, by increasing the
thermal distortion, by forming a built-up edge (BUE) and
by increasing the radial-spindle error. In this case, the
most likely cause for the change in the diameter error is
considered to be the change in the elastic deformation of
the workpiece.
The experimental results and the calculated S/N ratios
for the roundness are given Table 3. Figure 7 shows the
plots of the S/N ratios that were calculated for the
inner-diameter error. The S/N ratios of the factors for
each level are shown in Table 8. Hence, by considering
the S/N ratios from Table 8 and Figure 7, the optimum
cutting conditions for the inner-diameter error were V2
(V = 140 m/min), f3 (f = 0.08 mm/r) and MP2 (the 8th
machined part). The S/N ratios for the roundness of IDe
(the inner diameter error) were calculated with Equation
(3). Consequently, G and IDecal calculated for the opti-
mum cutting conditions were found to be 46.388 dB and
0.0048 μm, respectively.
Table 8: Response table of S/N ratios for inner-diameter error
Tabela 8: Tabela odgovorov razmerja S/N glede na napako notranjega
premera
Symbol
Mean S/N ratio
Rank
Level 1 Level 2 Level 3 
V 37.84 42.65 40.27 4.82 1
f 38.64 39.36 42.76 4.12 2
MP 40.32 41.48 38.96 2.51 3
Table 9: ANOVA for S/N ratios for inner-diameter error
Tabela 9: ANOVA za S/N-razmerje glede na napako notranjega pre-
mera
Source DF SS MS F P PCR
V 2 34.815 17.4075 54.87 0.018 46.99
f 2 29.1223 14.5612 45.9 0.021 39.32
MP 2 9.5031 4.7515 14.98 0.063 12.83
Error 2 0.6344 0.3172 0.86
Total 8 74.0748 100
The ANOVA results for the inner-diameter error are
shown in Table 9. The F-ratios from Table 9 and their
PCR were taken into consideration to identify the signi-
ficance levels of the variables. Table 9 shows that the
most effective variable for the inner-diameter error is the
cutting speed with a PCR of 46.99 %. The feed rate and
the number of the machined part also affect the inner-
diameter error with their PCR being 39.32 and 12.83 %,
respectively. Therefore, the number of the machined part
does not have a significant effect on the inner-diameter
error (P < 0.05). As the P-value of the number of the
machined part is less than 0.05, it is not significant.
3.4 Confirmation tests
Confirmation tests of the control factors were made
for the Taguchi method. The optimization was verified
with the confirmation tests after the determination of the
control factors that give the optimum results. The confir-
mation tests were conducted at the optimum factor levels
for each surface roughness, roundness and diameter
error. The confirmation tests performed at the optimum
variable levels determined for the surface roughness, the
roundness and the inner-diameter error were evaluated
by taking into consideration the confidence interval (CI)
calculated with Equations (4) and (5)16,17:
CI F V
n r
= +
⎛
⎝
⎜
⎞
⎠
⎟0 05 1
1 1
. ( , ) e e
eff
(4)
n
N
eff
T
=
+1 
(5)
M. BOY et al.: APPLICATION OF THE TAGUCHI METHOD TO OPTIMIZE THE CUTTING CONDITIONS ...
770 Materiali in tehnologije / Materials and technology 49 (2015) 5, 765–772
Figure 7: Main-effect plots for inner-diameter error
Slika 7: Diagrami glavnih u~inkov na napako notranjega premera
Figure 6: Interaction plots for inner-diameter error
Slika 6: Interakcijski diagrami za napako notranjega premera
where e is the error degree of freedom, Ve is the error
variance, neff is the repeating number of the experi-
ments, N is the total number of the experiments, T is
the variable’s degree of freedom and r is the number of
the confirmation tests.
Table 10 gives a comparison of the results of the
confirmation tests conducted according to the optimum
levels of the variables and the values calculated using
Equations (2) and (3). Besides, confidence-interval va-
lues for Ra, R and IDe were calculated using Equations
(4) and (5).
Table 10: Comparison between confirmatory test results and calcu-
lated values
Tabela 10: Primerjava med rezultati potrditvenih preizkusov in izra-
~unanimi vrednostmi
Confirmatory
experiment results Calculated values Differences
Ra mea/μm mea/dB Ra cal/μm cal/dB Ra mea –Ra cal
mea – cal
0.151 16.42 0.131 17.621 0.02 1.201
Rmea/μm mea/dB Rcal/μm cal/dB Rmea –Rcal
mea – cal
2.35 –7.4214 2.38 –7.5487 0.03 0.1273
IDemea/μm mea/dB IDecal/μm cal/dB IDemea –IDecal
mea – cal
0.005 46.020 0.0048 46.388 0.0002 0.368
Using Equation (5), confidence values of (2.56, 2.28
and 3.42) dB were obtained for the surface roughness
(Ra), the roundness (R) and the inner-diameter error
(IDe), respectively. Table 10 shows differences between
the values obtained with the confirmatory tests and the
values of the S/N ratios calculated with Equations (2) and
(3). It is seen that a difference of 1.201 dB is under the
5 % confidence interval for the surface roughness of 2.56
dB, a difference of 0.1273 dB is under the 5 % confi-
dence interval of for the roundness of 2.28 dB and, simi-
larly, a difference of 0.368 dB is under the 5 % confi-
dence interval for the inner-diameter error of 3.42 dB.
Therefore, the optimum-control-factor settings for all the
cutting conditions were confirmed as confident.
4 CONCLUSIONS
In this study, the cutting conditions for the surface
roughness, the inner-diameter error and the roundness
during the hard turning of an inner ring bore were opti-
mized with the Taguchi method. The results obtained
from this study are presented below:
According to the results of the statistical analysis, it
was found that the feed rate was the most significant
factor for the surface roughness with a PCR of 62.84 %.
The cutting speed was the most significant parameter for
the roundness and the inner-diameter error with PRC of
91.18 and 46.99 %, respectively.
The optimum levels of the control factors for mini-
mizing the surface roughness, the roundness and the
inner-diameter error using S/N ratios were determined.
The optimum conditions for the surface roughness were
observed at V2-f1-MP1 (i.e., cutting speed V = 140
m/min, feed rate f = 0.04 mm/r and the 4th machined
part). The optimum conditions for the roundness and the
inner-diameter error were observed at the same levels,
V2-f3-MP2 (i.e., cutting speed V = 140 m/min, feed rate f
= 0.08 mm/r and the 8th machined part).
Confirmation tests were carried out at the optimum
conditions. According to the confirmation-test results,
the measured values were found to be within the 95 %
confidence interval.
The outcomes for the roundness and the inner-dia-
meter errors are very close to the predicted values. Thus,
there is no need to carry out confirmation tests if the
cutting conditions found with the optimization procedure
are included in the cutting conditions within the Taguchi
experimental design.
Acknowledgement
The authors would like to thank the Ministry of
Science, Industry and Technology, Turkey (00980.STZ.
2011-2) and ORS Bearings, Turkey, for the financial
support of this study.
5 REFERENCES
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8506(07)63455-6
2 G. Byrne, D. Dornfeld, B. Denkena, Advancing cutting technology,
Annals of the CIRP, 52 (2003) 2, 483–507, doi:10.1016/S0007-
c8506(07)60200-5
3 W. König, A. Berktold, K. F. Koch, Turning versus grinding – a com-
parison of surface integrity aspects and attainable accuracies, Annals
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62387-7
4 F. Klocke, E. Brinksmeier, K. Weinert, Capability profile of hard
cutting and grinding processes, Annals of the CIRP, 54 (2005) 2,
22–45, doi:10.1016/S0007-8506(07)60018-3
5 M. N. Islam, Effect of additional factors on dimensional accuracy
and surface finish of turned parts, Machining Science and Techno-
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10910344.2012.747936
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and surface finish achievable in dry turning, Machining Science and
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doi:10.1080/10910340903451456
7 J. A. Malluck, S. N. Melkote, Modeling of deformation of ring
shaped workpieces due to chucking and cutting forces, Journal of
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doi:10.1115/1.1643079
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tification of causes for dimensional and form deviations of bearing
rings, CIRP Annals – Manufacturing Technology, 56 (2007) 1,
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Improving the shape quality of bearing rings in soft turning by using
a Fast Tool Servo, Prod. Eng. Res. Devel., 3 (2009), 469–474,
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11 D. Stöbener, B. Beekhuis, Application of an in situ measuring system
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thin-walled rings, CIRP Annals – Manufacturing Technology, 62
(2013) 1, 511–514, doi:10.1016/j.cirp.2013.03.129
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on roundness deviations of turned rings, Machining Science and
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tool, Int. J. Refract. Met. Hard., 28 (2010) 3, 349–361, doi:10.1016/
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772 Materiali in tehnologije / Materials and technology 49 (2015) 5, 765–772
L. GETSOV et al.: THERMAL FATIGUE OF SINGLE-CRYSTAL SUPERALLOYS: EXPERIMENTS, ...
773–778
THERMAL FATIGUE OF SINGLE-CRYSTAL SUPERALLOYS:
EXPERIMENTS, CRACK-INITIATION AND
CRACK-PROPAGATION CRITERIA
TOPLOTNO UTRUJANJE MONOKRISTALNIH SUPERZLITIN:
PREIZKUSI, MERILA ZA NASTANEK IN NAPREDOVANJE
RAZPOKE
Leonid Getsov1, Artem Semenov2, Sergey Semenov2, Alexander Rybnikov1,
Elena Tikhomirova3
1NPO CKTI, Saint-Petersburg, Russia
2Saint-Petersburg State Polytechnical University, Russia
3Klimov Company, Saint-Petersburg, Russia
guetsov@yahoo.com, semenov.artem@googlemail.com
Prejem rokopisa – received: 2014-10-03; sprejem za objavo – accepted for publication: 2014-11-06
doi:10.17222/mit.2014.251
Thermal-fatigue tests were performed on corset samples made from five single-crystal nickel-based superalloys with various
crystallographic orientations (〈 〉001 ,〈 〉011 ,〈 〉111 ) at different temperatures and cycle durations. The proposed criteria for the
thermal-fatigue crack initiation and crack propagation are discussed and compared with experimental results. The sensitivity of
durability to the crystallographic orientation and to thermal-cycle parameters is analyzed experimentally and by means of the
corresponding finite-element simulations.
Keywords: nickel-based superalloy, single crystal, thermal fatigue, crack initiation, crack propagation
Preizkusi toplotnega utrujanja so bili izvr{eni na vzorcih korzetne oblike. Vzorci so bili izdelani iz petih monokristalnih
nikljevih superzlitin z razli~no kristalografsko orientacijo (〈 〉001 ,〈 〉011 ,〈 〉111 ), preizkusi pa so bili narejeni pri razli~nih
temperaturah trajanja ciklov. Predlagana merila za nastanek razpok in njihovo napredovanje pri toplotnem utrujanju so opisana
in primerjana z eksperimentalnimi rezultati. Odvisnost zdr`ljivosti od kristalografske orientacije in parametrov toplotnih ciklov
je eksperimentalno analizirana in ugotovljena z ustrezno simulacijo s kon~nimi elementi.
Klju~ne besede: superzlitina na osnovi niklja, monokristal, toplotno utrujanje, nastanek razpok, napredovanje razpoke
1 EXPERIMENTAL DATA
The durability of the single-crystal blades of modern
gas turbines is mainly determined by their ability to
resist the thermal-cycling stress. In Russia, the testing of
the thermal fatigue of corset samples (Figure 1) in va-
cuum conditions, with and without lateral thermal-bar-
rier coatings is widespread. During the test, the shear-
band formation, the conditions of crack nucleation and
the rate of crack propagation are investigated.
Experimental studies were performed on five single-
crystal alloys (ZhS32, ZhS36 and others) developed by
the VIAM research centre, with different alloying and,
most importantly, different carbon contents. The chemi-
cal compositions of the alloys are given in Table 1.
The tests of the thermal-fatigue resistance were
carried out for different crystallographic orientations
(CGOs) of the specimens, various minimum (Tmin) and
maximum temperatures (Tmax) of the cycle, with and
without a dwell time at Tmax.1,2 The test results were cha-
racterized with the number of cycles to the macrocrack
initiation and to the failure of the samples and the
in-plane and out-of-plane residual strains (in, out) in the
central parts of the specimens. The changing of Tmax and
CGO has a significant influence on the thermal-cycling
durability (Table 2).
The results of a comparative analysis of the thermal-
fatigue durabilities of the five considered alloys are given
in Figure 2, for a thermal cycle of 150900 °C (Figures
2a and 2c) and for a cycle of 5001050 °C (Figures 2b
and 2d).
Materiali in tehnologije / Materials and technology 49 (2015) 5, 773–778 773
UDK 539.42:669.24:519.61/.64 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)773(2015)
Figure 1: Specimen for thermal-fatigue tests
Slika 1: Vzorec za preizkus toplotnega utrujanja
Table 2: Effects of cycle temperature and CGO (alloy 2)
Tabela 2: Vpliv temperature cikliranja in CGO (zlitina 2)
Orientation Ò/°Ñ Òmin/°Ñ Òmax/°Ñ
Number of
cycles to
failure
[001]
750
150 900 951
200 950 450
250 1000 63
500
450 950 2635
500 1000 1220
550 1050 356
[011] 750
150 900 560
250 1000 95
Table 3: Effects of dwell time at Tmax and of aging on the thermal-
fatigue resistance for alloy 3 at thermal cycle of 5001050 °C
Tabela 3: Vpliv zdr`ljivosti pri Tmax in staranja na odpornost proti to-
plotnemu utrujanju zlitine 3 pri toplotnem cikliranju 5001050 °C
Condition and test
mode
Number of
cycles to
macrocrack
Number of
cycles to
failure
åin/% åout/%
Without a dwell time
at Òmàõ
670 746 16 –1
Dwell time of 2 min
at Òmàõ
27 91 12 17
Aging at 850 °C,
500 h 85 242 13 4
in, out – residual strains in 2 directions transversal to the sample axis
The obtained results indicate a relatively weak in-
fluence of the chemical composition of the single-crystal
alloys on the thermal-fatigue durability (the differences
are less than 2–5 times). On the other hand, the long-
term strengths and creep-resistance values of the single-
crystal alloys with different compositions differ quite
significantly.3 None of the alloys among 1–5 demon-
strates the maximum durability throughout the whole
thermal-cycling-parameter range.
A significant decrease in the durability is observed
after the dwell time at the maximum temperature of the
cycle (Table 3). The preliminary aging at 850 °C also
leads to a considerable decrease in the durability and
residual strains.
2 CRITERION FOR THE TMF-CRACK
INITIATION
The prediction of the thermo-mechanical-fatigue
(TMF) failures of the single-crystal materials is per-
L. GETSOV et al.: THERMAL FATIGUE OF SINGLE-CRYSTAL SUPERALLOYS: EXPERIMENTS, ...
774 Materiali in tehnologije / Materials and technology 49 (2015) 5, 773–778
Figure 2: Comparison of thermal-fatigue resistance values for alloys
1–5 (number of cycles (a, b) and transversal residual strains at failure
(c, d) for cycles of 150900 °C (a, c) and 5001050 °C (b, d)).
Slika 2: Primerjava odpornosti proti toplotnemu utrujanju zlitin 1–5
({tevilo ciklov (a, b) in pre~ne preostale napetosti pri poru{itvi (c, d)
pri ciklih od 150900 °C (a, c) in 5001050 °C (b, d))
Table 1: Chemical compositions of Ni-based single-crystal alloys (w/%)
Tabela 1: Kemijska sestava monokristalnih zlitin na osnovi Ni (w/%)
C Cr Co Mo W Ta Re Al Ti Nb Re La Ru Ze Si
Alloy 1 0.12–0.18 4.3–5.6
8.0–
10.0 0.8–1.4 7.7–9.5 3.54–5 3.5–4.5 5.6–6.3 – 1.4–1.8 – – – – –
Alloy 2 0.05 4.5 9.0 1.5 12.0 2.0 5.9 1.0 1.0 2.0 – – – –
Alloy 3 <0.02 2.85 6.20 4.0 3.8 4.5 6.2 6.0 – – – – 3.9 <0.2
Alloy 4 <0.02 4.2 8.3 1.8 5.1 6.1 4.05 6.0 0.8 – – – – – <0.2
Alloy 5 <0.02 3.2–3.8 4.6–5.4 4.2–4.8 2.2–2.8 5.7–6.3 0.25–0.55 8.1–8.6 – – – 0.01–0.2 – <0.02 <0.2
formed on the basis of the deformation criterion.4–6 The
macrocrack-initiation criterion is the condition for
achieving the critical value of the total damage initiated
by different mechanisms:
D D D D1 2 3 4 1( ) ( ) ( ) ( )Δ Δ   eq
p
eq
c
eq
p
eq
c+ + + = (1)
Criterion (1) is based on the linear damage summa-
tion of the cyclic plastic strain:
D
T
i
ii
n
1
1
=
=
∑
( )
( )
Δ eq
p k
1C
(2)
cyclic creep strain:
D
T
i
ii
n
2
1
=
=
∑
( )
( )
Δ eq
c m
2C
(3)
one-sided accumulated plastic strain:
D
Tt t3 0
=
≤ ≤
max
( )max


eq
p
r
p (4)
and one-sided accumulated creep strain:
D
Tt t4 0
=
≤ ≤
max
( )max


eq
c
r
c (5)
where C1, C2, k, m,  r
p ,  r
c are the material parameters,
depending on the temperature and crystallographic
orientation. Usually the material parameters k = 2, m =
5/4, C1 = (rp)k, C2 = (3/4rc)m are used.4
Different norms of the strain tensor are considered as
the equivalent strains for (1)–(5), including the maxi-
mum shear strain in the slip system with normal n111 to
the slip plane and slip direction l011:
eq = n111 · e · l011 (6)
the maximum principal strain (the maximum eigenvalue
of the strain tensor):
eq = 1 = max argdet(e–	l) = 0 (7)
the von Mises strain intensity:
eq = [ ]
2
9
1
3
2 2 2 2 2 2( ) ( ) ( ) ( )      
 
 
x y y z z x xy yz zx− + − + − + + + (8)
the Hill strain intensity:
eq = [ ]
2
9
12 2 2 2 2 2( ) ( ) ( )
*
( )      
 
 
x y y z z x xy yz zxK
− + − + − + + +
(9)
and the maximum shear strain:
eq = ( )  1 3− (10)
Equivalent strains (6) and (9) correspond to the
crystallographic failure mode, while equivalent strains
(7), (8) and (10) correspond to the non-crystallographic
failure mode. In (9):
K
r
r
* =
⎛
⎝
⎜
⎞
⎠
⎟ −
⎡
⎣
⎢
⎤
⎦
⎥
〈 〉
〈 〉
−
9 4 1
001
011
2 1


3 RESULTS OF THE FINITE-ELEMENT
SIMULATION
In order to verify criterion (1) a stress-strain analysis
of the single-crystal corset samples (Figure 1) was
carried out using the finite-element (FE) program
PANTOCRATOR7 with an application of physical
models of elastoplasticity. These material models take
into account the fact that inelastic deformation occurs in
accordance with the crystal slip system due to a slip
mechanism and, therefore, the deformation is strongly
sensitive to the crystallographic orientation. The elasto-
plastic and visco-elasto-plastic material models8,9 with
nonlinear kinematic and isotropic hardening, also
including the self-hardening of each system and the
latent hardening10 between the slip systems, are used in
the FE simulations. The application of the visco-elastic
models leads to unrealistically overestimated levels of
the stress.
The obtained results for the inhomogeneous stress,
strain and damage fields allowed us to find the location
of the crack initiation. The damage field is determined
with criterion (1) on the basis of an analysis of the
strain-field evolution using the experimental data on
creep- and elastoplastic-deformation curves. The typical
damage field distribution after the 10th cycle
(20 °CTmax = 900 °CTmin = 150 °C) is presented in
Figure 3 for the sample from alloy 3 with the orientation
near to 〈 〉001 . The best prediction (in comparison with
L. GETSOV et al.: THERMAL FATIGUE OF SINGLE-CRYSTAL SUPERALLOYS: EXPERIMENTS, ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 773–778 775
Figure 4: Influence of the crystal orientation on the cyclic stress-
strain curve. Results of finite-element simulations of thermal cycles
with Tmin = 150 °C and Tmax = 900 °C (the central point of a speci-
men).
Slika 4: Vpliv orientacije kristala na cikli~no krivuljo napetost –
raztezek. Rezultati simulacije s kon~nimi elementi za toplotne cikle s
Tmin = 150 °C in Tmax = 900 °C (sredina vzorca).
Figure 3: Damage-field distribution after the 10th cycle for sample 43
from alloy 3 with near to 001 orientation
Slika 3: Razporeditev polja po{kodbe po 10. ciklu za vzorec 43 iz zli-
tine 3 z orientacijo blizu 001
the experiment) of the number of cycles for the crack
initiation is given for strain measure (6) in this loading
case.
The results of the FE simulations show that the cry-
stallographic orientation has a significant influence on
the stress-strain state of the samples (Figures 4 and 5)
which was also confirmed with the experiments.10
The results of the verification of criterion (1) for
alloys 2 and 3 are given in Figure 6. The computed num-
ber of the cycles to crack initiation demonstrates a
satisfactory accuracy.
The thermal-fatigue failure process for the corset spe-
cimens shows the presence of two distinct modes (cry-
stallographic and non-crystallographic) of the failure.
Their occurrence depends on the CGO and on tempera-
tures Tmin and Tmax of the cycle. The experimental results
and the Arrhenius approximation of the boundary bet-
ween the failure modes are given in Figure 7. The
difference in the failure mode is taken into consideration
for the damage computation based on equation (1) by
choosing the equivalent strain measure (6)–(10) and
analysing the crack propagation by choosing the direc-
tion of the crack growth.
The results of the experiments and computations
showed that the dwell time has a significant effect on the
thermal-fatigue durability. Figure 8 shows the computed
effect of the influence of the dwell time at Tmax on the
thermal-fatigue durability for the specimen from alloy 3
and for the thermal cycle of Tmin = 500 °C  Tmax = 1050 °C.
A comparison of the numbers of the cycles to crack
initiation computed with equation (1) and obtained with
the experimental results is given in Table 4.
L. GETSOV et al.: THERMAL FATIGUE OF SINGLE-CRYSTAL SUPERALLOYS: EXPERIMENTS, ...
776 Materiali in tehnologije / Materials and technology 49 (2015) 5, 773–778
Figure 8: Influence of the dwell time at Tmax on the thermal-fatigue
durability (alloy 3)
Slika 8: Vpliv ~asa zadr`anja pri Tmax na zdr`ljivost pri toplotnem
utrujanju (zlitina 3)
Figure 6: Comparison of experimental data with the results of com-
putations of thermal-fatigue durability for alloys 2 and 3
Slika 6: Primerjava eksperimentalnih podatkov z rezultati izra~unov
odpornosti proti termi~nemu utrujanju za zlitini 2 in 3
Figure 7: Diagram of thermal-fatigue failure modes (alloy 1)
Slika 7: Diagram vrste po{kodb pri toplotnem utrujanju (zlitina 1)
Figure 5: Comparison of deformed states (for clarity increased 10
times) of the central sections of the specimens with orientations 〈 〉001 ,
〈 〉011 and 〈 〉111 for the cooling phase (T = 150 °C) after the 10th cycle.
Colors show the damage-field distribution.
Slika 5: Primerjava deformacijskih stanj (za bolj{o prepoznavnost
pove~ano 10-krat) v srednjem prerezu vzorcev z orientacijo 〈 〉001 ,
〈 〉011 in 〈 〉111 pri fazi ohlajanja (T = 150 °C) po 10. ciklu. Barve pri-
kazujejo podro~je razporeditve po{kodbe.
Table 4: Comparison of the numbers of the cycles to crack initiation
obtained with the computation and with the experiment for various
dwell times at Tmax (alloy 3)
Tabela 4: Primerjava izra~unanega {tevila ciklov do nastanka razpoke
in eksperimentalnih rezultatov pri razli~nih ~asih zadr`anja na Tmax
(zlitina 3)
Dwell time
0 s 120 s
Computation with (1) 480 45
Experiment 670 27
Alloys 1–5 have significantly different creep charac-
teristics and creep-rupture strengths, therefore it can be
expected that the thermal-fatigue durability in the case of
a dwell time at Tmax varies significantly for different
alloys, contrary to the cases of no dwell time at Tmax.
4 CRITERION FOR THE TMF-CRACK
PROPAGATION
The direction of a crack propagation in a single cry-
stal is defined by the crystallographic structure (typically
it is plane 〈 〉111 for the considered alloys11,12 in the
crystallographic mode of failure) and, correspondingly, it
is defined by the stress state during the non-crystallogra-
phic failure mode (a high temperature).
The criterion for the thermal-fatigue-crack propaga-
tion,13 based on a linear sum of the contributions of the
fatigue and creep, is used in order to determine the rate
of crack propagation and the number of cycles to failure.
The criterion is based on using stress-intensity factor
Keff (for the description of the fatigue) and the C*-inte-
gral (for the description of the creep during the dwell
time at the maximum temperature of the cycle). The
mathematical formulation of the criterion represents a
generalization of the Paris power-type equation by taking
into account two mechanisms of cracking:13
d
d
deff
m qa
N
B K A C
ty
= + ∫( ) ( *( ))Δ  
0
(11)
where A, B, m, q are the material constants, defined
separately from the fatigue experiments as da/dN(Keff)
and creep experiments as da/dt(C*) crack-growth data.
The crack-propagation process in the corset sample
(Figure 1) from alloy 2 is simulated cycle by cycle using
the FE program ABAQUS.14 The specimen has its CGO
near to 〈 〉001 . The deviation from the axial orientation is
5.47 ° and from the azimuthal orientation it is 41.97 °
(the Euler angles for rotations ZX’Z’ are:  = 354.7 °, 
= 41.7 °,  = 89.8 °). The CGO was determined with
Laue diffraction patterns.
The location of the crack initiation is defined with an
FE computation in section 3 and it corresponds with the
experimentally obtained result. The crack propagates in
plane 〈 〉111 (Figure 9), except in the rupture stage. The
crack-propagation process is simulated only in plane
〈 〉111 , without out-of-plane deviations, using a step-by-
step technique with FE remeshing for the environment of
the crack front at every step. The initial crack front has
an elliptic form. The Jeff integral is used to calculate
Keff. The fracture zone is established by comparing the
calculated J integral with the JIc value. The details of the
simulation are described in.15
In the experiment the sample failed after 560 thermal
cycles (Tmin = 150 °C Tmax = 900 °C); the registered
crack-initiation stage lasted for 435 cycles and the
crack-propagation stage took 125 cycles. A comparison
of the computed number of cycles for the crack-propa-
gation stage based on the FE simulation (Figure 10) with
the experimental results is given in Table 5.
A simulated crack-front evolution is shown in Figure
11a. It demonstrates a good correlation with the data of a
fractography analysis (Figure 11b).
Table 5: Comparison of the estimated number of cycles for the
crack-propagation stage with the experimental results (alloy 2)
Tabela 5: Primerjava dolo~enega {tevila ciklov za fazo napredovanja
razpok z eksperimentalnimi rezultati (zlitina 2)
Number of cycles
FE computation with (11) 192
Experiment 125
L. GETSOV et al.: THERMAL FATIGUE OF SINGLE-CRYSTAL SUPERALLOYS: EXPERIMENTS, ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 773–778 777
Figure 10: Plastic-strain-intensity-field distribution in the specimen
from alloy 2 with a crack of a length of 0.5 mm
Slika 10: Razporeditev polja intenzitete plasti~ne deformacije v vzor-
cu iz zlitine 2 z dol`ino razpoke 0,5 mm
Figure 9: Views of the specimen from alloy 2 after the failure: a) left
half, b) right half
Slika 9: Videz vzorca zlitine 2 po poru{itvi: a) leva polovica, b) desna
polovica
5 CONCLUSIONS
1. Thermal-fatigue tests were performed on corset
samples from five Russian single-crystal nickel-based
superalloys with various crystallographic orientations
(〈 〉001 , 〈 〉011 , 〈 〉111 ) under different temperatures and
cycle durations with the aims to determine the cha-
racteristics of the thermal-fatigue resistance, compare
the alloys and determine the thermal-fatigue-failure
criterion.
2. The crack-initiation criterion for the single-crystal
alloys under thermal cyclic loadings is proposed and
discussed. A satisfactory correlation is observed bet-
ween the proposed deformation criterion (1) and the
obtained experimental results.
3. The finite-element simulations of the single-crystal
corset specimens under the thermal cyclic loading
were performed using different material models. The
obtained results indicate an applicability of the pro-
posed deformation criteria for failure for the five
considered single-crystal alloys at temperatures of up
to 1100 °C.
4. A high sensitivity of the thermal-fatigue durability to
the crystallographic orientation and to the thermal-
cycle parameters is observed experimentally and also
in the corresponding finite-element simulations of the
corset specimens.
5. The thermal-fatigue failures of the corset specimens
show the presence of crystallographic and non-crys-
tallographic modes of failure. The occurrence of
these modes depends on the crystallographic orien-
tation and the temperature parameters of the cycle.
6. The proposed crack-propagation criterion (11) under
the thermal cyclic loadings demonstrates a satisfac-
tory accuracy in comparison with the experimentally
obtained number of cycles to failure and the crack-
front evolution indicated by the data of a fracto-
graphy analysis.
7. To obtain reliable analysis results for the thermal-
fatigue strength and durability of the single-crystal
blades of GTE, further detailed studies of the mate-
rial characteristics, including the crack-resistance
parameters, carried out over a wide temperature
range, are required.
Acknowledgements
The present study is supported by Russian Science
Foundation, grant No.15-19-00091.
6 REFERENCES
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13 A. S. Semenov, S. G. Semenov, A. A. Nasarenko, L. B. Getsov,
Mater. Tehnol., 46 (2012) 3, 197–203
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L. GETSOV et al.: THERMAL FATIGUE OF SINGLE-CRYSTAL SUPERALLOYS: EXPERIMENTS, ...
778 Materiali in tehnologije / Materials and technology 49 (2015) 5, 773–778
Figure 11: Evolution of the thermal-fatigue crack-front propagation
on a fracture surface: a) results of finite-element simulations, b) data
of a fractography analysis (specimen from alloy 2)
Slika 11: Razvoj napredovanja fronte toplotne utrujenostne razpoke
na povr{ini preloma: a) rezultati simulacije s kon~nimi elementi, b)
podatki iz analize preloma (vzorec iz zlitine 2)
M. B. BILGIN: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE BUILT-UP-LAYER ...
779–784
INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON
THE BUILT-UP-LAYER AND BUILT-UP-EDGE FORMATION
DURING THE MACHINING OF AISI 310 AUSTENITIC STAINLESS
STEELS
PREISKAVA VPLIVOV PARAMETROV REZANJA NA NASTANEK
NAKOPI^ENE PLASTI IN NAKOPI^ENEGA ROBA MED
STRU@ENJEM AVSTENITNEGA NERJAVNEGA JEKLA AISI 310
Mehmet Burak Bilgin
Amasya University, Technology Faculty, Dept. of Mechanical Engineering, Amasya, Turkey
mehmetburak.bilgin@amasya.edu.tr
Prejem rokopisa – received: 2014-10-06; sprejem za objavo – accepted for publication: 2014-11-27
doi:10.17222/mit.2014.253
This study experimentally investigated the effects of machining parameters on the built-up-layer (BUL) and built-up-edge
(BUE) formation and the wear behavior of cutting tools during the machining of the AISI 310 austenitic stainless steel with
titanium-carbide cutting tools. Five different cutting speeds (50, 75, 100, 125 and 150) m/min, three feed rates (0.15, 0.20 and
0.25) mm/r and two cutting depths (1.5, 2) mm were used as the cutting parameters. The highest accumulation of the BUL and
the BUE was observed at a cutting speed of 50 m/min, a feed rate of 0.15 mm/r and a cutting depth of 1.5 mm.
Keywords: built-up edge, built-up layer, machining, stainless steels
Ta {tudija eksperimentalno preu~uje vpliv parametrov rezanja na nakopi~eno plast (BUL) in nakopi~en rob (BUE) ter na vedenje
orodja za odrezovanje med stru`enjem AISI 310 avstenitnega nerjavnega jekla z orodjem iz titanovega karbida. Kot parametri
odrezovanja je bilo uporabljeno pet razli~nih hitrosti (50, 75, 100, 125 in 150) m/min, tri hitrosti podajanja (0,15, 0,20 in 0,25)
mm/r in dve globini rezanja 1,5 mm in 2 mm. Najve~ja akumulacija BUL in BUE je bila opa`ena pri hitrosti rezanja 50 m/min,
hitrosti podajanja 0,15 mm/r in globini rezanja 1,5 mm.
Klju~ne besede: nakopi~en rob, nakopi~ena plast, strojna obdelava, nerjavna jekla
1 INTRODUCTION
In the machining process, workpiece materials may
adhere to the cutting tool in two ways during the cutting.
First, the workpiece becomes welded to the cutting edge,
leading to the formation of what is known as a built-up
edge (BUE). Second, workpiece materials may distribute
and accumulate over a large portion of the cutting tool’s
rake face, leading to the formation of what is known as a
built-up layer (BUL). These situations may arise separa-
tely on a cutting edge, or may occur simultaneously.
AISI 310 stainless steels are materials that are com-
monly used in the manufacturing industry despite their
characteristically poor machinability. The most signifi-
cant machinability problems of stainless steels are the
tool smearing and the BUE formation.
Carrilero and Marcos1 observed that the tool wear
occurs due to a combination of load factors affecting the
cutting edge (mechanical, thermal, chemical, abrasive
factors). Individually, each load factor can be effective at
a certain stage of the tool wear. Generally, the loads do
not act separately and as the machining process conti-
nues the combined effect of the loads increases gra-
dually.
Carrilero et al.2 suggested that, in machining, the tool
wear generally determines the tool life, along with the
criteria such as the cutting strength or the surface rough-
ness that vary in relation to the cutting parameters.
Korkut et al.3 observed that the cutting-tool wear
occurs at a faster rate during the machining of stainless
steel. By increasing the wear of the cutting edges, a BUE
leads to a poor surface finish on a workpiece. While
warm chips are led away from the workpiece, they form
a continuous wire that wears the cutter and adversely
affects the surface of the workpiece. To prevent this, the
operator must clean the chips from each machined work-
piece, which, in turn, hampers the productivity.
Sanchez et al.4 found that certain temperatures,
depending on the cutting parameters, might be generated
during the machining processes, increasing the occur-
rence of wear mechanisms (adhesion, oxidation, fatigue,
abrasion, diffusion). They also suggested that facilitating
the ability of tools to cut within a reasonable period of
time is one of the most important topics that should be
investigated in the studies conducted on the cutting tool
wear.
List et al.5 searched the wear behavior in the machin-
ing of an aluminum-copper alloy (2024). They suggested
that under low cutting conditions, built-up edges form on
the tool’s rake face and take on the function of the
cutting edge. They observed that the continuous sliding
Materiali in tehnologije / Materials and technology 49 (2015) 5, 779–784 779
UDK 621.9:621.941:669.14.018.8 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)779(2015)
of BUE fragments between the tool and chips increases
the tool wear. They also suggest the use of a large rake
angle and a polished tool surface at a low cutting speed
as the adhesion mechanism is more mechanical than
physical.
Rubio et al.6 studied the surface roughness (Ra) of an
AA7050 alloy. They observed that the Ra value of the
AA7050 aluminum pieces obtained through the turning
process had a certain tendency to decrease with the
machined length. They also observed that Ra values
slightly increase with the cutting speed. They suggested
taking into account that parts of the machined material
adhere to the tool on the edge (BUE) and on the rake
face (BUL) during the cutting process.
In addition to the mechanical properties of a material,
Özçatalbaº and Aydin7 discovered that the machining
parameters such as the cutting speed, the feed rate, the
cutting depth and the cutting-edge geometry are also
important to ensure a proper cut.
Gökkaya and Nalbant8 investigated the effects of
different cutting speeds, built-up layers and built-up
edges on aluminum machining. They suggested that
increasing the cutting speed decreased BUL and BUE
formations, but did not eliminate them entirely.
Liew and Ding9 investigated the wear progression in
the low-speed end milling of stainless steel. They ob-
served that a strong bonding between an adherent work
material and a tool surface can result in the formation of
a BUE. When the BUE and the adherent work material
on the tool become unstable, they adhere to the under-
side of the work piece, resulting in a deterioration of the
surface finish.
Katuku et al.10 investigated the wear, cutting forces
and chip characteristics of austempered ductile iron
under finishing conditions. Using cutting speeds ranging
from 50 m/min to 800 m/min, they observed that, at
cutting speeds lower than 150 m/min, the abrasion wear
was the main wear mechanism. They suggested that, at
these cutting speeds, the fragmentation of chips and the
instability of the BUE controlled the dynamic cutting
forces.
Thakur et al.11 studied the machinability of Inconel
718 during high-speed turning. They used cutting speeds
within the range of 40–60 m/min; however, the BUE for-
mation was not observed at the aforementioned machin-
ing parameters.
Chattopadhyay et al.12 studied the wettability and
machinability of pure aluminum for uncoated and coated
carbide cutting-tool inserts. They observed that a large
built-up edge was formed on the uncoated and all the
coated carbide inserts, excluding diamond. They con-
cluded that the flank-wear measurement confirms that
the diamond-coated tool is superior.
Neugebauer et al.13 investigated the velocity effects in
metal forming and machining processes. They suggested
that the adherent material may, assuming a low cutting
velocity, form built-up edges, while at high speeds, a thin
flow layer with extremely high shear deformations tends
to develop.
Zhou et al.14 investigated the surface damage in the
high-speed turning of Inconel 718. They found that, at a
low feed rate, the tendency to form built-up edges is also
higher than at a higher feed rate, due to an increase in the
size of the plastic-deformation area at the interface of the
tool and the workpiece.
Khan et al.15 investigated the tool wear/life in the
finish turning of Inconel 718. They found that, at the
lowest cutting speed (150 m/min), a severe grooving and
a built-up-edge (BUE) formation were observed on the
wear-scar micrographs in all the experiments. They ob-
served that as the cutting speed increased to 300 m/min,
the presence of the grooving and BUE diminished.
Gomez-Parra et al.16 investigated the built-up-edge
and built-up-layer formations in the turning of aluminum
alloys. They suggested that the changes in the BUL and
BUE took place and the formation mechanisms were
related to the changes observed in the roughness profile
of the machined pieces and evaluated through the ave-
rage surface roughness, Ra. They confirmed that the BUE
growth is responsible for a decrease in Ra; this is due to
the fact that a higher BUE thickness can be related to a
lower value of the tool-position angle and, thus, to a
lower value of the maximum height of the cutting finger
on the workpiece surface.
In this study, the machining of the AISI 310 auste-
nitic stainless-steel material, i.e., the CNC turning ma-
chining under dry conditions was performed using a
titanium-carbide cutting tool at five different cutting
speeds, three different feed rates and two different cutt-
ing depths (1.5 mm, 2 mm) as the cutting parameters.
The aim for this study was to establish the ideal cutting-
condition parameters by determining the wear tendencies
of the cutting edge and the effects of the cutting speed,
the feed rate, and the cutting depth on the built-up-layer
and built-up-edge formations with the aid of a scanning
electron microscope (SEM).
2 MATERIALS AND METHODS
A spectral analysis of the AISI 310 stainless-steel
material used in the experiments was performed using a
11814/00 optic emission spectrophotometer. The chemi-
cal composition of the test samples is indicated in Table 1.
All the test samples were extracted by cutting a
single-sized part so that they all consisted of the same
crucible material. Thus, it was ensured that the test
samples had the same physical and chemical charac-
teristics. The sample dimensions used in the tests are
shown in Figure 1.
During the study, the cutting parameters indicated in
Table 2 were applied to the test samples listed in Table
1. Throughout the experiments, a TC 35 Johnford CNC
turning machine with a Fanuc control unit was used
according to the parameters indicated in Table 2. The
780 Materiali in tehnologije / Materials and technology 49 (2015) 5, 779–784
M. B. BILGIN: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE BUILT-UP-LAYER ...
attachment of the test samples to the turning machine
and their machining are schematically represented in
Figure 2.
Due to dry cutting being recommended by a cutter’s
catalogue, the cutting fluid was not used during the
machining of the test samples. A titanium-carbide-co-
vered SNMG 120408-MS US735-shaped cutting edge,
with an ISO M30 certification, recommended for auste-
nitic stainless steel by the Mitsubishi company, a cutt-
ing-tool manufacturer, was used, along with its matching
SSBCR112525 tool holder, for the machining of the test
samples. The choice of the cutting parameters to be used
during the tests was determined by taking into conside-
ration the manufacturing company’s data, prepared
according to ISO 3685. These cutting parameters are
listed in Table 2. Different machining parameters such
as the cutting speed in addition to the feed rate and the
cutting depth were used for each sample.
The optimum cutting speeds were 80–120 m/min for
the chosen cutting edge. To observe the results obtained
with the cutting speeds below and above the recom-
mended cutting speed, cutting speeds of 50–150 m/min
were also included in the tests. The feed rate and the
cutting depth suitable for a radius 0.8 mm cutting-tool
edge were determined based on the ISO 3685 reference
values.
To determine the wear tendency of the cutting edge
resulting from the machining performed with different
cutting parameters, a different cutting edge was used for
each test. Thus, an attempt was made to elucidate the
contribution of each parameter to the built-up-layer and
built-up-edge formations.
Table 2: Cutting parameters used in machining experiments
Tabela 2: Parametri rezanja pri preizkusih strojne obdelave
Expe-
riment
No.
V/
(m/min)
f/
(mm/r) a/mm
Expe-
riment
No.
V/
(m/min)
f/
(mm/r) a/mm
1 50
0.15 1.5
16 50
0.15 2
2 75 17 75
3 100 18 100
4 125 19 125
5 150 20 150
6 50
0.20 1.5
21 50
0.20 2
7 75 22 75
8 100 23 100
9 125 24 125
10 150 25 150
11 50
0.25 1.5
26 50
0.25 2
12 75 27 75
13 100 28 100
14 125 29 125
15 150 30 150
To determine the cutting forces involved in the chip
formation during the machining performed with the
CNC turning machine, a KISTLER 9257B piezoelec-
tric-crystal-based dynamometer was used. To assess the
wear rate, the wear and the smearing behavior of the
cutting edges were examined using the JEOL JSM 6060
LV scanning electron microscope (SEM).
3 DISCUSSION
In this study, machining was performed at five diffe-
rent cutting speeds, three different feed rates and two
different cutting depths using a titanium-carbide-covered
tool. The analysis of the cutting edges with scanning
microscopy revealed that the highest BUE formation
occurred on the samples machined at lower cutting
speeds. Hence, despite the use of different cutting para-
meters, the SEM images obtained for a cutting speed of
50 m/min and with the highest amounts of the built-up
edge and built-up layer were evaluated. The scanning-
electron-microscopy images of the BUE formed on the
M. B. BILGIN: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE BUILT-UP-LAYER ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 779–784 781
Table 1: Chemical composition of AISI 310 austenitic stainless steel (w/%)
Tabela 1: Kemijska sestava AISI 310 avstenitnega nerjavnega jekla (w/%)
C Si Mn P S Cr Mo Ni
< 0.0050 0.3885 2.173 > 0.0960 0.0344 17.67 0.2983 14.71
Al Cu Nb Ti V Fe
0.0145 0.0806 < 0.0050 < 0.0010 0.0457 61.70
Figure 2: Attachment of the test samples and measurement of cutting
forces
Slika 2: Namestitev preizkusnih vzorcev in merjenje sil pri rezanju
Figure 1: Dimensions of the samples used in the experiments
Slika 1: Dimenzije vzorcev, uporabljenih pri preizkusih
cutting edge with the test samples machined at the cutt-
ing speed 50 m/min are provided in Figure 3. A uni-
formly distributed BUE can be seen in Figure 3, while a
non-uniformly distributed BUE can be seen in other
figures.
It is possible to claim that within the experimental
ranges, the feed rate and cutting depth were not as signi-
M. B. BILGIN: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE BUILT-UP-LAYER ...
782 Materiali in tehnologije / Materials and technology 49 (2015) 5, 779–784
Figure 4: a), b) SEM images and c) a cutting-force diagram of the
cutting tool under the machining conditions of 50 m/min cutting
speed, 0.15 mm/r feed rate and 1.5 mm cut depth
Slika 4: a), b) SEM-posnetka in c) diagram sile rezanja rezalnega
orodja pri strojni obdelavi: hitrost rezanja 50 m/min, hitrost podajanja
0,15 mm/r in globina rezanja 1,5 mm
Figure 3: SEM images of BUE and BUL formed on the radius of the
cutting tool used for the experimental sample at a cutting speed 50
m/min
Slika 3: SEM-posnetki BUE in BUL, ki nastajajo na radiusu rezilnega
orodja, uporabljenega pri vzorcu, s hitrostjo rezanja 50 m/min
Table 3: Wear of the cutting edges used in the experiments
Tabela 3: Obraba rezilnih robov pri preizkusih
Expe-
riment
No
V/
(m/min)
f/
(mm/r)
a/
mm Evaluation
Expe-
riment
No
V/
(m/min)
f/
(mm/r)
a/
mm Evaluation
1 50
0.15 1.5
There is a smooth BUE at the
edge 16 50
0.15 2
There is a notch at the cutting
edge because of the BUE
2 75 There is very little wear at theedge 17 75
The edge is clean
3 100 The edge is clean; there is verylittle accumulation at the far end 18 100
The edge is clean, there is a
crush inside
4 125 The edge is clean 19 125 The edge is clean, there is alittle accumulation inside
5 150 There is a little notch at theedge 20 150
The edge is clean
6 50
0.20 1.5
There is a BUE 21 50
0.20 2
There is crater erosion at the
edge; there is a crush inside
7 75 There is a notch at the far end 22 75 There is notching at the sidesurface; there is a crush inside
8 100 There is very little BUE at theside surface 23 100
The edge is clean; there are
residues
9 125 The edge is clean 24 125 There are notches at the edgeand side surface
10 150 The edge is clean 25 150 The edge is clean
11 50
0.25 1.5
Plenty of BUE; there is a notch
at the far end 26 50
0.25 2
There is an accumulation at the
side surface
12 75 There is a little BUE at the sidesurface 27 75
The edge is clean; there is a
crush inside
13 100 There is some BUE at the sidesurface 28 100
The edge is clean; there is a
crush inside
14 125 There is a BUE line at the farend 29 125
The edge is clean; there is a
crush inside
15 150 The edge is clean 30 150 The edge is clean; there is acrush inside
ficant as the cutting speed for the formation of the
built-up edge and built-up layer. However, when the
images in Figure 3 are examined, it can be seen that the
feed rate and the cutting depth also had a slight effect on
the BUE and BUL, in addition to the cutting speed. It
can be stated that increasing the cutting depth may
slightly increase the BUE and BUL, while increasing the
feed rate may lead to a slight decrease in the BUE and
BUL.
The SEM images of the formations observed on the
cutting tool under the machining conditions of the
cutting speed 50 m/min, feed rate 0.15 mm/r, and cutting
depth 1.5 mm are provided in Figure 4. As can be seen
in Figure 4, if a built-up layer forms on the rake face of
the cutting tool, the built-up edge forms uniformly along
the main cutting edge and the tool-nose radius progresses
towards the cutting tool’s rake face. It was noted that the
cutting forces increased when the BUE also increased.
This is demonstrated in a two-dimensional image shown
as Figure 4a and a three-dimensional image of Figure
4b. The cutting-force diagram is given in Figure 4c.
The images obtained with scanning electron micro-
scopy for the other cutting edges used in the machining
of the test samples are not provided. Instead, an evalu-
ation of all the cutting parameters of the cutting edges is
provided in Table 3. When the results obtained with the
SEM for the cutting edges used during the test are
examined, a decrease in the cutting-tool wear corres-
ponding to a decrease in the smearing tendency is ob-
served, as seen in Figures 5 and 6. In relation to the
decrease in the smearing tendency, a decrease in the
cutting-tool wear can also be observed when Figures 5
and 6 are examined. In the tests conducted with the
cutting speed of 50 m/min, the feed rate of 0.15 mm/r
and the cutting depth of 1.5 mm, the wear of the cutting
tool was observed to decrease as the cutting speed
increased, despite the presence of a notch in the cutting
edge.
In Figure 5, it is possible to see the traces of the
built-up-layer and built-up-edge formations on the
cutting tool after the machining at 75–100 m/min.
In Figure 6, the built-up layer can be observed once
again on the rake face of the cutting tool. On the other
hand, a built-up edge formed on the main cutting edge.
At cutting speeds of 125 m/min and 150 m/min, smaller
built-up layers and built-up edges formed. However, the
built-up edge extended to the cutting tool’s rake face due
to an accumulation. A decrease in the BUE, correspond-
ing to a decrease in the cutting strength, can be seen in
Figure 6c.
It was observed that the cutting condition affected the
formation of the built-up layer and built-up edge. The
built-up-layer and built-up-edge thicknesses were ob-
served to be on an increase up to the critical level as long
as the machining was performed. Following this in-
crease, the built-up edge was identified as being plasti-
cally deformed and extending towards the cutting-tool
surface. For this reason, the previous built-up layer was
covered and remained below the new layer. The built-
up-layer and built-up-edge formations disrupt the geo-
metry of the cutting tool. Disrupting the tool geometry
and efficiency causes a workpiece to have a rough sur-
face and reduced efficiency. The lowest surface quality
was achieved with the cutting speed of 50 m/min, the
0.25 mm/d feed rate, and the cutting depth of 2 mm,
which resulted in a surface-roughness value (Ra) of 5.75
μm. The formation of a BUE, particularly on the chip
roots, contributed to the increase in the surface rough-
ness.
In the assessments that were performed, the sizes of
the built-up layer and built-up edge formed at the cutting
speed of 50 m/min were observed as being larger than
the built-up layer and built-up edge formed at the cutting
speed of 150 m/min. At the 150 m/min cutting speed, no
built-up edge was identified in the environs of the tool-
nose radius that extended towards the rake face. This
M. B. BILGIN: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE BUILT-UP-LAYER ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 779–784 783
Figure 6: Cutting-edge radius used for the test samples machined at
cutting speeds of: a) 125 m/min, b) 150 m/min and c) cutting-force
diagram
Slika 6: Uporabljeni radij rezilnega orodja pri obdelavi vzorca s
hitrostjo rezanja: a) 125 m/min, b) 150 m/min in c) diagram sile
rezanja
Figure 5: Cutting-edge radius used for the test samples machined at
cutting speeds of: a) 75 m/min and b) 100 m/min
Slika 5: Radius rezilnega roba pri vzorcu, obdelanem s hitrostjo
rezanja: a) 75 m/min in b) 100 m/min
situation demonstrated that the cutting speed is an
important parameter for the formation of a built-up edge.
The causes of this decrease can be explained with the
effects of the factors such as the ease of deformation pro-
cesses associated with the increasing temperatures at
high cutting speeds, the ease of deformation of the work-
piece material around the cutting edge and tool-nose
radius, and the plastic flow formed at high temperatures.
The temperatures generated by the increase in the
cutting speed result in soft layers that create a plastic
flow. This, in turn, prevents the formation of built-up
layers and built-up edges. Based on the observations
from the experiments performed, it can be stated that
cutting speeds above 100 m/min are necessary to prevent
the formation of built-up layers and built-up edges.
4 CONCLUSIONS
Cutting tools were examined under a SEM to deter-
mine the effects of machining parameters on the forma-
tion of built-up layers and built-up edges on the cutting-
tool surfaces during machining processes. The results
obtained within the experimental ranges are provided
below:
• The effects of the feed rate and cutting depth on the
formation of built-up edges and built-up layers were
not as significant as the effect of the cutting speed.
• built-up edge formed on the main cutting edge of the
cutting tool.
• The built-up-edge formation along the main cutting
edge of the cutting tool did not affect the characte-
ristics of the main cutting edge.
• The wear of the cutting tools was observed mainly at
low cutting speeds, particularly at 50 m/min. It is
possible to claim that an increasing cutting speed
leads to a decrease in the cutting wear.
• A built-up layer formed on the rake face of the cutt-
ing tool.
• A built-up edge formed at a low cutting speed of 50
m/min was observed to progress from the environs of
the tool-nose radius towards the surface of the rake
face.
• In the machining of the AISI 310 austenitic stainless
steel, the built-up-layer and built-up-edge formations
decreased as the tested cutting speeds increased.
However, higher speeds could not entirely prevent the
formations of built-up layers and built-up edges.
• It can be asserted that, during the machining of the
AISI 310 austenitic stainless steel with titanium-
carbide cutting tools, cutting speeds higher than 100
m/min are necessary to prevent the built-up-layer and
built-up-edge formations, as well as the cutting-tool
wear.
5 REFERENCES
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Aluminium Alloys, J of Mech Behav of Mater, 7 (1996) 3, 179–191,
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M. B. BILGIN: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE BUILT-UP-LAYER ...
784 Materiali in tehnologije / Materials and technology 49 (2015) 5, 779–784
M. GUTMAN et al.: MORTAR-TYPE IDENTIFICATION FOR THE PURPOSE OF RECONSTRUCTING ...
785–790
MORTAR-TYPE IDENTIFICATION FOR THE PURPOSE OF
RECONSTRUCTING FRAGMENTED ROMAN WALL PAINTINGS
(CELJE, SLOVENIA)
ANALIZA OMETOV ZA REKONSTRUKCIJO FRAGMENTOV
RIMSKIH STENSKIH POSLIKAV (CELJE, SLOVENIJA)
Maja Gutman1, Martina Lesar Kikelj1, Jelka Kuret1, Sabina Kramar2
1Institute for the Protection of Cultural Heritage of Slovenia, Restoration Centre, Poljanska 40, 1000 Ljubljana, Slovenia
2Slovenian National Building and Civil Engineering Institute, Dimi~eva 12, 1000 Ljubljana, Slovenia
maja.gutman@rescen.si
Prejem rokopisa – received: 2014-10-08; sprejem za objavo – accepted for publication: 2014-11-07
doi:10.17222/mit.2014.256
Archaeological excavations of Roman Celeia (present-day Celje, Slovenia) carried out about 30 years ago revealed the remains
of a Roman residential building where more than 9000 fragments of wall paintings were found. The past manual reassembling
of fragments proved to be time consuming, limiting the amount of the material that can be examined and reconstructed. How-
ever, recently work has been resumed using a software application specifically developed for fragment reassembly. Information
on mortar types can provide additional data about the fragments and thus help in the reconstruction process. Fragments of the
wall paintings consist of up to three preserved mortar layers, differing in their thicknesses, aggregate compositions and binder
colours. Based on the mineralogical/petrographic analyses of the mortars, the investigated fragments of the wall paintings were
divided into several groups. The results revealed that the first-mortar-layer aggregates of all the fragments consist of carbonate
grains, while the other-mortar-layer aggregates consist of carbonate, silicate/carbonate, silicate or silicate/ceramic grains.
Keywords: wall paintings, Roman mortars, petrography, reassembling, fragments
Arheolo{ka izkopavanja rimske Celeie (dana{nje Celje) pred pribli`no 30 leti so razkrila ostanke anti~nega stanovanjskega
objekta, v katerem je bilo najdenih ve~ kot 9000 fragmentov stenskih poslikav. Sestavljanje tak{nega {tevila fragmentov zahteva
veliko ~asa. Nedavno je bila za sestavljanje fragmentov razvita posebna ra~unalni{ka aplikacija. Informacije o sestavi malte se
uporabljajo kot dodatni podatki pri sestavljanju fragmentov. Fragmente stenskih poslikav sestavljajo do tri ohranjene plasti
ometov, ki se razlikujejo po debelini, sestavi agregata in barvi veziva. Analizirani fragmenti so bili na podlagi
mineralo{ko-petrografskih preiskav razdeljeni v ve~ skupin. Rezultati so pokazali, da pri vseh fragmentih zgornjo plast ometa
sestavlja agregat iz karbonatnih zrn, medtem ko agregat spodnjih plasti sestavljajo bodisi karbonatna, silikatna/karbonatna,
silikatna ali silikatna/kerami~na zrna.
Klju~ne besede: stenske poslikave, anti~ni ometi, petrografija, sestavljanje, fragmenti
1 INTRODUCTION
The characterization of ancient wall paintings pro-
vides useful information about the painting technique
and pigments used. Romans used lime of high purity,
while aggregates most often contained siliceous sand,
crushed marble and crushed ceramic.1–3 A similar situ-
ation can be seen on the Roman wall paintings in Slo-
venia where, for instance, crushed carbonate grains are
characteristic of the first mortar layer of the wall
paintings from Roman Emona (today’s Ljubljana), while
in the lower mortar layers, besides silicate sand, grains of
crushed ceramics also appear.4 On the other hand, fluvial
dolomite sand with an addition of crushed ceramic5 or, in
some cases, crushed marble prevails6 in the wall-painting
mortars of the Roman-villa rustica-bath complex near
Mo{nje.
Archaeological excavations in Celje in 1978 unco-
vered the remains of a Roman residential building from
the first century A.D. where more than 9000 fragments
of wall paintings were found. The elongated room of the
residential building, measuring 4 m × 13 m, was heated
with a hypocaust; the floors were decorated with a
black-and-white mosaic with a simple rosette pattern,
while the preserved walls were embellished with wall
paintings.
Due to the large amount of fragments concentrated in
the elongated room of the archaeological site, it is still
unknown whether the fragments actually belonged to the
Roman residential building or perhaps to a pit where the
Romans deposited the material during a renovation
process. The majority of the fragments are monochroma-
tic (most often of white, red, green, blue, black or yellow
colour), but on some of them various motifs such as
animals, flowers and various patterns can be recognized.
The pigments identified were Egyptian blue, vermilion,
red ochre, yellow ochre and green earth,7,8 which were
all common pigments of the Roman period.9,10 The ex-
cavated fragments were first transported to the Regional
Museum in Celje and in 1989 to the Restoration Centre
of the Institute for the Protection of Cultural Heritage of
Slovenia. The manual reassembly process of the frag-
mented wall paintings performed 30 years ago proved to
be time consuming and limiting the amount of the mate-
Materiali in tehnologije / Materials and technology 49 (2015) 5, 785–790 785
UDK 75:693.611:555.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)785(2015)
rial that could be examined and reconstructed. Figure 1
shows three wall paintings reconstructed in 1989,
measuring about 200 cm in length and 30 cm to 90 cm in
width.
A total of 9521 fragments thus still remained dis-
assembled and are now documented and stored in 166
boxes. Given the large scale of the fragments and a high
quality of the wall paintings, a decision for a reconstruc-
tion of the wall paintings and a presentation of the re-
maining fragments was made. In June 2010, a systematic
cleaning and consolidation of the fragments began. The
fragments were consolidated using nanoparticles of
calcium hydroxide (nanolime) which is compatible with
the lime-mortar layers of the wall paintings.
Due to the large number of fragments, it is difficult to
have a comprehensive overview of the entity of the
excavated wall paintings, but fortunately several systems
for a virtual reconstruction of the wall-painting frag-
ments, such as the computer-assisted reassembly of the
wall paintings from the Akrotiri excavations in Thera
(Santorini), Greece11 or the 3D reassembling of the parts
of the wall paintings belonging to the Mycenaean
civilization (c. 1300 BC), enabled us to overcome this
problem.12 Another advantage of a virtual reconstruction
is also that it eliminates the physical contact with the
fragments and thus prevents additional damage to the
original pieces. Additionally, a system for a computer-
assisted documentation and reassembling of the wall-
painting fragments named Pedius13 was developed espe-
cially for the Celje wall paintings.14,15 Its online version,
e-Pedius, is a newly developed computer program
intended not only for restorers but also non-experts.
By means of the Pedius system, all the fragments
were digitized (in 2D) and assigned a unique identifica-
tion number (Figure 2). Capturing images of the
fragments was carried out in two ways: with a scanner
(Figure 3) or with a digital camera.
After the upload of the images in the database, a user
can start an interactive assembling of the fragments (Fig-
ure 4). Since there is no need to manipulate the actual
valuable fragments, the reassembling can also be per-
formed by non-expert users. With the computer support,
the assembling is much accelerated since the program
provides some suggestions based on the colour of the
surface, the lines, etc.
M. GUTMAN et al.: MORTAR-TYPE IDENTIFICATION FOR THE PURPOSE OF RECONSTRUCTING ...
786 Materiali in tehnologije / Materials and technology 49 (2015) 5, 785–790
Figure 3: Scanning a wall-painting fragment
Slika 3: Skeniranje fragmenta stenske poslikave
Figure 2: Fragment with an identification number
Slika 2: Fragment z identifikacijsko {tevilko
Figure 1: Wall paintings with various motifs reconstructed in 1989
Slika 1: Leta 1989 rekonstruirane stenske poslikave z razli~nimi
motivi
Figure 4: Example of interactive assembling of wall-painting frag-
ments in Pedius
Slika 4: Zgled interaktivnega sestavljanja fragmentov stenskih po-
slikav v Pediusu
As the information on the mortar types can provide
additional data about the fragments and help in the
reconstruction process, the study deals with a mineralo-
gical/petrographic characterization of the mortars of the
Roman wall-painting fragments. The results obtained
may not only help in the reconstruction of the frag-
mented wall paintings but also contribute new knowl-
edge about Roman mortars and the technology of wall
paintings in Slovenia and worldwide.
2 EXPERIMENTAL WORK
2.1 Materials
A total of thirty samples were selected among the
fragments (Table 1). The fragments of the wall paintings
consist of up to three preserved mortar layers that vary in
their thicknesses, aggregates and binder colours.
2.2 Methods
In order to characterise the textures of the mortars
and define the contents of particular aggregate compo-
nents, the content of the binder and their ratios, our first
approach was to examine thin sections of the fragments
by means of optical microscopy. Polished thin sections
of the mortar samples were studied with light micro-
scopy using an Olympus BX-60 equipped with a digital
camera (Olympus JVC3-CCD).
Raman spectra of the studied areas of the polished
thin sections of the samples were obtained with a Horiba
Jobin Yvon LabRAM HR800 Raman spectrometer
equipped with an Olympus BXFM light microscope.
Measurements were made using a laser-excitation line
785 nm, and a Leica 50× objective. The spectral resolu-
tion was about 1 cm–1.
M. GUTMAN et al.: MORTAR-TYPE IDENTIFICATION FOR THE PURPOSE OF RECONSTRUCTING ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 785–790 787
Table 1: Groups of analyzed samples
Tabela 1: Skupine analiziranih vzorcev
I. White binder of 1st layer
Group Subgroup Composition Sample
A –
1st layer thickness:
2–5 mm
A – 1
1st layer: Carbonate grains
2nd layer: Carbonate grains
3rd layer: Silicate grains
HTM 9, HTM 12, HTM 21
A – 2
1st layer: Carbonate grains
2nd layer: Carbonate grains
3rd layer: Silicate grains (with ceramic grains)
HTM 16, HTM 47, HTM 48
A – 3
1st layer: Carbonate grains
2nd layer: Silicate grains
3rd layer: Silicate grains
HTM 6, HTM 7, HTM 8, HTM 11,
HTM 14, HTM 45, HTM 50, HTM 53
A – 4
1st layer: Carbonate grains
2nd layer: Carbonate grains : Silicate grains = 2 : 1
3rd layer: Silicate grains (with ceramic grains)
HTM 17, HTM 20
A – 5
1st layer: Carbonate grains
2nd layer: Carbonate grains : Silicate grains = 2 : 1
3rd layer: Silicate grains
HTM 49, HTM 52
A – 6
1st layer: Carbonate grains
2nd layer: Silicate grains (with ceramic grains)
3rd layer: Silicate grains (with ceramic grains)
HTM 46
B –
1st layer thickness:
6–10 mm
B – 1
1st layer: Carbonate grains
2nd layer: Carbonate grains
3rd layer: Silicate grains
HTM 13
B – 2 1
st layer: Carbonate grains
2nd layer: Silicate grains HTM 10, HTM 15
B – 3
1st layer: Carbonate grains
2nd layer: Silicate grains
3rd layer: Silicate grains (with ceramic grains)
HTM 18
B – 4 1
st layer: Carbonate grains
2nd layer: Silicate grains (with ceramic grains) HTM 19
II. Red binder of 1st layer
C –
1st layer thickness:
2–5 mm
C – 1
1st layer: Carbonate grains
2nd layer: Carbonate grains
3rd layer: Silicate grains
HTM 51, HTM 54
C – 2
1st layer: Carbonate grains
2nd layer: Carbonate grains
3rd layer: Carbonate grains
HTM 58
C – 3 1
st layer: Carbonate grains (calcite)
2nd layer: Silicate grains HTM 55
C – 4 1
st layer: Carbonate grains (calcite)
2nd layer: Carbonate grains : Silicate grains = 2 : 1 HTM 56, HTM 57
3 RESULTS AND DISCUSSION
Each fragment has a multilayered structure typical
for the wall paintings that consist of two or three sepa-
rate mortar layers followed by the paint layer (Figure 5).
As seen in Table 1, the first mortar layer (from the top
outer layer) always contains a white carbonate aggregate
but varies in the binder colour and layer thickness (the
latter could vary between 2 mm and 10 mm). The aggre-
gate of the second mortar layer, which is normally bet-
ween 5 mm and 16 mm, most often consists of carbonate
grains, followed by silicate grains or even a mixture of
the two. Furthermore, the third mortar layer (the inner
layer), with a thickness of up to 30 mm, mainly consists
of silicate grains, sometimes with an addition of crushed
ceramics. What binds them is the air lime or, in the cases
of ceramic grains, the hydraulic binder and air lime.
Lime lumps were present in all the samples from all the
mortar layers, and they varied in size from 0.2 mm to
3 mm. Fissures were observed in most of the lower-layer
samples with a silicate-sand aggregate, while the upper
outer layers with carbonate grains were generally more
compact.
The fragments were divided into two main groups
based on the binder colour of the first mortar layer: (I)
the fragments with white binder (most frequent – 24
fragments) and (II) the fragments with red binder (rare –
6 fragments). The samples of the first main group were,
according to the thickness of the first layer, furthermore
divided into two groups; (A) the samples with the
thickness of the first mortar layer between 2 mm and 5
mm (19 fragments) and (B) the samples with the thick-
ness of the first mortar layer between 6 mm and 10 mm
(5 fragments). On the other hand, the second main group
with the coloured binder consists of merely one sub-
group, containing the samples whose first-layer thick-
ness is between 2 mm and 5 mm. Considering the mine-
ralogical/petrographic compositions of the mortar layers,
the samples were grouped into 14 subgroups.
The common characteristic of the first mortar layers
of all the samples is an exclusive presence of carbonate
grains. Carbonate grains, determined by Raman miscro-
spectroscopy, are mostly represented by dolomite
(Figure 6a), while calcite (Figure 6b) is observed only
in a few samples (HTM 55, HTM 56, and HTM 57).
Most of the dolomite is represented by semi-angular
grains (Figure 7a) with a small amount of subrounded
grains. Crushed dolomitic rock was most probably used
as an aggregate (cf. the wall paintings from the Roman
site near Mo{nje with round grains of dolomite indicat-
ing a fluvial deposit).5 In addition, a coarse-grained
sparry calcite (Figure 7b) may indicate the use of
crushed marble.1,6,16
Furthermore, the grains in the aggregate are poorly
sorted, between 0.02 mm and 3.0 mm in size. The binder
is lime, mostly compact with rare fissures. In all three
samples with the coarse-grained sparry calcite, the
M. GUTMAN et al.: MORTAR-TYPE IDENTIFICATION FOR THE PURPOSE OF RECONSTRUCTING ...
788 Materiali in tehnologije / Materials and technology 49 (2015) 5, 785–790
Figure 6: Raman spectra of: a) dolomite, sample HTM 54 and b) cal-
cite, sample HTM 55
Slika 6: Ramanski spekter: a) dolomita, vzorec HTM 54 in b) kalcita,
vzorec HTM 55
Figure 5: Wall-painting fragments representing various groups with
different mortar layers: a) group B-2, sample HTM 10, b) group B-4,
sample HTM 19, c) group C-1, sample HTM 54, d) group A-2, sample
HTM 47, e) group A-3, sample HTM 7, and f) group B-1, sample
HTM 13; Legend: 1: 1st mortar layer, 2: 2nd mortar layer, 3: 3rd mortar layer;
light microscope, reflective light, crossed polars, 5-times magnifications
Slika 5: Fragmenti stenskih poslikav, razdeljeni v razli~ne skupine
glede na sestavo ometov: a) skupina B-2, vzorec HTM 10, b) skupina
B-4, vzorec HTM 19, c) skupina C-1, vzorec HTM 54, d) skupina
A-2, vzorec HTM 47, e) skupina A-3, vzorec HTM 7 in f) skupina
B-1, vzorec HTM 13; legenda: 1: prva plast ometa, 2: druga plast ometa, 3:
tretja plast ometa; svetlobni mikroskop, refleksijska svetloba, navzkri`ni nikoli,
5-kratna pove~ava
binder is red due to the addition of red ochre (hematite)
pigment (Figure 8). A slightly red-coloured binder was
also observed in some Emona samples, which was due to
the addition of red ochre,4 whereas Weber16 reported
additions of red ochre or cinnabar to the lime.
Regarding the second mortar layers, in general, three
types of aggregate were observed: layers with predomi-
nant grains of carbonate, layers with carbonate grains and
an addition of silicate grains or layers with prevailing
silicate grains. Silicate aggregates such as quartz, chert and
lithic grains of sedimentary and magmatic rocks were
observed. Grains are rounded, subrounded and angular,
measuring between 0.03 mm and 3.75 mm. In contrast to
the wall paintings from some other archaeological sites
in Slovenia where ceramic grains were observed
exclusively in the first mortar layers,4,5 ceramic grains
measuring between 0.04 mm and 2.90 mm were present
in the second layer of the two samples (Figure 7c).
In all the samples, the aggregate of the third mortar
layer consists mostly of silicate grains, except for the
HTM 58 sample where carbonate grains are present. The
silicate aggregate is composed of subrounded quartz,
rounded lithic grains of chert, quartz sandstone, magma-
tic rocks, some angular grains of feldspar and mica
(Figure 7d). The grains measure between 0.06 mm and
3.50 mm in size. In some samples, significant amounts
of ceramic grains in a size between 0.5 mm and 3 mm
are present (for comparison: ceramic grains in the Emo-
na samples are rare and observed only in a few
samples.)4
The results show that the wall-painting fragments
with more sophisticated motifs or patterns, such as mul-
tiple colours (groups A-1, A-4, A-5, B-1, C-2, C-4), had
two or even three mortar layers prepared with a carbo-
nate aggregate or at least a mixture of carbonate and
silicate, while those with monochromatic colours usually
had only one carbonate layer followed by silicate sand.
Vitruvius17 wrote in his book about the correct method of
plastering walls and ceilings and of making a quality
base for wall paintings that plaster would not crack and
would be without defects if the walls were covered with
three layers of sand mortar and as many layers of marble
mortar. The use of marble powder in the top layers
improved the polishing effect and produced a mirror-like
sheen on the surface. It is presumed that for the mortar
layers with a mixture of carbonate and silicate grains
(the second layer) crushed dolomite was mixed with
fluvial sand. Dolomite was crushed most probably to
obtain the so-called "polvere di marmo", but normally
limestone was used.18
An addition of ceramic powder reduces the water
permeability and increases the mechanical strength,
which is attributed to the hydraulic reaction with lime in
the presence of water at the edge of the ceramic parti-
cles.19 This highly hydraulic lime has been used since the
Roman times for rendering and plastering the buildings
situated in the places with a high relative humidity, such
as baths.20
The fragments of the same group display very similar
micro-stratigraphic sequences and materials and there-
fore very probably belong to the same construction phase
or room. Thus, such a large number of groups could indi-
cate the existence of several different phases of the
construction and different rooms.
The petrographic composition of the mortar aggre-
gate could reflect local geological conditions. Since
deposits of the upper and middle Triassic dolomite occur
in the Celje area,21 we could assume that the dolomite
aggregate originated from the quarry in the vicinity. Fur-
M. GUTMAN et al.: MORTAR-TYPE IDENTIFICATION FOR THE PURPOSE OF RECONSTRUCTING ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 785–790 789
Figure 8: Raman spectra of red ochre (hematite), sample HTM 56
Slika 8: Ramanski spekter rde~ega okra (hematit), vzorec HTM 56
Figure 7: a) Dolomite aggregate in the first mortar layer, group B-2,
sample HTM 10; light microscope, transmission light, parallel polars,
b) coarse-grained calcite and slightly red-coloured binder in the first
mortar layer, group C-4, sample HTM 56; light microscope, trans-
mission light, parallel polars, c) ceramic particles in the second mortar
layer, group B-4, sample HTM 19; light microscope, transmission
light, crossed polars, d) silicate grains in the third mortar layer, group
A-3, sample HTM 8; light microscope, transmission light, crossed
polars
Slika 7: a) Zrna dolomita v zgornji plasti ometa, skupina B-2, vzorec
HTM 10; svetlobni mikroskop, transmisijska svetloba, vzporedni
nikoli, b) debelozrnat kalcit in rahlo rde~e obarvano vezivo v zgornji
plasti ometa, skupina C-4, vzorec HTM 56; svetlobni mikroskop,
transmisijska svetloba, vzporedni nikoli, c) zrna keramike v drugi pla-
sti ometa, skupina B-4, vzorec HTM 19; svetlobni mikroskop, trans-
misijska svetloba, navzkri`ni nikoli, d) silikatna zrna v tretji plasti
ometa, skupina A-3, vzorec HTM 8; svetlobni mikroskop, transmisij-
ska svetloba, navzkri`ni nikoli
thermore, the round-grain aggregate from the lower mor-
tar layers suggests a fluvial origin from the rivers nearby
the archaeological site, most probably the alluvial depo-
sits of the Savinja River. A possible source of the
coarse-grained calcite could be the abandoned marble
quarry in Pohorje (40 km from the archaeological site)
active in the Roman times.
4 CONCLUSIONS
The characterisation of the wall-painting mortars
from the archaeological site in Celje revealed differences
between the mortars employed and contributed to the
knowledge of the Roman wall-painting technique. The
samples were divided into two main groups based on the
binder colour and three groups based on the thickness of
the first mortar layer. According to the mineral/petrogra-
phic compositions of the layers, the fragments were
further divided into 14 different subgroups. Such a large
number of groups may indicate the existence of several
different phases of the construction and different rooms,
keeping in mind that the fragments did not necessarily
belonged only to the room in which they were found.
The results revealed that the common characteristic
of the first layers of all the samples was the exclusive
presence of carbonate grains. In majority, carbonate
grains were represented by dolomite, while coarse-
grained calcite was found only in three samples that also
indicated the use of crushed marble. The aggregates from
the lower mortar layers consisted of only carbonate
grains or silicate sand, or even of a mixture of both. An
addition of ceramic grains was observed in some cases.
Since crushed ceramic was commonly used in damp
places such as hypocaust, baths and the lower parts of
walls, we can say that some of the analyzed samples
might have belong to such rooms, maybe also being part
of the hypocoust where the fragments had been found.
A reassembly of wall-painting fragments requires
that the fragments be put together in the right way to
form the original artwork. Any knowledge of the
compositions of the mortars (thickness, colour, mineral
composition) not only gives us information about the
technique, but also helps us reconstruct the object. A
combination of an analysis and specific computer
programs, such as Pedius, can help restorers reassemble
the fragments faster and more effectively.
Acknowledgements
This research was financially supported by the
Ministry of Culture of the Republic of Slovenia.
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790 Materiali in tehnologije / Materials and technology 49 (2015) 5, 785–790
P. MAJERI^ et al.: Au-NANOPARTICLE SYNTHESIS VIA ULTRASONIC SPRAY PYROLYSIS ...
791–796
Au-NANOPARTICLE SYNTHESIS VIA ULTRASONIC SPRAY
PYROLYSIS WITH A SEPARATE EVAPORATION ZONE
SINTEZA NANODELCEV ZLATA Z ULTRAZVO^NO RAZPR[ILNO
PIROLIZO Z LO^ENO CONO IZHLAPEVANJA
Peter Majeri~1, Bernd Friedrich2, Rebeka Rudolf1,3
1University of Maribor, Faculty of Mechanical Engineering, Smetanova ulica 17, 2000 Maribor, Slovenia
2RWTH Aachen University, IME Process Metallurgy and Metal Recycling, Germany
3Zlatarna Celje d.d., Kersnikova ulica 19, 3000 Celje, Slovenia
peter.majeric@um.si
Prejem rokopisa – received: 2014-10-17; sprejem za objavo – accepted for publication: 2014-11-11
doi:10.17222/mit.2014.264
Some experiments were conducted in connection with the gold-nanoparticle production in an effort to produce a more accurate
model for determining the gold-nanoparticle synthesis with a modified ultrasonic spray pyrolysis (USP). As previous experi-
ments with gold nanoparticles yielded nanoparticles of various shapes (spherical, triangular, cylindrical, etc.), a focus on
synthesizing only spherical nanoparticles is underway, as a mixture of different shapes is difficult to characterize and utilize.
One of the factors for the particle formation is droplet evaporation. In an attempt to produce the optimum conditions for droplet
evaporation with the solvent diffusion and precipitation, a separate furnace and a separate reaction-gas inlet were used. This
modification separates the evaporation stage from the reaction stage, compared to the standard USP set-up. Using relatively low
temperatures, from 60 °C up to 140 °C, for the evaporation stage provides for more time and better conditions for the diffusion
of the solvent into the center of a droplet and a higher probability of forming a spherical particle.
Keywords: ultrasonic spray pyrolysis, gold nanoparticles, aerosol droplet evaporation, TEM microscopy
Prispevek opisuje sintezo nanodelcev zlata, ki smo jo izvedli z modificirano ultrazvo~no razpr{ilno pirolizo (USP) s ciljem
postavitve ~im bolj realnega teoreti~nega modela. Pri prej{njih eksperimentih sinteze nanodelcev zlata smo dobili nanodelce
razli~nih oblik (sferi~ne, trikotne, cilindri~ne, itd.), zato smo se pri tej raziskavi osredinili na izdelavo sferi~nih nanodelcev.
Morfolo{ko razli~ni nanodelci so ovira ne samo pri karakterizaciji, ampak tudi kasneje pri sami uporabi. Pri sintezi nanodelcev
ima velik vpliv na kon~no morfologijo izhlapevanje kapljic aerosola. Za zagotovitev razmer za izhlapevanje kapljic, ki omogo-
~ajo nastanek sferi~nih nanodelcev, smo v procesu USP uporabili lo~eno ogrevalno obmo~je in lo~en vnos reakcijskega plina. Ta
modifikacija namre~ lo~uje stopnjo izhlapevanja od stopenj reakcij za nastanek nanodelcev. Ni`ja temperatura v fazi izhlape-
vanja (med 60 °C in 140 °C) omogo~a bolj{e razmere za difuzijo topljenca do sredi{~a kapljice in ve~jo verjetnost nastanka
sferi~nih nanodelcev.
Klju~ne besede: ultrazvo~na razpr{ilna piroliza, zlati nanodelci, izhlapevanje kapljic aerosola, TEM-mikroskopija
1 INTRODUCTION
Spray pyrolysis is a well-known method for obtaining
nanoparticles of various shapes and sizes from a wide
variety of materials. This is a process that creates pow-
ders or suspensions from droplets of a solution with a
dissolved selected material1,2. There are variants of this
process in terms of droplet generation, droplet-transpor-
tation orientation and final nanoparticle collection. How-
ever, the process is always the same and follows these
steps: droplet generation, droplet transportation into the
heating zone, particle formation from the droplet, parti-
cle collection.
Ultrasonic spray pyrolysis (USP) uses ultrasound to
produce aerosol droplets from a precursor solution, a
carrier gas for the transportation of the aerosol to the
reactor furnace, a reaction gas for the formation of nano-
particles (if needed), a reactor furnace for the nanoparti-
cle formation and a collection system for nanoparticles.
The principle of nanoparticle production lies in the use
of heavily diluted precursor solutions and the creation of
small droplets. Ultrasound produces very small aerosol
droplets in the range of a few micrometers in diameter
(from 1 μm to 10 μm), or more, depending on the ultra-
sound frequency and the parameters of the used solu-
tion3. Using small solution concentrations the micron-
sized droplets are evaporated and the quantity of the
dried solute particles left behind is in the nano-range. By
Materiali in tehnologije / Materials and technology 49 (2015) 5, 791–796 791
UDK 669.21:620.3:537.533.35 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)791(2015)
Figure 1: Schematic of the USP appliance previously used for the
nanoparticle synthesis
Slika 1: Shema naprave USP, uporabljene za sintezo nanodelcev
increasing or decreasing the solute concentration inside
the droplet, greater or smaller nanoparticles can be
acquired. In theory, every aerosol droplet produces one
nanoparticle. This is known as the one-particle-per-drop-
let mechanism2.
For the gold-nanoparticle production with USP we
used a solution of hydrogen tetrachloroaurate (HAuCl4)
with nitrogen as the carrier gas for the aerosol droplets
and hydrogen for the reduction of dried gold-chloride
particles into gold nanoparticles. In the previous publi-
cations we reported synthesizing gold nanoparticles of
different sizes and differing shapes such as irregular,
spherical, triangular and cylindrical size4. We also no-
ticed some unexpected processes, such as the formation
of nanoparticles inside the ultrasonic aerosol generator.
This caused a change in the color of the initial precursor
solution from a clear yellow to a dark brown solution
after the synthesis.
The formation of gold nanoparticles occurs at rela-
tively low temperatures during spray pyrolysis (theore-
tically, a decomposition of HAuCl4 begins at 258 °C)5.
This means that the conditions of the aerosol formation
via ultrasound and the presence of the reduction hydro-
gen gas inside the generator may cause the nanoparticles
to start forming inside the generator (Figure 1). A solu-
tion was proposed for this problem, where the reduction
gas would be introduced at a later stage when the aerosol
is already formed. This also provides for an opportunity
to connect a separate evaporation stage to the current
synthesis set-up.
The separate evaporation stage would ensure the for-
mation of dried spherical particles of gold chloride
before the reaction stage involving hydrogen. The sepa-
rate heating zone (the evaporation stage) could accom-
plish this using the low temperatures required for a slow
evaporation of the solvent. Theoretically, when the sol-
vent (water) evaporation is slow enough, the solute (gold
chloride) will diffuse into the center of a droplet, form-
ing a dried spherical particle (Figure 2)6 and, without the
reaction gas present at this stage, the particles of gold
chlorides cannot form on the surfaces of the droplets.
After the evaporation, hydrogen gas is introduced
into the next heating zone (the reaction stage) and dried
particles are reduced from gold chloride to gold. The
particles are then collected for analysis. The dried-nano-
particle form can also be an indicator of the final nano-
particle form (Figure 2). For this reason, spray drying
was utilized, with which we acquired dried particles only
from HAuCl4 solution droplets. These experiments are
only discussed briefly in this article. There are descrip-
tions of several models of the solute precipitation from a
drying solution droplet6, where the predicted dried-parti-
cle form depends largely on the solution concentration,
the droplet size, the evaporation rate and the solubility of
the solute, while other parameters, such as the relative
humidity and the aerosol-droplet quantity, do not affect
the dried-particle formation so much.
In our research for the application of the produced
gold nanoparticles in medicine it was shown that the
sizes of around 50 nm nominally lead to the best results
regarding cytotoxicity and biocompatibility7. From a
theoretical standpoint this means that low concentrations
of gold should be used in the precursor solution. Low
concentrations of the material in the precursor solution
for the nanoparticle production also provide better con-
ditions for a solid, spherical nanoparticle formation8.
The separate evaporation zone with USP allows a
better control of the particle formation during the first
stages of the synthesis process. This, in turn, allows a
production of more uniformly spherical nanoparticles,
which are better suited for the applications in medicine
compared to the mixture of nanoparticles of spherical,
triangular and cylindrical shapes.
2 MATERIALS AND METHODS
For the aerosol-droplet generation we used an ultra-
sonic nebulizer Gapusol 9001 RBI (France), with an
ultrasonic frequency of 2.5 MHz. The carrier gas, nitro-
gen, was passed through the nebulizer into the tubes
leading towards the furnaces. Two tube furnaces were
used; the first one was kept at the lower temperatures
required for the evaporation, while the second one was
kept at the higher temperatures, required for the nano-
particle-synthesis reactions. Hydrogen gas was released
into the tubes between the furnaces to prevent a prema-
ture nanoparticle reduction before the droplets were
completely evaporated (Figure 3). The washing bottles
were filled with water for particle collection.
P. MAJERI^ et al.: Au-NANOPARTICLE SYNTHESIS VIA ULTRASONIC SPRAY PYROLYSIS ...
792 Materiali in tehnologije / Materials and technology 49 (2015) 5, 791–796
Figure 3: Schematic of the modified USP appliance used for experi-
ments
Slika 3: Shema spremenjene naprave USP za izvedene eksperimente
Figure 2: Simple presentation of the evaporation of an aerosol droplet
with solute precipitation
Slika 2: Enostavna predstavitev izhlapevanja kapljice aerosola in izlo-
~anje topljenca
The precursor solution in the aerosol generator was
kept at 21 °C with a thermostat. The generated aerosol
was carried into the evaporation heating zone via the
quartz tubes with a diameter of 20 mm by nitrogen gas at
a flow rate of 1.5 L/min. The residence time of the drop-
lets/dried particles in the evaporation zone was 3.77 s.
Both heating zones were 30 cm long, with a 5 cm gap
between them for the reduction-gas inlet. The reduction
gas, hydrogen, was passed into the tubes between the
heating zones at a flow rate of 1 L/min. The residence
time of the particles inside the second heating zone was
2.26 s.
The temperature profile of the two furnaces was
measured prior to the nanoparticle synthesis using only
water for the aerosol generation and nitrogen gas at both
gas inlets (Figure 4).
In order to determine the feasibility of dried-nano-
particle collection a few experiments were performed
with only the first furnace and without hydrogen gas
(Figure 5). The aim of these experiments was to deter-
mine the morphology of dried particles from aerosol
droplets before the final nanoparticle reactions. The
particles were collected on a filter paper as they were
assumed to be the particles of gold chloride. As gold
chloride is highly soluble it would not be possible to
collect it in wash bottles containing water, alcohol, or
similar.
The experiments performed are presented in Table 1.
TEM images with an EDS analysis were used for the
nanoparticle characterization.
3 RESULTS AND DISCUSSION
A number of thermal measurements were performed
to ensure the best possible conditions for the particle
formation. The temperatures measured were from 60 °C
to 140 °C for the first furnace and from 250 °C to 400 °C
for the second furnace. At the very low temperature sett-
ings (60–80 °C) for the first furnace, more condensation
was observed in the tubes, resulting in an aerosol droplet
loss with the synthesis. To avoid the needless conden-
sation, the temperature at which no condensation was
observed in the first furnace was selected. Condensation
of aerosol droplets reduces the nanoparticle output.
A reasonable temperature was selected for the second
furnace, where the temperature has to be high enough to
allow for the gold-nanoparticle formation while ensuring
a complete reaction so that no unreacted particles exit in
the second furnace. The temperature profile of the heat-
ing zones is shown in Figure 5.
In the previous experiments of the nanoparticle syn-
thesis with USP gold nanoparticles of different shapes
were obtained without a separate evaporation stage. The
collected nanoparticles were of spherical, irregular,
triangular and cylindrical shapes4. In comparison, sepa-
rating the evaporation stage from the reaction and densi-
fication stages yielded mostly non-ideally spherical
nanoparticles and some irregularly shaped nanoparticles.
Lower concentrations of gold in the precursor yielded
less irregular shapes of the nanoparticles.
The average nanoparticle sizes were also reduced by
separating the droplet evaporation from the rest of the
process. Using a precursor solution with 2.5 g/L Au and
an ultrasound frequency of 2.5 MHz, sizes ranging from
10–250 nm were obtained without a separate evaporation
zone4 (Table 2). With a separate evaporation zone, sizes
ranging from 10–80 nm were obtained (Figure 6 and
Table 2). It seems there are less agglomeration and par-
P. MAJERI^ et al.: Au-NANOPARTICLE SYNTHESIS VIA ULTRASONIC SPRAY PYROLYSIS ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 791–796 793
Figure 5: Schematic of the device for dried-nanoparticle-collection
experiments
Slika 5: Shema naprave za preizkus zbiranja posu{enih nanodelcev
Figure 4: Temperature profile of the heating zone for aerosol-droplet
evaporation (left side, before the gas inlet) and nanoparticle reactions
(right side, after the gas inlet)
Slika 4: Temperaturni profil ogrevanega obmo~ja za izhlapevanje kap-
ljic aerosola (leva stran grafa – pred vstopom plina) in reakcije nano-
delcev (desna stran grafa – po vstopu plina)
Table 1: Experiments performed with the USP device
Tabela 1: Eksperimenti, izvedeni z USP-napravo
Experi-
ment Type
Au solution
concentration
(g/L)
Collection
medium
1 Droplet evaporation 0.5 Filter paper
2 Droplet evaporation 1.25 Filter paper
3 Droplet evaporation 2.5 Filter paper
4 Nanoparticle synthesis 0.5 Water
5 Nanoparticle synthesis 1.25 Water
6 Nanoparticle synthesis 2.5 Water
7 Nanoparticle synthesis 5 Water
ticle growth present when the nanoparticles are synthe-
sized with a separate evaporation zone.
Table 2: Comparison of nanoparticle sizes after USP synthesis with
and without a separate evaporation zone
Tabela 2: Primerjava velikosti nanodelcev pri sintezi z USP brez lo~e-
nega obmo~ja za izhlapevanje in z njim
USP
Gold concentration
in precursor
solution
Synthesized
gold-nanoparticle
size distribution
Without a separate
evaporation zone 2.5 g/L 10–250 nm
With a separate
evaporation zone 2.5 g/L 10–80 nm
The sizes of the nanoparticles are generally accepted
to be mainly dependent upon the precursor concentration
and droplet size. In our experiments, the droplet sizes are
considered to be the same in each experiment, while the
precursor concentration was changed for each experi-
ment. With a gold concentration of 5 g/L in the precursor
solution we obtained nanoparticles in the size range from
around 50 nm up to 300 nm. While there are small nano-
particles present, the bulk of the nanoparticles have a
diameter of about 200 nm. There are also a bigger num-
ber of irregularly shaped nanoparticles, with less sphe-
rical nanoparticles present. More agglomeration is also
observed when using such a precursor solution concen-
tration. The separate evaporation zone seems to be
beneficial for the production of more uniformly shaped
nanoparticles as the irregularly shaped nanoparticles
have round edges even though they are not spherical.
When not using the separate evaporation zone, nanopar-
ticles have sharper edges in the form of triangles and
discs and irregular shapes with sharp edges4. Such sha-
pes did not appear in the case of the separate evaporation
zone. The separate evaporation zone is, therefore,
favorable in order to obtain only spherical shapes of the
nanoparticles produced with USP.
With a gold concentration of 2.5 g/L in the precursor
solution, we obtained sizes ranging from 10 nm to 80 nm.
The average size of the nanoparticles with such a con-
centration is 55 nm. The shapes of the nanoparticles are
still not ideally spherical as there are still particles with
irregular shapes present. However, even though they are
not ideally spherical, their shapes appear to have a high
order of sphericity, with the smaller nanoparticles being
more spherical than the bigger ones.
A gold concentration of 1.25 g/L in the precursor
solution yielded nanoparticles in the range from 10 nm
to 40 nm. The bulk of the nanoparticle sizes are about
20 nm. With a precursor concentration of 0.5 g/L of gold
in the solution the nanoparticle sizes obtained were from
10 nm to 40 nm, with the average nanoparticle size being
somewhat smaller, around 15 nm.
The shapes of the nanoparticles with gold concentra-
tions of (2.5, 1.25 and 0.5) g/L in the precursor solution
are very similar. With all three concentrations we ob-
tained a mixture of spherical and irregularly shaped
nanoparticles (Figures 6 and 7).
The separate evaporation zone yielded better results
regarding the uniformity of the shapes and sizes with
higher gold concentrations (2.5 g/L and 5 g/L) in the pre-
cursor. Even though the shapes are still not ideally uni-
form, less diverse shapes are produced while using the
separate evaporation zone in the USP process. The sizes
of the obtained nanoparticles are also less dispersed. In
this regard, the separated evaporation zone is useful for
producing gold nanoparticles of more similar shapes and
sizes with USP. This is not apparent when using lower
gold concentrations (0.5 g/L and 1.25 g/L). With lower
concentrations, solid spherical shapes are formed more
readily with USP, and the benefits of the separate evapo-
ration zone are not seen so easily (Figure 7). From a
theoretical standpoint, using a low concentration in the
precursor solution is effective for producing solid and
more uniform nanoparticles with the standard USP
P. MAJERI^ et al.: Au-NANOPARTICLE SYNTHESIS VIA ULTRASONIC SPRAY PYROLYSIS ...
794 Materiali in tehnologije / Materials and technology 49 (2015) 5, 791–796
Figure 7: TEM image of gold nanoparticles obtained with USP with a
separate evaporation zone, gold concentration in precursor solution:
0.5 g/L
Slika 7: TEM-posnetek nanodelcev zlata, pridobljenih z USP z
lo~enim obmo~jem za izhlapevanje; koncentracija zlata v za~etni
raztopini: 0,5 g/L
Figure 6: TEM image of gold nanoparticles obtained with USP with a
separate evaporation zone; gold concentration in precursor solution:
2.5 g/L
Slika 6: TEM-posnetek nanodelcev zlata, pridobljenih z USP z
lo~enim obmo~jem za izhlapevanje; koncentracija zlata v za~etni
raztopini: 2,5 g/L
process. Even though spherical shapes of the
nanoparticles were formed in our experiments with the
modified USP process, irregular shapes were also
present.
As a result, it is difficult to determine the effective-
ness of the separate evaporation zone when using lower
concentrations of gold in the precursor solution. More
work is needed with regard to the other synthesis para-
meters in order to determine the feasibility of using the
separate evaporation zone when using lower gold con-
centrations in the precursor solution. However, from a
technological point of view, the separate evaporation
zone eliminates the reactions with hydrogen and the par-
ticle formation inside the ultrasound generator. The pre-
cursor solution remained the same before and after the
nanoparticle synthesis, which is a good indicator of this
technical issue being resolved, compared to the standard
USP set-up where the precursor solution changes after
the experiment.
The separate evaporation zone used in our experi-
ments gave some favorable results with respect to
reducing the nanoparticle-size distribution and producing
more uniformly shaped nanoparticles. In order to ad-
vance the research in the direction of producing uniform
shapes, other parameters, such as the gas flow, should be
studied experimentally, or further modifications should
be made to the USP process.
In the experiments with the droplet evaporation the
produced particles were collected on a filter paper made
from cellulose in order to prevent a re-dissolving of the
gold chloride particles. This collection yielded two sub-
stances, or types of particles, visible on the filter paper
after the collection. On the edges of the used filter paper
a purple substance was collected, while at the center of
the filter paper we collected a yellowish substance. This
suggests that two chlorides of gold were collected: AuCl
appeared in the form of yellow crystals, while AuCl3
appeared in the form of red crystals9.
It is assumed that the nanoparticle morphology is
dependent mainly upon the evaporation stage. The opti-
mum drying of the aerosol droplets theoretically pro-
duces solid, spherical particles of the dried solute. In
order to check the feasibility of dried-nanoparticle
collection after the evaporation, the particles were first
collected on filter paper (Figure 8). The particles
collected in this way are difficult to characterize and an
improved method for collecting dried nanoparticles will
be utilized in future research for an easier particle
characterization after the collection.
After the collection of dried particles, the correlation
of dried particles with the final nanoparticles is
examined. This insight into the separate synthesis stages
provides a better understanding of the process and allows
us to produce a more accurate model for determining the
gold-nanoparticle synthesis with a separate evaporation
zone.
4 CONCLUSION
Gold nanoparticles were synthesized using USP with
a separate evaporation zone. The gold concentration in
the precursor solution used for the aerosol generation
was the primary variable parameter in order to determine
the effects of the separate evaporation zone on the final
nanoparticle sizes and shapes. When using higher con-
centrations of gold in the precursor solution, the experi-
ments showed more uniformly shaped nanoparticles in
the case of the separate evaporation zone compared to
the standard USP set-up. A narrower size distribution of
the gold nanoparticles was also obtained in comparison
to the standard USP set-up. When using lower gold con-
centrations in the precursor solution, these advantages
are not so obvious. As the separate evaporation zone
does not entirely prevent the presence of irregularly
shaped nanoparticles, similar nanoparticles can be ob-
tained without a separate evaporation zone at low precur-
sor-solution concentrations. However, it is apparent that
separating the evaporation stage from the reaction and
densification stages improved the process to some extent
so that more nanoparticles of spherical shapes and more
uniform sizes are produced. The produced nanoparticles
are smaller in size (10–80 nm for 2.5 g/L Au in the pre-
cursor solution) compared to the nanoparticles produced
without a separate evaporation zone (10–250 nm for 2.5
g/L Au in the precursor solution). Further work on the
modifications and parameter optimization is now needed
to improve the process in pursuit of gold nanoparticles
with uniform shapes and as narrow a nanoparticle-size
distribution as possible.
Acknowledgment
We would like to thank Dr. Darja Jenko for per-
forming the TEM analysis on the Au nanoparticles. This
research was made possible by the Slovenian Research
Agency (ARRS), the Young Researcher Program and the
P. MAJERI^ et al.: Au-NANOPARTICLE SYNTHESIS VIA ULTRASONIC SPRAY PYROLYSIS ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 791–796 795
Figure 8: Nanoparticles collected on filter paper after evaporation,
created with a gold concentration of 1.25 g/L in the precursor solution
Slika 8: Nanodelci, zbrani na filter papirju po izhlapevanju, narejeni s
koncentracijo zlata 1,25 g/L v za~etni raztopini
Slovenian Public Fund for the Development of Human
Resources.
5 REFERENCES
1 G. L. Messing, S. C. Zhang, G. V. Jayanthi, Ceramic Powder Syn-
thesis by Spray Pyrolysis, J. Am. Ceram. Soc., 76 (1993) 11,
2707–2726, doi:10.1111/j.1151-2916.1993.tb04007.x
2 S. C. Tsai, Y. L. Song, C. S. Tsai, C. C. Yang, W. Y. Chiu, H. M. Lin,
Ultrasonic spray pyrolysis for nanoparticles synthesis, J. Mater. Sci.,
39 (2004) 11, 3647–3657, doi:10.1023/B:JMSC.0000030718.76690.
11
3 J. Bogovi}, A. Schwinger, S. Stopic, J. Schröder, V. Gaukel, H. P.
Schuchmann, B. Friedrich, Controlled droplet size distribution in
ultrasonic spray pyrolysis, Metall, 65 (2011) 10, 455–459
4 S. Stopic, R. Rudolf, J. Bogovic, P. Majeric, M. Colic, S. Tomic, M.
Jenko, B. Friedrich, Synthesis of Au nanoparticles prepared by
ultrasonic spray pyrolysis and hydrogen reduction, Mater. Tehnol.,
47 (2013) 5, 577–583
5 R. Dittrich, S. Stopic, B. Friedrich, Mechanism of nanogold forma-
tion by ultrasonic spray pyrolysis, Proceeding of the EMC Confe-
rence, Düsseldorf, Germany, 2011, 385
6 G. V. Jayanthi, S. C. Zhang, G. L. Messing, Modeling of Solid Par-
ticle Formation During Solution Aerosol Thermolysis: The Evapo-
ration Stage, Aerosol Sci. Technol., 19 (1993) 4, 478–490,
doi:10.1080/02786829308959653
7 S. Tomi}, J. Ðoki}, S. Vasiliji}, N. Ogrinc, R. Rudolf, P. Pelicon, D.
Vu~evi}, P. Milosavljevi}, S. Jankovi}, I. An`el, J. Rajkovi}, M. S.
Rupnik, B. Friedrich, M. ^oli}, Size-Dependent Effects of Gold
Nanoparticles Uptake on Maturation and Antitumor Functions of
Human Dendritic Cells In Vitro, PLoS ONE, 9 (2014) 5, e96584,
doi:10.1371/journal.pone.0096584
8 T. T. Kodas, M. J. Hampden-Smith, Aerosol Processing of Materials,
1st edition, Wiley-VCH, New York 1998
9 D. R. Lide (Ed.), CRC Handbook of Chemistry and Physics, 80th
Edition, CRC Press, Boca Raton 1999
P. MAJERI^ et al.: Au-NANOPARTICLE SYNTHESIS VIA ULTRASONIC SPRAY PYROLYSIS ...
796 Materiali in tehnologije / Materials and technology 49 (2015) 5, 791–796
R. KAYIKCI et al.: DETERMINATION OF THE CRITICAL FRACTION OF SOLID ...
797–800
DETERMINATION OF THE CRITICAL FRACTION OF SOLID
DURING THE SOLIDIFICATION OF A PM-CAST ALUMINIUM
ALLOY
DOLO^ANJE KRITI^NEGA DELE@A STRJENE FAZE MED
STRJEVANJEM V PM ULITE ALUMINIJEVE ZLITINE
Ramazan Kayikci1, Murat Colak2, Selcuk Sirin1, Engin Kocaman1, Neset Akar3
1Sakarya University, Faculty of Technology, Dept. of Met. Mat. Eng., Sakarya, Turkey
2Bayburt University, Faculty of Eng., Dept. of Mat. Sci. and Nanotech. Eng., Bayburt, Turkey
3Gazi University, Faculty of Technology, Dept. of Met. Mat. Eng., Ankara, Turkey
rkayikci@sakarya.edu.tr
Prejem rokopisa – received: 2014-10-20; sprejem za objavo – accepted for publication: 2014-11-27
doi:10.17222/mit.2014.266
During the solidification of aluminium alloys, a flow of the interdendritic liquid through the mushy zone plays an important role
in the feedability and shrinkage-driven porosity formation. Therefore, for almost all the studies about casting modelling, a good
knowledge of the mushy-zone behaviour is necessary.
In this study, the limit of the mushy-zone permeability for permanent-mould (PM)-cast aluminium alloys was investigated. A
novel experimental technique was developed to quantify the value of the critical fraction of solid (CFS) at which the
interdendritic feeding stops. Grain-refined and non-grain-refined aluminium alloys were used to evaluate the effect of the grain
size on the CFS limit of both the non-grain-refined and fully grain-refined alloys. Three different initial mould temperatures
were used to find if the mould temperature has any effect on the feeding of a casting. Alloys of varying solidification intervals
were also used to determine the critical fraction of solid (CFS) for the feeding limit of solidifying castings.
The results showed that the CFS values for solidifying aluminium alloys may vary from 34 % to 57 % depending on the casting
conditions. The results also showed that for A356 aluminium-alloy castings the CFS value is not an alloy property, yet it
significantly depends on the parameters such as the grain refinement and the modification of eutectic silicon.
Keywords: critical fraction of solid, mushy-zone permeability, PM casting
Med strjevanjem aluminijevih zlitin ima tok meddendritne taline skozi testasto podro~je pomembno vlogo pri napajanju in
nastajanju poroznosti zaradi kr~enja. Zato je skoraj v vseh modelih ulitkov potrebno dobro poznanje vedenja testastega
podro~ja.
V tej {tudiji je bila iskana meja prepustnosti testastega podro~ja pri aluminijevih zlitinah, ulitih v stalne forme (PM). Razvita je
bila nova eksperimentalna tehnika za oceno kriti~nega dele`a strjene faze (CFS), pri kateri se meddendritno napajanje ustavi.
Uporabljene so bile aluminijeve zlitine z udrobnjenimi in neudrobnjenimi zrni za oceno vpliva velikosti zrn na mejno CFS pri
obeh vrstah zlitin. Izmerjene so bile tri razli~ne temperature kokil, da bi ugotovili, ali temperatura kokile vpliva na napajanje
ulitka. Uporabljene so bile tudi zlitine z razli~nim intervalom strjevanja, da bi ugotovili kriti~en dele` trdne faze (CFS), ki
pomeni mejo napajanja pri strjevanju ulitkov.
Rezultati so pokazali, da vrednosti CFS variirajo med 34 % in 57 %, odvisno od razmer pri ulivanju. Rezultati so tudi pokazali,
da pri ulitkih iz aluminijeve zlitine A356 vrednost CFS ni odvisna samo od zlitine, temve~ je zelo odvisna od parametrov, kot
sta zmanj{anje zrn in modifikacija evtekti~nega silicija.
Klju~ne besede: kriti~en dele` trdne faze, prepustnost testastega podro~ja, ulivanje v stalne forme (PM)
1 INTRODUCTION
The solidification of aluminium alloys generally
starts with dendrites of a aluminium and they continue to
grow into a coherent network. With the progress of soli-
dification, several intermetallics might precipitate and
the solidification usually completes with one or more
eutectic phases.1 While the solidification continues, the
viscosity of liquid aluminium alloys increase with a
decrease in the fluidity. During the solidification, alumi-
nium alloys pass through a long mushy region during
which the liquid requirement for a contraction of soli-
difying dendrites can be compensated with a flow of the
interdendritic liquid.2 These mechanisms continue until a
certain stage of the solidification in the mushy zone.
Dendrite coherency is established in a later stage and the
dendrite network gains some rigidity which becomes
resistant against the flow of the interdendritic liquid.3
This phenomenon continues until the liquid flow stops
completely due to the blockage of the growing solid den-
drites.3 In the literature about casting solidification, this
stage of the solidification is defined as either šthe critical
fraction of solid’ or šthe critical fraction of liquid’.4
At the critical fraction of the solid (CFS) stage, if
there is still a need for the liquid flow for feeding the
solidifying liquid, a pressure drop is created due to the
resistance to the liquid flow. If this negative pressure in
the mushy zone cannot be balanced by the flow of the
liquid due to the dendrite blockage it may lead to some
shrinkage porosity.5 Shrinkage cavities that can often be
seen in casting processes are typical examples of this
porosity. Today high-performance technology requires
Materiali in tehnologije / Materials and technology 49 (2015) 5, 797–800 797
UDK 669.715:621.74:536.421.4 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)797(2015)
high-quality and light-section castings and therefore
cannot even tolerate micro-porosity6. To avoid the forma-
tion of micro- and macro-porosities within the critical
sections of cast parts, a variety of 3D casting-modelling
software packages, often referred to as casting simula-
tion packages, have been used during the design stages
of casting processes. Since the CFS point is practically
the end of the feeding action within a mushy region, it
has been used as one of the most critical boundary con-
ditions in all the commercial casting-simulation software
packages. Therefore, to model aluminium casting pro-
cesses with different alloys and cooling-solidification
conditions, a good knowledge about the critical fraction
of solid (CFS) is necessary.
In this study, using commercial A356 aluminium
casting alloys, the CFS point of the mushy zone was
investigated. Permanent mould castings with three diffe-
rent alloy conditions, namely, the as-cast, the grain-
refined, and the grain-refined and modified conditions
were used. Casting-simulation techniques, X-ray radio-
graphy, the Archimedes density measurement and metal-
lographic techniques were also employed.
2 EXPERIMENTAL WORK
The experimental part of the present study is twofold.
In the first part an A356 aluminium alloy was cast in a
steel mould and the resulting shrinkage was examined
for varying casting conditions. In the second part, the
casting and mould geometries were modelled using the
SOLIDCast 3D casting-simulation software. Finally, the
amounts of the shrinkage porosity obtained from the
actual castings were compared with the results from the
computer modelling and the feeding threshold value (the
critical fraction of solid) was determined for each casting
condition.
The casting geometry and the solidified cast part
within the mould are shown in Figures 1a and 1b,
respectively. The casting geometry was specially
designed to obtain two thick sections connected to each
other with a narrow feeding neck to create an insufficient
feeding path so that it could produce a measurable
amount of the shrinkage porosity at the bottom part of
the casting, used for comparison and evaluation.
A chemical analysis of the A356 alloy was carried
out with a Foundry-Master optical emission spectrometer
before and after each addition to the master alloy as
shown in Table 1. Primary ingots of 8 kg were melted in
a SiC crucible in a bell-type electric-resistance-heated
furnace. The alloy was degassed with dry nitrogen at 720
°C for 6 min before being cast into a steel mould which
was preheated to 300 °C. The casting process was re-
peated and followed by an addition of an Al5Ti1B-type
grain refiner or an Al-10Sr modification of the master
alloy. The castings were radiographically tested on both
sides to check whether any shrinkage porosity had
occurred. Their upper and bottom parts were separated
by cutting them apart from the top of the narrow neck
and the cast density of the bottom parts was measured,
using the Archimedes system of weighing the parts in air
and in water. Finally, the castings were sectioned along
the vertical axis to reveal any shrinkage in the central
regions of the castings.
Table 1: Chemical composition of the A356 alloy used in the experi-
ments (Al balance), w/%
Tabela 1: Kemijska sestava zlitine A356, uporabljene pri preizkusih
(osnova je Al), w/%
Elements Fe Si Cu Mn Mg Ti Sr
As-cast 0.35 6.9 0.1 0.35 0.38 0.03 0
TiB modified 0.35 6.8 0.1 0.34 0.36 0.12 0
Sr modified 0.35 6.82 0.1 0.3 0.3 0.1 0.08
A commercial 3D casting-simulation package,
SOLIDCast, was used for modelling the castings. The
system solves the Fourier heat-transfer equation with the
finite-difference method and calculates the volume
change within the casting with respect to the tempe-
rature. It assumes that the feeding takes place under the
gravity effect and the volume of the total final shrinkage
798 Materiali in tehnologije / Materials and technology 49 (2015) 5, 797–800
R. KAYIKCI et al.: DETERMINATION OF THE CRITICAL FRACTION OF SOLID ...
Figure 2: Scanned view of the radiographic test films of A356 cast-
ings: a) as cast, b) with added TiB, c) with added TiB + Sr
Slika 2: Skenirani posnetki filmov rentgenskih posnetkov ulitkov iz
A356: a) lito stanje, b) dodan TiB, c) dodan TiB + Sr
Figure 1: a) Casting geometry, b) solidified casting within a perma-
nent mould
Slika 1: a) Geometrija ulitkov, b) strjen ulitek v stalni kokili
pores is equal to the volume of the unfed contracted
liquid metal. The system can allow the feed-liquid
movement as long as the feeding path is open, i.e., the
fraction of solid (FS) is below the critical value in the
mushy zone.
3 RESULTS AND DISCUSSION
A set of the radiographic test results representing the
as-cast, TiB-modified and TiB + Sr-modified castings is
shown in Figure 2. The potential shrinkage areas in the
bottom sections of the castings are highlighted with
circles. Dark smoky marks in the circled areas indicate
the shrinkage porosity. The largest shrinkage occurred in
the casting in the as-cast condition indicating the lowest
feeding action of this particular casting. This was asso-
ciated with a shorter CFS time due to an early blockage
of coarser dendrites. On the other hand, the casting that
was fully grain refined and modified with an addition of
TiB + Sr revealed the smallest amount of the shrinkage
porosity. This showed that the longest feeding action
took place during the solidification. Only the grain-
refined casting (TiB-modified only) showed an interme-
diate shrinkage porosity. These results are in good
agreement with the report on the grain refinement of an
aluminium alloy by Birol.7
Figure 3 shows macrostructures of the cross-sections
of the castings. The shrinkage patterns of the castings
reveal characteristics similar to the X-ray radiography
pattern shown in Figure 2. The percentages of the
shrinkage porosity of the castings were determined by
subtracting their measured Archimedes densities from
the theoretical density of a well-fed PM cast of the A356
alloy. The results obtained from the simulations of the
castings are shown in Figure 4. This figure shows that
increasing the CFS values of the castings resulted in
longer feeding times as the time required to reach a given
CFS value of the mushy zone also increased. Modelled
porosities were obtained directly from the simulation
results. These values are given in Table 2. The corres-
ponding CFS values for different casting conditions are
also given in Table 2. A number of simulations were run
to change the CFS point of a casting as a boundary con-
dition from 30 % to 70 % and their resulting porosity
values were compared with the measured porosity values
of the actual castings. Determination of the actual CFS
value of each casting was carried out by comparing the
measured and modelled densities. When a close match
was obtained between the two, the corresponding value
of the CFS that was already used as an input boundary
condition in that particular casting simulation was
assumed to be the right value of the CFS. As seen in
Table 2, the lowest CFS value of 34 % and the shorter
feeding time were obtained for the PM-cast A356 alloy.
The CFS value significantly increases with the grain
refinement achieved with the TiB addition to the melt,
reaching 52 %. The modification of the eutectic silicon
with Sr provides further improvement, causing the CFS
to rise to 57 %.
Table 2: Comparison of measured and modelled porosity fraction of
the castings with their CFS values
Tabela 2: Primerjava izmerjenih in z modelom dolo~enih dele`ev po-
roznosti ulitkov z vrednostmi CFS
Experimental
castings Modelling
Alloy
condition
Measured
porosity (%)
Calculated
porosity (%) CFS value (%)
As-cast 2.25 2.22 34
TiB modified 1.83 1. 82 52
TiB + Sr
modified 1.36 1.36 57
4 CONCLUSIONS
• An experimental casting technique coupled with a
casting simulation was studied to determine the CFS
values of the commercial A356 alloy in the as-cast,
TiB-modified and TiB + Sr-modified conditions.
• The results showed that the CFS point of the mushy
zone is not a constant value as it might vary signi-
ficantly with different conditions of the cast alloy.
R. KAYIKCI et al.: DETERMINATION OF THE CRITICAL FRACTION OF SOLID ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 797–800 799
Figure 4: Representative results obtained with a casting simulation
with four different CFS conditions
Slika 4: Zna~ilni rezultati, dobljeni s simulacijo ulivanja pri {tirih raz-
li~nih CFS-razmerah
Figure 3: Cross-sections of A356 castings: a) as cast, b) with added
TiB, c) with added TiB + Sr
Slika 3: Prerez ulitka iz A356: a) lito stanje, b) dodan TiB, c) dodan
TiB + Sr
• The lowest CFS value of 34 % was obtained for the
as-cast A356 alloy, while it increased up to 57 % with
a grain refinement and modification of the eutectic
silicon by adding TiB + Sr to the alloy.
Acknowledgement
The authors are grateful to The Scientific and
Technological Research Council of Turkey (TÜBÝTAK)
for providing a research grant to the 112M422 project.
5 REFERENCES
1 M. C. Flemings, Solidification Processing, McGraw-Hill, New York
1974
2 B. O. Danylo, Ö. N. Doðan, An examination of effects of solidifi-
cation parameters on permeability of a mushy zone in castings,
Journal of Materials Science, 43 (2008), 1471–1479, doi:10.1007/
s10853-007-2325-z
3 Ø. Nielsen, S. O. Olsen, Experiment for Quantification of Feedability
and Permeability in Industrial Aluminium Alloys, International
Journal of Cast Metals Research, 19 (2006) 1, 32–37, doi:10.1179/
136404606225023336
4 D. Sun, S. V. Garimella, Numerical and Experimental Investigation
of Solidification Shrinkage, Numerical Heat Transfer, Part A:
Applications, 52 (2007) 2, 145–162, doi:10.1080/1040778060
1115079
5 Ch. Pequet, M. Gremaud, M. Rappaz, Modelling of Microporosity,
Macroporosity, and Pipe Shrinkage Formation during the Solidifica-
tion of Alloys Using a Mushy-Zone Refinement Method: Applica-
tions to Aluminum Alloys, Metallurgical and Materials Transactions,
33A (2002) 7, 2095–2106, doi:10.1007/s11661-002-0041-5
6 M. Tiryakioðlu, J. Campbell, Guidelines for Designing Metal Cast-
ing Research: Application to Aluminium Alloy Castings, Internatio-
nal Journal of Cast Metals Research, 20 (2007) 1, 25–29,
doi:10.1179/136404607X186509
7 Y. Birol, Grain Refinement in Aluminium Shape Casting, Proc. of
the 3rd Int. Aluminium Conference, Arak, Iran, 2012, 1–12
R. KAYIKCI et al.: DETERMINATION OF THE CRITICAL FRACTION OF SOLID ...
800 Materiali in tehnologije / Materials and technology 49 (2015) 5, 797–800
F. TEHOVNIK et al.: MICROSTRUCTURAL EVOLUTION OF INCONEL 625 DURING HOT ROLLING
801–806
MICROSTRUCTURAL EVOLUTION OF INCONEL 625 DURING
HOT ROLLING
MIKROSTRUKTURNI RAZVOJ INCONELA 625 MED VRO^IM
VALJANJEM
Franc Tehovnik, Jaka Burja, Bojan Podgornik, Matja` Godec, Franci Vode
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
Prejem rokopisa – received: 2015-08-25; sprejem za objavo – accepted for publication: 2015-09-09
doi:10.17222/mit.2015.274
This research provides an overview of the structural changes that occur during the hot rolling of the nickel superalloy Inconel
625. It is well known that microstructure control is of paramount importance concerning the mechanical properties of a material.
The microstructure also plays an important role in processing materials at elevated temperature. In this work the hot-rolling
behaviour of the Inconel 625 superalloy has been investigated. The specimens were hot rolled at a temperature of 1200 °C using
different numbers of passes. During the hot rolling the loads were measured and recorded. A light microscope and an electron
microscope, employing the electron-backscatter-diffraction (EBSD) technique, were employed to investigate the microstructure
evolution, revealing a necklace dynamic-recrystallization mechanism.
Keywords: nickel superalloy, hot rolling, dynamic recrystallization, necklace mechanism
Raziskava obsega pregled strukturnih sprememb, ki se pojavijo med vro~im valjanjem nikljeve superzlitine Inconel 625. Znano
je, da je kontrola mikrostrukture odlo~ilnega pomena za mehanske lastnosti materiala. Mikrostruktura ima pomemben vpliv na
predelavo materiala pri povi{anih temperaturah. Preu~evali smo superzlitino Inconel 625 med vro~im valjanjem. Vzorci so bili
vro~e valjani pri temperaturi 1200 °C z razli~nimi {tevili prevlekov, med poizkusi smo merili tudi sile valjanja. Vzorce smo
preiskali s svetlobnim in elektronskim mikroskopom s tehniko difrakcije povratno sipanih elektronov (EBSD). Raziskali smo
razvoj mikrostrukture in potrdili dinami~no rekristalizacijo preko “mehanizma ogrlice”.
Klju~ne besede: nikljeva superzlitina, vro~e valjanje, dinami~na rekristalizacija, mehanizem ogrlice
1 INTRODUCTION
The Inconel nickel-chromium superalloy 625 is well
known for its high strength and outstanding corrosion
resistance. The strength of the Inconel alloy is derived
from the stiffening effect of the molybdenum and nio-
bium on its nickel-chromium matrix. This combination
of elements is also responsible for its superior resistance
to a wide range of corrosive environments of high seve-
rity as well as the high-temperature effects such as
oxidation and carburization. Because of these properties,
as well as fatigue and creep resistance, nickel-based
superalloys are widely used in modern aero engines and
gas turbines1,2.
During the hot deformation of metals and alloys, the
material flow behaviour is often very complex. In order
to optimize the final mechanical properties, control of
the microstructure is of great importance3. Studies show
that during hot deformation the work hardening, dyna-
mic recovery (DRV) and dynamic recrystallization
(DRX) often occur in the metals and alloys with a low
stacking-fault energy4–6. Generally, DRX is not only an
important softening mechanism, but also an effective
method to refine the crystal grain size. Therefore, in the
case of nickel-based alloys, which retain high strengths
at high processing temperatures, causing high rolling
loads, DRX plays a decisive role. The aim of the current
work was to investigate the microstructure evolution and
the nucleation mechanisms of the dynamic recrystalliza-
tion of Inconel superalloy 625 during hot rolling.
2 EXPERIMENTAL
2.1 Material
Flat specimens of Inconel 625, with the chemical
composition given in Table 1, were used for hot rolling
tests. The specimens were cut from a thick hot-rolled
plate; the specimens were 11 mm high, 145 mm long and
46 mm wide.
Table 1: Chemical composition of Inconel 625
Tabela 1: Kemijska sestava Inconela 625
w/% C N Ni Mo Nb Fe Cr Ti Si
Inconel
625 0.012 0.018 63.4 8.7 3.5 0.15 21.8 0.29 0.24
2.2 Hot rolling
Before the hot-rolling experiment the specimens were
heated to a temperature of 1200 °C and soaked for 30
min in order to obtain a homogeneous temperature field
of 1200 °C before each rolling experiment, regardless of
the number of passes used. Up to five rolling passes were
employed in this study. The logarithmic deformation per
Materiali in tehnologije / Materials and technology 49 (2015) 5, 801–806 801
UDK 669.245:621.77:620.186 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(5)801(2015)
pass was 0.223. The logarithmic deformation rate during
the rolling was calculated to be 7.12 s–1 for the first, 7.96
s–1 for the second, 8.90 s–1 for the third, 9.95 s–1 for the
fourth and 11.12 s–1 for the fifth pass. During the hot
rolling the rolling loads were monitored using a measu-
rement system with 2000 recordings per second, and
converted to stress by dividing the loads by the contact
area. The specimen temperature was measured at the
beginning and after each rolling pass. After the hot
rolling the rolled specimens were air cooled and pre-
pared for microscopic examination.
2.3 Microstructural analysis
Light microscopy, SEM and SEM-based electron-
backscatter-diffraction (EBSD) analyses were performed
on the metallographic samples to reveal the micro-
structure and microstructure evolution during the hot
rolling. The microstructure was examined optically using
a Nikon Microphot FXA, while the EBSD data were
collected in a FE-SEM JEOL JSM 6500F field-emission
scanning electron microscope using an HKL Nordlys II
EBSD camera and Channel 5 software. The EBSD
mapping analyses were performed with samples tilted by
70 ° using a 15 kV accelerating voltage and a 1.5 nA
probe current. The EBSD map was generated in steps of
2 μm using 619 × 503 grids with 4 × 4 binning and mini-
mum 5 and maximum 6 band detection and an indexing
rate between 92 % and 97 %. The samples were prepared
using a classic metallography preparation procedure,
including grinding, polishing and etching, and colloidal
silica polishing for the EBSD analysis.
3 RESULTS AND DISCUSSION
3.1 Hot rolling of Inconel superalloy
Although the degrees of deformation and the
deformation rates used in this study were relatively high,
no cracks or failures were observed at the edges of the
hot-rolled specimens. The samples after hot rolling and
air cooling are presented in Figure 1. However, the tem-
perature drop after each rolling pass was found to be
significant. The temperature at the end of first rolling
pass was 1130 °C, and the dropped to only 840 °C at the
end of the fifth pass.
The corresponding loads recorded during the five
rolling regimes (number of passes) are given in Figure 2.
As shown, the loads increase with the increasing number
of passes. The values recorded during the first pass were
found to be similar for all five regimes, which also
applies to all the other sequences. Furthermore, the tem-
peratures measured during each pass did not vary by
more than 15 °C between consecutive tests and the five
rolling regimes employed.
Nickel superalloys exhibit work hardening at high
strains. When comparing the hot rolling of the nickel
superalloy Inconel 625 to the superaustenitic steel 904L,
done in previous work7, the nickel superalloy shows a
greater deformation resistance. For this reason the initial
degree of deformation during the industrial processing of
nickel superalloy slabs has to be small. The deformation
degrees can be higher when the onset of recrystallization
begins at sufficiently high temperatures. However, spe-
cial care must be taken when processing as-cast micro-
structures. Otherwise, the ductility of nickel alloys is
very good and does not represent any concern. In the
present case the microstructure of the initial specimens
was already wrought, eliminating any concern about
defects occurring due to insufficient ductility.
Figure 3 shows the stress evolution during each
rolling pass for up to five passes. The corresponding
logarithmic strain rates and temperatures measured at the
end of each rolling pass are provided in the figure cap-
tion.
F. TEHOVNIK et al.: MICROSTRUCTURAL EVOLUTION OF INCONEL 625 DURING HOT ROLLING
802 Materiali in tehnologije / Materials and technology 49 (2015) 5, 801–806
Figure 2: Rolling loads with different numbers of rolling passes
Slika 2: Sile valjanja z razli~nimi {tevili prevlekov valjanja
Figure 1: Hot rolled and air cooled specimens of Inconel 625, sub-
jected to different numbers of rolling passes
Slika 1: Vro~e valjani in na zraku ohlajeni vzorci Inconela 625 z raz-
li~nimi {tevili prevlekov
The stress during the first rolling pass reached values
between 720 MPa and 900 MPa, while the specimen
temperature dropped from 1200 °C to 1130 °C. During
the second rolling pass the stress was between 790 MPa
and 1180 MPa and the temperature was 1040 °C. For the
third rolling pass the stress increased to 990 and 1360
MPa and final temperature dropped to 980 °C. During
the fourth and fifth rolling passes the stress was between
1180 MPa and 1710 MPa, and between 1610 MPa and
2000 MPa, respectively. On the other hand, the specimen
temperature decreased to 940 °C after the fourth, and
even down to 840 °C after the last, rolling pass. Due to
the elongation of the specimen the rolling time also
increased from 0.48 s for the first pass to 0.74 s for the
fifth pass.
The material’s resistance to deformation increases for
each rolling pass. This is the product of the decreased
temperature, increased strain rate and the accumulation
of strain. However, DRX did take place and partially
softened the material, as was revealed by the microstruc-
ture analysis given in Section 3.2. The dispersion of the
load data during rolling shown in Figure 2 is attributed
F. TEHOVNIK et al.: MICROSTRUCTURAL EVOLUTION OF INCONEL 625 DURING HOT ROLLING
Materiali in tehnologije / Materials and technology 49 (2015) 5, 801–806 803
Figure 4: Light micrograph of the studied superalloy microstructure
before hot deformation
Slika 4: Svetlobni posnetek mikrostrukture preiskovane superzlitine
pred vro~o deformacijo
Figure 3: Calculated stress during rolling at different strain rates and
the end of rolling temperatures, A first pass  = 7.12 s–1, T = 1131 °C,
B second pass  = 7.96 s–1, T = 1040 °C, C third pass  = 8.90 s–1, T =
983 °C, D fourth pass  = 9.95 s–1, T = 940 °C, E fifth pass  = 11.12
s–1, T = 840 °C
Slika 3: Izra~unane napetosti med valjanjem pri razli~nih hitrostih
deformacije in kon~nih temperaturah valjanja, A prvi prevlek  = 7,12
s–1, T = 1131 °C, B drugi prevlek  = 7,96 s–1, T = 1040 °C, C tretji
prevlek  = 8,90 s–1, T = 983 °C, D ~etrti prevlek  = 9,95 s–1, T = 940
°C, E peti prevlek  = 11,12 s–1, T = 840 °C
Figure 5: Microstructure of specimens after the hot rolling: a) after the first pass, b) after the second pass, c) after the third pass, d) after the
fourth pass, e) after the fifth pass
Slika 5: Mikrostruktura vro~e valjanih vzorcev: a) po prvem prevleku, b) po drugem prevleku, c) po tretjem prevleku, d) po ~etrtem prevleku in
e) po petem prevleku
to the measurement error and DRX, while the gradual
increase of the stress during a single pass is attributed to
the sample cooling.
3.2 Microstructure
Figure 4 shows the initial microstructure of the
Inconel 625 before hot deformation. The specimen was
annealed at 1200 °C for 30 min, followed by water
quenching, used to obtain a fine, homogeneous  phase
with carbides solute in the matrix. It is clear that the
microstructure consists of equiaxed grains with a mean
grain size of 80 μm and a large number of annealing
twins in the austenite grains, typical for alloys with a low
stacking-fault energy6.
The microstructure of the hot-rolled and air-cooled
specimens obtained for a different number of rolling
passes is given in Figure 5.
The start of the recrystallization process is visible
from Figure 5a, showing the microstructure of the Inco-
nel 625 alloy subjected to a single rolling pass. The DRX
occurs through the so-called necklace mechanism, form-
ing at the grain boundaries. After starting at some indivi-
dual boundaries, as seen in Figure 5a, the recrystallized
grains start to appear along most of the grain boundaries
as the number of rolling passes and the deformation rate
increase (Figure 5b). The recrystallization proceeds by
increasing in number and decreasing in size of the new
grains (Figure 5c), forming along all the boundaries
(Figure 5d) and finally representing a significant part of
the microstructure and regularly appearing at the defor-
mation twin boundaries (Figure 5e).
Figure 5 shows that newly formed twins exist in the
original deformed grains. It has been reported that the
formation of twinning plays an important role for the
nucleation process during DRX in materials with a low
stacking-fault energy6. Higher grain-boundary mobility
results in the nucleation of the twins, which act as the
main activate nucleation mechanism of DRX for nickel-
based superalloys deformed at higher temperatures. This
accelerates the bulging and the separation of bulged parts
from the original grains8. The amount of twin boundaries
increases with increased strain. Figure 5e clearly shows
a large amount of recrystallized grains on the twin grain
boundaries.
During the hot deformation of low-stacking-fault-
energy metals, the plastic deformation leads to a serra-
tion of the grain boundaries. For the sufficient deforma-
tion level, the serrations can lead to the creation of new
grains by bulging from the pre-existing grain bounda-
ries9. Miura et al.10 showed that the nucleation of new
grains occurs preferentially at the triple junctions. The
triple junctions acting as nucleation sites can either be
constituted solely of three or more high-angle grain
boundaries, or also include some twins among the inter-
secting boundaries. The initial stages of DRX appear on
the grain boundaries. The new recrystallized grains
nucleate on the deformed grain boundaries and form a
necklace around it7,11–13, as visible in Figure 5c.
F. TEHOVNIK et al.: MICROSTRUCTURAL EVOLUTION OF INCONEL 625 DURING HOT ROLLING
804 Materiali in tehnologije / Materials and technology 49 (2015) 5, 801–806
Figure 6: EBSD grain maps of samples: a) after one pass, b) after
three passes and c) after five passes
Slika 6: EBSD ploskovna porazdelitev zrn: a) po prvem prevleku, b)
po tretjem prevleku, c) po petem prevleku
It has been reported that during deformation the
cooling is slower due to the conversion of deformation
energy to heat by adiabatic heating14. It is well known
that the temperature rise during hot deformation can not
only induce the process of DRX, but also lead to a rapid
reduction in the work-hardening rate15. A higher strain
rate will result in increased dislocation density, which in
turn will dramatically reduce the recrystallization tempe-
rature16,17.
The deformation of the grains and the occurrence of
new recrystallized grains were confirmed by the EBSD
mapping shown in Figure 6. The microstructures pre-
sented correspond to the Inconel 625 superalloy at the
end of hot deformation after: a) one pass, b) three passes
and c) five rolling passes.
The analysis of the EBSD mapping images reveals
large, non-recrystallized grains with twin grain bounda-
ries and small recrystallized grains in the Inconel 625
alloy microstructure subjected to one rolling pass (Fig-
ure 6a). The evolution of the DRX structures clearly
indicates that DRX starts predominantly at the bounda-
ries of the deformed grains and not inside the grains18.
Normal, large-angle grain boundaries are more suscep-
tible to recrystallization than twin boundaries. In the case
of a sample analysed after three rolling passes (Figure
6b) the non-recrystallized grains are slightly deformed,
reflected in misorientation inside the grains and small
recrystallized grains that are smaller compared to the
sample after one rolling pass. This is due to the higher
final deformation and the lower final temperature. The
recrystallized grains do not remain isolated, but soon
form closed networks, resulting in the well-known
necklace structure19. For the most deformed sample
obtained after five rolling passes (Figure 6c), large non-
recrystallized grains are severely deformed, which is
reflected in the misorientation inside the grains. Further-
more, small recrystallized grains are smaller than in the
case of three rolling passes, caused by an even higher
final deformation rate and a very low final temperature.
The analysis of the EBSD data of the different passes
clearly shows that the recrystallized fraction increases
with strain.
4 CONCLUSIONS
This research provides an overview of the structural
changes occurring during hot rolling of the superalloy
Inconel 625. Specimens were hot rolled from a tempera-
ture of 1200 °C with different deformation rates, ob-
tained by changing the number of rolling passes. A light
microscope and an electron microscope, employing the
electron-backscatter-diffraction technique, were em-
ployed to investigate the microstructure evolution and
the nucleation mechanisms of dynamic recrystallization.
The necklace DRX mechanism was found to be the
dominant material-softening mechanism. The recrystal-
lization begins at the deformed grain boundaries, being
consistent with the necklace mechanism. However, at
high deformation rates, new recrystallized grains also
appear at twin boundaries.
The misorientation in the deformed grains increases
with the lower rolling temperature of the deformation
and the increased strain rate. The fraction of recrystal-
lized grains increases with a higher degree of defor-
mation, while the size of the recrystallized grains is
reduced with the higher deformation rate and the lower
temperature.
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F. TEHOVNIK et al.: MICROSTRUCTURAL EVOLUTION OF INCONEL 625 DURING HOT ROLLING
P. LICHÝ et al.: THERMOPHYSICAL PROPERTIES AND MICROSTRUCTURE OF MAGNESIUM ALLOYS ...
807–811
THERMOPHYSICAL PROPERTIES AND MICROSTRUCTURE OF
MAGNESIUM ALLOYS OF THE Mg-Al TYPE
TERMOFIZIKALNE LASTNOSTI IN MIKROSTRUKTURA
MAGNEZIJEVIH ZLITIN TIPA Mg-Al
Petr Lichý1, Jaroslav Beòo1, Ivana Kroupová1, Iveta Vasková2
1V[B-Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of Metallurgy and Foundry Engineering,
17. listopadu 15/2172, 708 33 Ostrava, Czech Republic
2Technical University of Ko{ice, Faculty of Metallurgy, Department of Ferrous and Foundry Metallurgy, Letna 9, Ko{ice, Slovak Republic
petr.lichy@vsb.cz
Prejem rokopisa – received: 2013-10-01; sprejem za objavo – accepted for publication: 2014-11-24
doi:10.17222/mit.2013.196
Generally speaking, magnesium alloys of the Mg-Al type are used as structural materials, but their disadvantage lies in their low
heat resistance. The addition of suitable alloying elements can be positive, as it makes it possible to achieve good
thermo-mechanical properties. For the application of a specific material for thermally stressed cast parts it is necessary to
consider the extent of their linear and volumetric changes at elevated and high temperatures. The aim of this paper is to study
the behaviour of selected magnesium alloys based on Mg-Al during heat stress conditions in order to simulate real conditions
using dilatometric analyses. These properties were evaluated on samples of alloys prepared by gravity casting in metallic
moulds. The effect of the metallurgical processing of the alloys on the studied parameters was also investigated.
Keywords: castings, magnesium alloys, thermophysical properties, microstructure
V splo{nem se zlitine Mg-Al uporabljajo kot konstrukcijski materiali, toda njihova pomanjkljivost je, da imajo majhno toplotno
odpornost. Dodatek ustreznih zlitinskih elementov pozitivno vpliva na izbolj{anje njihove toplotne stabilnosti oz. odpornosti.
Pri uporabi nekega materiala, ki bo toplotno obremenjen, moramo upo{tevati njegove linearne in volumenske spremembe pri
povi{anih in visokih temperaturah. Namen tega prispevka je {tudij vedenja izbranih Mg-zlitin tipa Mg-Al med toplotnimi
obremenitvami, da bi tako simulirali njihovo vedenje v realnih razmerah. Njihovo vedenje smo analizirali z dilatometri~nimi
analizami na vzorcih zlitin, ki so bili izdelani z gravitacijskim litjem v kovinskih modelih. Raziskovali smo tudi vpliv
metalur{kih procesov na izbrane zlitine.
Klju~ne besede: ulitki, magnezijeve zlitine, termofizikalne lastnosti, mikrostruktura
1 INTRODUCTION
Aluminium is the main alloying element in magne-
sium alloys. Mg-Al based alloys belong to the most
widely used group for the foundry industry and they are
the oldest of the foundry magnesium alloys. Aluminium
is one of the few metals that easily dissolves in magne-
sium. These alloys may also contain additional alloying
elements (e.g., Si, Mn, Zr, Th, Ag, Ce). Their properties
are the result of a relatively large area of the solid
solution  in the equilibrium diagram of a Mg-Al alloy
and by the possibility of their alloying also with other
elements. The most common alloys contain 7–10 % of
Al.1
Alloys containing more than 7 % Al are hardenable,
and during their hardening a discontinuous precipitate of
the Mg17Al12 phase is formed, the alloys are usually
alloyed with small quantities of zinc and manganese. An
increasing aluminium content significantly increases the
solidification interval and thus also the width of the
two-phase zone. Such alloys have during gravity casting
a strong tendency to form micro-shrinkages and shrink-
age porosities. For this reason the content of Al in the
alloys for gravity casting does not exceed 5 %. A brittle
intermetallic Mg17Al12 phase is formed above the solu-
bility limit. The limit of solubility of aluminium at the
eutectic temperature is at the amount fraction x = 11.5 %
(mass fraction, w = 12 %), and approximately 1 % at
room temperature. As a result of this, the Mg17Al12 phase
plays a dominant role and it decides what the properties
will be.2,3
2 PROPERTIES OF MAGNESIUM ALLOYS AT
ELEVATED TEMPERATURES
The use of magnesium alloys in the automotive indu-
stry is currently limited to several chosen applications
(such as car dashboard, steering wheel, structure of seats,
etc.).4 The alloys used in these applications are based on
the Mg-Al system, for example the series AM and AZ.
The alloys based on the Mg-Al system are on an
industrial scale the most acceptable from the perspective
of economics. These alloys offer a good combination of
strength and ductility at room temperature. Another
advantage is their good corrosion resistance and
excellent pourability. The main areas of the increasing
use of Mg alloys for car manufacturers are such compo-
nents as gear boxes and engine blocks. These applica-
Materiali in tehnologije / Materials and technology 49 (2015) 5, 807–811 807
UDK 669.721.5:621.74.043:620.186 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(5)807(2015)
tions require exploitation in operating conditions at a
temperature of 150–200 °C. Commercial magnesium
alloys of the type AM and AZ do not have these pro-
perties, this is related to their poor mechanical properties
at elevated temperatures,5 and the cause is the low struc-
tural stability, and thus the creep resistance. Develop-
ments in this area brought about the introduction of new
alloys, i.e., Mg-Al-Sr (AJ) and Mg-Al-RE (AE). Magne-
sium alloys alloyed with rare-earth metals, which do not
contain aluminium, are in the long term recognised as
one of the most resistant to creep. From the viewpoint of
the use of castings made from Mg-alloys at elevated tem-
peratures the volumetric and linear changes of the
material are also critical. Their magnitude is determined,
among other things, by the chemical composition and by
the microstructure of the material (e.g., by presence and
share of individual phases, by grain size).
3 DESCRIPTION OF USED ALLOYS
For the experimental evaluation both commonly used
alloys AZ91, AM60, AMZ40 and the AJ62 alloy were
used. All the compared materials were supplied by a
Czech manufacturer and Table 1 shows their chemical
composition according to the supplier’s certificate.
4 PREPARATION OF THE TEST SAMPLES
Magnesium alloys were melted in an electric resis-
tance furnace in a metal crucible made from low-alloyed
steel. Magnesium alloys are highly reactive thanks to the
high affinity of magnesium for oxygen. For this reason
the material was treated during melting with an agent
having the trade mark EMGESAL. This material in a
form of covering and refining flux served for limiting the
alloy oxidation and for cleaning the melt of possible
inclusions. The castings were produced by gravity cast-
ing into the ingot mould made of cast iron, which was
prior to pouring pre-heated to a temperature of 450 °C ±
30 °C in order to extend its service life and to achieve
sufficient fluidity of the metal. Casting temperatures and
operating temperatures were kept in a narrow range in
order to achieve as high as possible limitation of the
influence of different cooling effects. Part of the melt
was treated metallurgically by an agent with the com-
mercial designation MIKROSAL MG T 200, which was
based on hexachloretane, and which introduced into the
melt the nuclei for crystallisation to achieve the fine-
grained structure. This was followed by the manufacture
of samples from the castings for an evaluation of the
microstructure, as well as samples for a determination of
the thermo-physical properties, and test samples for the
metallographic analyses.
5 MICROSTRUCTURE OF THE CAST SAMPLES
The microstructure of the AZ91 alloy is shown in
Figure 1. It is a structure with a share of eutectics and
with a significant share of the Mg17Al12 phase. The
addition of the forced crystallisation nuclei did not have
any significant effect on this material (Figure 2). A large
share of plate-type precipitates is also evident. Figure 3
P. LICHÝ et al.: THERMOPHYSICAL PROPERTIES AND MICROSTRUCTURE OF MAGNESIUM ALLOYS ...
808 Materiali in tehnologije / Materials and technology 49 (2015) 5, 807–811
Table 1: Chemical composition of the used magnesium alloys
Tabela 1: Kemijska sestava uporabljenih magnezijevih zlitin
alloy
Element, w/%
Zn Al Si Cu Mn Fe Ni Ca Be Sr
AZ91 0.56 8.80 0.06 0.004 0.20 0.004 0.001 0.000 0.001 0.00
AM60 0.07 5.78 0.03 0.001 0.33 0.003 0.001 0.000 0.001 0.00
AMZ40 0.14 3.76 0.02 0.001 0.34 0.003 0.000 0.000 0.001 0.00
AJ62 0.01 5.78 0.04 0.001 0.35 0.003 0.001 0.008 0.001 2.92
Figure 2: Inoculated material AZ91
Slika 2: Cepljena zlitina AZ91
Figure 1: Non-inoculated material AZ91
Slika 1: Necepljena zlitina AZ91
shows the microstructure of the AM60 alloy, which
consists of a matrix containing a solid solution of -Mg,
a secondary phase of Al-Mn compounds and eutectics.
Figure 4 illustrates the structure of the AM60 alloy after
the metallurgical treatment (inoculation). The structure is
much finer, and its homogeneity was generally im-
proved, which means that it is possible to achieve better
mechanical properties.
The effect of inoculation was not manifested in the
achieved structure of the AMZ40 alloys (Figures 5 and
6). This alloy is characterised by the low content of
alloying elements; its structure is formed by the primary
-Mg phase and by eutectics without the presence of
precipitates.
Figures 7 and 8 show the microstructure of the alloy
AJ62, consisting of the basic solid solution , as well as
P. LICHÝ et al.: THERMOPHYSICAL PROPERTIES AND MICROSTRUCTURE OF MAGNESIUM ALLOYS ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 807–811 809
Figure 8: Inoculated material AJ62
Slika 8: Cepljena zlitina AJ62
Figure 5: Non-inoculated material AMZ40
Slika 5: Necepljena zlitina AMZ40
Figure 7: Non-inoculated material AJ62
Slika 7: Necepljena zlitina AJ62
Figure 4: Inoculated material AM60
Slika 4: Cepljena zlitina AM60
Figure 6: Inoculated material AMZ40
Slika 6: Cepljena zlitina AMZ40
Figure 3: Non-inoculated material AM60
Slika 3: Necepljena zlitina AM60
several types of intermetallic phases, i.e., (Al, Mg)4Sr,
Al3Mg13Sr and very small quantity of Mn5Al8. In the case
of the inoculated alloy (Figure 8) it is possible to find
some differences; it is, however, impossible to identify
unequivocally the finer structure.
6 MEASUREMENT OF THE PHYSICAL
PROPERTIES
Dilatometric analyses were performed in order to
verify the behaviour of the materials at elevated tempe-
ratures. Use of the given material for thermally stressed
components is related, among others, to the dimensional
stability of the cast part.
Heat expansion (Equation (1)) of the material is
usually characterised by the mean temperature
coefficient (coefficient of linear expansion):
α T
T T
T0
d
d
=
−
−
= ⎛⎝
⎜ ⎞
⎠
⎟
l l
l T T l
l
T
0
0 0
1
( )
(1)
where:
T – coefficient of linear heat expansion
l0 – sample length under reference (e.g. laboratory) tem-
perature
dl – change of the sample length
dT – difference in temperatures.
The linear heat expansion was measured on the
above-described samples with use of a Netzsch DIL
402C/7 dilatometer. The experiments ran in the tempera-
ture interval from (20 ± 5) °C up to 350 °C with heating
and cooling rates of 15.0 K/min, with a holding time of
30 min at the maximum temperature (isotherm) in a pro-
tective argon atmosphere (99.9999 % Ar) with a constant
gas flow of 20 mL/min. The size of the used samples was
as follows: a mean length of 20 mm and a mean diameter
of 6 mm. Samples of the non-inoculated (marked in the
name of the sample – as “non”) and inoculated materials
(marked as “in”) were divided into two groups. The first
group was used in the as-cast state without previous heat
stressing and the second one after their heat stressing
(250 °C, 30 min) in the argon atmosphere for checking
the influence of the increased temperature during the
stressing of the castings in real conditions.
Table 2 contains the results of the measurements of
the coefficient of linear heat expansion (T) for the
chosen temperature interval ((20 ± 5) °C up to 350 °C).
The table gives the values for the samples without heat
stressing (Tlab) and after heat stressing (250 °C, 30 min) –
T250. Figures 9 and 10 show the greatest change of length
under a temperature of 350 °C (max l350 °C).
Table 2: Overview of thermophysical parameters of the Mg alloys
Tabela 2: Pregled toplotno-fizikalnih parametrov Mg-zlitin
Specimen
Tlab T250
T × 10–6/K–1 T × 10–6/K–1
AZ91 non 29.6319 27.7702
AZ91 in 28.8887 27.3736
AM60 non 28.7385 26.2843
AM60 in 26.2465 26.0252
AMZ40 non 28.2639 18.1402
AMZ40 in 17.3364 17.1782
AJ62 non 28.8727 26.1917
AJ62 in 18.1686 17.5744
In the as-cast state, the highest absolute value of T
(28.8727 × 10–6) was achieved for the AZ91 alloy, and
the lowest (17.3364 × 10–6 was achieved for the AMZ40
alloy after its metallurgical treatment. It is possible to
find from the measured values a correlation between the
measured value of the coefficient of linear thermal
dilatation and the addition of the inoculant. The structure
influenced by the inoculation shows substantially lower
values of T.
A similar effect was observed in the thermally
stressed samples of the magnesium alloys. In these cases
the achieved values of the coefficient of linear thermal
expansion were also the highest in the AZ91 alloy,
followed by the AJ62 alloy, while the lowest values were
achieved in the AMZ40 alloy.
The values of the maximum length change achieved
at the temperature of 350 °C also correspond with these
P. LICHÝ et al.: THERMOPHYSICAL PROPERTIES AND MICROSTRUCTURE OF MAGNESIUM ALLOYS ...
810 Materiali in tehnologije / Materials and technology 49 (2015) 5, 807–811
Figure 10: Thermal expansion of material after heat stress
Slika 10: Toplotni raztezek zlitin po toplotni obremenitvi
Figure 9: Thermal expansion of material without heat stress
Slika 9: Toplotni raztezek zlitin brez toplotne obremenitve
results. The positive impact of the inoculation was
unequivocally demonstrated here, particularly in the case
of the AMZ40 alloy. In case of the samples without ther-
mal stressing (Figure 9) it was possible to note a consi-
derable difference between the non-inoculated and
inoculated materials. On the other hand, in the case of
the thermally stressed samples (Figure 10), this diffe-
rence is much smaller. It is therefore possible to assume
a distinct influence of a possible heat treatment on these
properties.
7 CONCLUSIONS
This study was focused on an evaluation of the ther-
mophysical properties and microstructure of the magne-
sium alloys AZ91, AM60, AMZ40 and AJ62. On the
basis of the realised experiments, it is possible to see a
positive impact of the inoculation on the achievement of
the smaller dilations of materials, which are important
from the perspective of the use of casting for heat-
stressed parts. The lowest values of the coefficient of
linear thermal expansion and the maximum value of
expansion at the temperature of 350 °C were achieved
with samples of the alloys AJ62 and AMZ40, which
were metallurgically processed. Despite the high cooling
effect of the metallic mould, the positive effect of inocu-
lation on the evaluated properties and microstructure of
majority of the magnesium alloys were visible.
Acknowledgements
This work was conducted within the frame of the
research project TA02011333 (Technology Agency of
the CR).
8 REFERENCES
1 M. J. F. Gándara, Mater. Tehnol., 45 (2011) 6, 633–637
2 B. L. Mordike, T. Ebert, Materials Science and Engineering A, 302
(2001) 1, 37–45, doi:10.1016/S0921-5093(00)01351-4
3 S. Roskosz, J. Adamiec, M. Blotnicki, Archives of Foundry Engi-
neering, 7 (2007) 1, 143–146
4 E. Adámková, P. Jelínek, S. [tudentová, Manufacturing Technology,
13 (2013) 3, 255–262
5 P. Lichý, M. Cagala, J. Beòo, Mater. Tehnol., 47 (2013) 4, 503–506
P. LICHÝ et al.: THERMOPHYSICAL PROPERTIES AND MICROSTRUCTURE OF MAGNESIUM ALLOYS ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 807–811 811

M. OSTRÝ et al.: MICRO-ENCAPSULATED PHASE-CHANGE MATERIALS FOR LATENT-HEAT STORAGE: ...
813–816
MICRO-ENCAPSULATED PHASE-CHANGE MATERIALS FOR
LATENT-HEAT STORAGE: THERMAL CHARACTERISTICS
MIKROENKAPSULIRANI MATERIALI S FAZNO PREMENO ZA
SHRANJEVANJE LATENTNE TOPLOTE: TOPLOTNE
ZNA^ILNOSTI
Milan Ostrý1, Darina Dostálová1, Tomá{ Klubal1, Radek Pøikryl2, Pavel Charvát3
1Brno University of Technology, Faculty of Civil Engineering, Institute of Building Structures, Veveøí 95, 602 00 Brno, Czech Republic
2Brno University of Technology, Faculty of Chemistry, Purkyòova 464/188, 612 00 Brno, Czech Republic
3Brno University of Technology, Faculty of Mechanical Engineering, Technická 2, 616 69 Brno, Czech Republic
ostry.m@fce.vutbr.cz
Prejem rokopisa – received: 2013-10-01; sprejem za objavo – accepted for publication: 2014-11-24
doi:10.17222/mit.2013.210
A significant problem of the utilization of renewable energy is the mismatch between the demand and availability. From this
point of view, heat storage contributes to the use of renewable energy for the heating and cooling of buildings. Heat can be
stored using sensible- or latent-heat-storage techniques. Nowadays, activated concrete structures are used as an alternative
cooling-and-heating system. Light-weight timber structures with an equivalent heat-storage capacity can be used instead of
these heavy-weight structures. The latent-heat storage using micro-encapsulated phase-change materials (PCMs) represents a
way to provide an adequate thermal mass within a small thickness of thermally activated structures. A series of experiments in a
lab and test rooms were carried out to investigate the thermal behavior of micro-encapsulated PCMs for building applications.
The Micronal®DS 5040 X microencapsulated PCM was used as the latent-heat-storage medium in combination with gypsum
plaster. Differential scanning calorimetry (DSC) was used for a thermal analysis of the latent-heat storage medium. The storage
medium underwent thermal cycling for an assessment of the heat-storage capacity changes during the proposed service life. The
floor structure and the plasterboards fixed onto the ceiling and walls of an experimental room with an integrated PCM were
thermally activated with capillary tubes. The performance of the system was continuously observed and evaluated.
Keywords: phase-change materials (PCMs), latent-heat storage (LHS), differential scanning calorimetry (DSC), thermal
cycling, building, thermal capacity
Pomemben problem pri izkori{~anju obnovljive energije je neskladje med potrebo in razpolo`ljivostjo. S tega vidika
shranjevanje toplote prispeva k uporabi obnovljive energije pri ogrevanju in hlajenju zgradb. Toplota se lahko smiselno shrani s
tehniko shranjevanja latentne toplote. Danes se uporabljajo aktivirane betonske strukture kot alternativni ogrevni ali ohlajevalni
sistemi. Lahke lesene strukture z enakovredno kapaciteto shranjevanja toplote se lahko uporabijo namesto teh te`kih struktur.
Shranjevanje latentne toplote z uporabo mikroenkapsuliranega materiala s fazno premeno (PCM) je na~in za zagotovitev
ustrezne koli~ine toplote kljub majhni debelini toplotno aktiviranih zgradb. Izvr{ena je bila serija preizkusov v laboratoriju in v
preizkusnih sobah za ugotovitev toplotnega vedenja mikroenkapsuliranega PCM za uporabo pri zgradbah. Kot medij za
shranjevanje latentne toplote je bil uporabljen mikroenkapsulirani PCM Micronal®DS 5040 X v kombinaciji z ometom iz
mavca. Diferen~na vrsti~na kalorimetrija (DSC) je bila uporabljena za toplotno analizo medija za shranjevanje latentne toplote.
Shranjevalni medij je bil izpostavljen toplotnim ciklom za ugotovitev sprememb kapacitete shranjene toplote med predlaganim
delovanjem. Tla in mav~ne plo{~e z vgrajenim PCM, pritrjene na strop in stene eksperimentalne sobe, so bile toplotno
aktivirane s kapilarnimi cevkami. Vedenje sistema je bilo opazovano in ocenjeno.
Klju~ne besede: materiali s fazno premeno (PCM), latentno shranjevanje toplote (LHS), diferen~na vrsti~na kalorimetrija
(DSC), toplotno cikliranje, zgradba, toplotna kapaciteta
1 INTRODUCTION
Heat storage can be in the forms of sensible heat in a
liquid or solid medium, heat of fusion, chemical energy
or products in a reversible chemical reaction.1 Latent-
heat storage represents a more efficient form of heat
storage compared to sensible-heat storage due to a high
storage capacity per volume or mass. Common storage
media, e.g., brick walls or concrete structures, can store
heat only through an increase in their temperature. This
is a problem when the required change in ambient tem-
perature is narrow. A typical example is indoor climate
in buildings. The occupants require a stable indoor cli-
mate with very low changes in the temperature during
the year. The heat-storage capacity of a building enve-
lope helps to maintain the indoor air temperature in the
required range without additional cooling and heating
systems. This principle is strongly dependent on the
potential of heat loses or heat gains. Building envelopes
of modern office buildings are made of light-weight
materials. Regarding the thermal comfort, the main dis-
advantages of light-weight structures are their low ther-
mal mass, low thermal inertia and a potential comfort
problem such as overheating.2
Therefore, light-weight building materials with a
high heat-storage capacity represent a way of decreasing
the energy consumption and operating cost.
Phase-change materials (PCMs) are latent-heat-sto-
rage media. The PCMs are able to store large amounts of
heat during the melting process. The solidifying process
Materiali in tehnologije / Materials and technology 49 (2015) 5, 813–816 813
UDK 536.2:536.65:66.017 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(5)813(2015)
is accompanied by a release of the stored heat. The
PCMs should have the melting and solidification
temperature range in the practical range of an application
and they need to have a high latent heat of fusion and a
high thermal conductivity for the effective thermal sto-
rage2.
There are some possible techniques to integrate
PCMs directly into the building materials:3
• a simple immersion of a porous material in a liquid
PCM;
• a vacuum impregnation of a PCM in the porous
aggregates of a building material;
• a direct mixing of an encapsulated PCM into concrete
or plaster during a concrete or plaster mixing.
Microcapsules are small containers encapsulating a
PCM that can be mixed with many building materials.
Micro-encapsulation is a process in which PCM particles
are surrounded or coated with a continuous film of a
polymeric material to produce capsules in a micrometer
to millimeter range.4 The PCM forms the core and the
polymeric material creates the shell. The microcapsules
are usually of a spherical shape though other shapes are
also possible. The microcapsules allow for the volume
changes of the PCM, reducing the reaction of the PCM
with the outside environment.5 This technology of encap-
sulation increases the heat-transfer area and allows for an
easier handling of PCMs. There is no risk of a massive
leakage of the PCMs since the PCM volume in a capsule
is very small. The building structures with incorporated
PCMs are similar to the commonly used building mate-
rials. The micro-encapsulation can be carried out with
physical or chemical methods.4 The physical methods
are: pan coating, air-suspension coating, centrifugal
extrusion, vibrational nozzle and spray drying. There are
several chemical methods: interfacial polymerization,
in-situ polymerization and matrix polymerization.
2 MATERIAL AND METHODS
There are several important aspects of a practical use
of the latent-heat-storage technology:
• the kind of the phase-change material;
• the kind of encapsulation;
• the thermal characteristics and thermal stability of
PCMs;
• the integration in building structures;
• the heat transfer.
Moreover, there is a requirement for a long-term sta-
bility of the storage medium and the PCM-container
system. Other relevant aspects are the useful life of these
systems and the number of cycles they can withstand
without a degradation of their properties.6
The research and development at the Brno University
of Technology dealing with thermal storage is focused
on the integration of PCMs into building envelopes to
provide a thermal comfort in the range required by the
legislation. A micro-encapsulated paraffin-based PCM is
used for this purpose. A team of investigators from the
Brno University of Technology developed and patented
heat-storage modules for a stabilization of the thermal
environment in buildings. The composition of a module
is shown in Figure 1.
The system can effectively stabilize the thermal mi-
croclimate in buildings with low operating costs. It
allows not only a passive cooling of the rooms in
summer, but it can also be used for a low-temperature
heating in winter. The main technological components of
the module are the basic slab, the heat-storage plaster
and the capillary tubes for a thermal activation of the
module. The system of the capillary tubes is placed in a
layer of gypsum plaster with integrated micro-pellets
containing a PCM for an advanced latent-heat/cold sto-
rage. Micronal®DS 5040 X was used as the latent-heat-
storage medium (PCM). The capillary tubes integrated in
the heat-storage plaster were connected to the distribu-
tion and return pipes. It is necessary to provide a direct
exposure of the heat-storage layer to the indoor environ-
ment for a good heat transfer; therefore, the heat-storage
structures containing PCMs cannot be obstructed by
furniture or suspended ceilings.
Gypsum plaster with a special composition for the
proposed use was modified with a mass fraction w =
30 % of Micronal®DS 5040 X.
Two important properties were tested in the experi-
ments:
• the change in the thermal characteristics of the PCM
after 100 melting-freezing cycles;
• the change in the thermal conductivity of the plaster
with the microencapsulated PCM.
The thermal-energy-storage properties, the thermal
reliability and the thermal durability of the composite
with the PCM were determined with differential scann-
ing calorimetry (DSC).7
DSC is a widely used method from the group of ther-
mal-analysis methods. Nowadays, DSC is used for ther-
mal analyses of new materials, e.g., in the study of the
phase changes and testing quality of polymer materials.
DSC compensation is a method whose principle consists
of measuring the temperature and the heat flux during
the heating or cooling of a tested material. The process is
realized at a constant rate of heating or cooling and
carried out in a closed chamber. The test result can be
used to determine the melting temperature of a sample –
the onset/offset of the phase-transition temperature – the
M. OSTRÝ et al.: MICRO-ENCAPSULATED PHASE-CHANGE MATERIALS FOR LATENT-HEAT STORAGE: ...
814 Materiali in tehnologije / Materials and technology 49 (2015) 5, 813–816
Figure 1: Detail of a heat-storage structure
Slika 1: Detajl strukture za shranjevanje toplote
crystallization, or various types of crystallization, the
heat capacity and the temperature degradation. The diffe-
rential scanning calorimeter is a twin instrument, com-
prising a sample and a reference calorimeter within a
common thermal enclosure, where the two calorimeters
are usually assumed to be identical. A PYRIS1 Perkin
Elmer calorimeter was used in the tests.
Micronal®DS 5040 X was tested before and after the
thermal cycling. A sample of Micronal®DS 5040 X was
placed in a glass laboratory test tube with a thermo-
couple sensor. The tube was immersed in liquid nitrogen
until the temperature measured by the thermocouple
dropped to –10 °C, then the tube was immersed in hot
water (60 °C). The tube remained in nitrogen and warm
water for about 1 min and 1.5 min. The whole procedure
was repeated 100-times.
The heat transfer between the indoor environment
and the heat storage depends mainly on the thermal con-
ductivity of the plaster with the micro-encapsulated
PCM. The thermal conductivity was determined with the
guarded-hot-plate method. In this method a solid sample
of a material is placed between two plates. One plate is
heated and the other is cooled. The temperatures of the
plates are monitored until they are constant. A constant
heat flow flows through the test sample in the station-
ary-temperature state. The steady-state temperatures, the
thickness of the sample and the heat input to the hot
plate are used to calculate the thermal conductivity.
Several samples of gypsum plaster with various com-
positions were tested:
• referential sample (sample no. 0) – plain gypsum pla-
ster;
• modified plaster (sample no. 1) – w = 68 % gypsum
plaster and w = 32 % PCM;
• modified plaster (sample no. 2) – w = 65 % gypsum
plaster, w = 30 % PCM and w = 5 % natural graphite;
• modified plaster (sample no. 3) – w = 61.5 % gypsum
plaster, w = 29 % PCM and w = 9.5 % natural gra-
phite;
• modified plaster (sample no. 4) – w = 67.9 % gypsum
plaster, w = 31.8 % PCM and w = 0.3 % expanded
graphite;
• modified plaster (sample no. 5) – w = 67.7 % gypsum
plaster, w = 31.7 % PCM and w = 0.6 % expanded
graphite.
3 RESULTS AND DISCUSSION
3.1 Thermal stability of the latent-heat-storage me-
dium
The characteristics of all the samples were tested
with DSC twice, before and after the cycling. The expe-
riments were carried out at a temperature rate of 1.0
K/min. The onset and peak temperatures during the
heating and cooling depend on the temperature rate. The
melting temperatures rise with the temperature rate. The
risk of supercooling rises with the heating rate.
The results are shown in Tables 1 and 2. Both tables
show a minimum difference in the enthalpy of the phase
change of the tested samples. The difference between the
onset temperatures is 0.1 K.
The DSC curves in Figure 2 show minimum diffe-
rences between the characteristics of the samples.
Table 1: Characteristics of Micronal®DS 5040 X before thermal
cycling
Tabela 1: Zna~ilnosti Micronal®DS 5040 X pred toplotnim ciklira-
njem
Heating Cooling
Onset temperature, °C 18.8 21.8
Peak temperature, °C 24.5 Multi-peak
Enthalpy, J/g 96.2 92.6
Table 2: Characteristics of Micronal®DS 5040 X after thermal
cycling
Tabela 2: Zna~ilnosti Micronal®DS 5040 X po toplotnem cikliranju
Heating Cooling
Onset temperature, °C 18.9 21.9
Peak temperature, °C 24.5 Multi-peak
Enthalpy, J/g 96.5 93.4
3.2 Thermal conductivity of the plaster with PCM
Thermal conductivity is negatively influenced by the
micro-encapsulated PCM. The increase in the thermal
conductivity can be achieved with an addition of natural
or expanded graphite. The influence of graphite on the
thermal conductivity is shown in Table 3. As can be
seen, the best material for the heat-conductivity improve-
ment is natural graphite. A low amount of expanded
graphite (w = 0.3 %) has no positive effect on the heat
conductivity.
M. OSTRÝ et al.: MICRO-ENCAPSULATED PHASE-CHANGE MATERIALS FOR LATENT-HEAT STORAGE: ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 813–816 815
Figure 2: DSC results for Micronal®DS 5040 X before and after ther-
mal cycling
Slika 2: Rezultati DSC za Micronal®DS 5040 X pred toplotnim
cikliranjem in po njem
Table 3: Thermal conductivity of plaster with different compositions
Tabela 3: Toplotna prevodnost ometa pri razli~nih sestavah
Number of sample
Pour density
of plaster
kg/m3
Thermal
conductivity
W/(m K)
Referential sample (no. 0) 1080 0.3866
Plaster with PCM (no. 1) 677 0.1760
w = 5 % natural graphite (no. 2) 676 0.1996
w = 9.5 % natural graphite (no. 3) 665 0.2336
w = 0.3 % expan. graphite (no. 4) 647 0.1598
w = 0.6 % expan. graphite (no. 5) 589 0.1766
4 CONCLUSION
The PCM used in the experiment did not significantly
change its thermal properties even after 100 thermal
cycles. From this point of view, the tested organic PCM
is suitable for integration in building structures. The
advantage of this kind of PCM is the possibility of
integration in gypsum plaster that can be combined with
capillary tubes for charging or discharging the heat. The
modified plaster acting as a thermally activated layer can
be used in radiant heating or cooling in residential
buildings. However, the addition of a PCM causes a
significant decrease in the thermal conductivity of the
plaster. This problem can be overcome by adding natural
graphite, e.g., w = 9.5 % of natural graphite increased the
thermal conductivity of the plaster with the PCM from
0.1760 W/(m K) up to 0.2336 W/(m K).
Acknowledgement
This work paper has been worked out under project
No. P104/12/1838 "Utilization of latent heat storage in
phase change materials to reduce primary energy con-
sumption in buildings", supported by Czech Science
Foundation and under the project No. LO1408 "AdMaS
UP – Advanced Materials, Structures and Technologies",
supported by Ministry of Education, Youth and Sports
under the "National Sustainability Programme I".
5 REFERENCES
1 F. Agyenim, N. Hewitt, F. Eames, M. Smyth, A review of materials,
heat transfer and phase change problem formulation for latent heat
thermal energy storage systems (LHTESS), Renewable and Sustain-
able Energy Reviews, 14 (2010) 2, 615–628, doi:10.1016/j.rser.
2009.10.015
2 N. Soares, J. J. Costa, A. R. Gaspar, P. Santos, Review of passive
PCM latent heat thermal energy storage systems towards buildings'
energy efficiency, Energy and Buildings, 59 (2013), 82–103,
doi:10.1016/j.enbuild.2012.12.042
3 T. C. Ling, C. S. Poon, Use of phase change materials for thermal
energy storage in concrete: An overview, Construction and Building
Materials, 46 (2013), 55–62, doi:10.1016/j.conbuildmat.2013.04.031
4 V. V. Tyagi, S. C. Kaushik, S. K. Tyagi, T. Akiyama, Development of
phase change materials based microencapsulated technology for
buildings: A review, Renewable and Sustainable Energy Reviews, 15
(2011) 2, 1373–1391, doi:10.1016/j.rser.2010.10.006
5 L. Bayés-García, L. Ventola, R. Cordobilla, R. Benages, T. Calvet,
M. A. Cuevas-Diarte, Phase Change Materials (PCM) microcapsules
with different shell compositions: Preparation, characterization and
thermal stability, Solar Energy Materials and Solar Cells, 94 (2010)
7, 1235–1240, doi:10.1016/j.solmat.2010.03.014
6 B. Zalba, J. M. Marín, L. F. Cabeza, H. Mehling, Review on thermal
energy storage with phase change: materials, heat transfer analysis
and applications, Applied Thermal Engineering, 23 (2003) 3,
251–283, doi:10.1016/S1359-4311(02)00192-8
7 A. Sari, A. Biçer, Preparation and thermal energy storage properties
of building material-based composites as novel form-stable PCMs,
Energy and Buildings, 51 (2012), 73–83, doi:10.1016/j.enbuild.
2012.04.010
M. OSTRÝ et al.: MICRO-ENCAPSULATED PHASE-CHANGE MATERIALS FOR LATENT-HEAT STORAGE: ...
816 Materiali in tehnologije / Materials and technology 49 (2015) 5, 813–816
J. ZACH et al.: CERAMIC MASONRY UNITS INTENDED FOR THE MASONRY RESISTANT ...
817–820
CERAMIC MASONRY UNITS INTENDED FOR THE MASONRY
RESISTANT TO HIGH HUMIDITY
KERAMI^NI GRADBENI ELEMENTI, NAMENJENI ZA ZGRADBE,
ODPORNE PROTI VISOKI VLAGI
Jiøí Zach, Vítìzslav Novák, Jitka Hroudová, Martin Sedlmajer
Brno University of Technology, Faculty of Civil Engineering, Veveøí 331/95, 602 00 Brno, Czech Republic
zach.j@fce.vutbr.cz, novak.v@fce.vutbr.cz, hroudova.j@fce.vutbr.cz, sedlmajer.m@fce.vutbr.cz
Prejem rokopisa – received: 2014-08-01; sprejem za objavo – accepted for publication: 2014-10-14
doi:10.17222/mit.2014.170
The Faculty of Civil Engineering has been developing modern masonry blocks for several years. The aim is to develop masonry
units that provide a good thermal insulation, mechanical and acoustic properties and a reduction in the energy consumption
required for their production. Given the ever increasing frequency of natural disasters, which also affect landlocked countries, a
part of this research also focuses on the development of ceramic blocks resistant to high humidity. High humidity is one of the
most harmful effects occurring in a building construction. It causes an overall deterioration of the building construction and can
lead to its degradation. Ceramic-brick constructions are already under stress due to a high humidity during the building period
and also afterwards during their use. However, natural disasters such as floods, which have often occurred in the Czech Republic
in recent years, cause the destruction of building constructions and their subsequent demolition. The paper deals with the
possibility of a preventive use of hydrophobic ceramic masonry units intended for masonry plinths. This hydrophobisation
serves as the prevention in the cases, in which the conventional waterproofing fails and also in the case of an extremely high
humidity.
Keywords: ceramic masonry blocks, high humidity, hydrophobic, waterproofing, capillary absorption
Fakulteta za gradbeni{tvo `e ve~ let razvija moderne gradbene elemente. Namen je razviti moderno gradbeno enoto, ki bo imela
dobro toplotno izolacijo, mehanske in akusti~ne lastnosti ter zmanj{ano porabo energije za njeno proizvodnjo. Zaradi vedno
ve~je pogostosti naravnih katastrof, ki vplivajo tudi na oto{ke dr`ave, je del razvoja usmerjen tudi v razvoj kerami~nih opek,
odpornih proti visoki vlagi. Ta je med najbolj {kodljivimi vplivi na gradbeno konstrukcijo. Povzro~a splo{no poslab{anje
zgradbe in lahko povzro~i celo njen razpad. Konstrukcije iz kerami~nih opek so izpostavljene veliki vlagi `e pri gradnji, kot tudi
kasneje pri uporabi. Pri naravnih nesre~ah, kot so poplave, ki so bile zadnja leta pogoste v ^e{ki Republiki, so bile vzrok za
uni~enje gradbenih konstrukcij in posledi~no ru{enje le-teh. ^lanek obravnava mo`nost preventivne uporabe hidrofobne
kerami~ne opeke za podlago zgradbe. Ta hidrofobizacija se uporablja kot za{~ita v primeru, kjer navadna hidroizolacija odpove,
in tudi v primerih ekstremno visoke vla`nosti.
Klju~ne besede: kerami~ni gradbeni elementi, visoka vla`nost, hidrofoben, hidroizolacija, kapilarna absorpcija
1 INTRODUCTION
Two waterproofing principles can be applied to pro-
tect a building construction from the negative effects of
water and moisture:1,2
The indirect waterproofing principle – following this
principle, we minimise/reduce the moisture stress on a
construction, e.g., by placing the building in an optimum
environment, using the drainage of the adjacent building,
increasing the surface temperature of the construction,
etc.
The direct waterproofing principle – following this
principle, we prevent water and moisture from penetra-
ting into the structure. This is achieved by means of
waterproof (or moisture-proof) materials, e.g., conti-
nuous-sheet waterproofing, waterproof concrete, injec-
tions, surface penetration and impregnation, surface
hydrophobisation,3–5 etc. It is also possible to prevent
water or moisture penetration using active methods for
reducing the moisture stress on a construction, e.g.,
electokinetic methods.
Masonry constructions, especially their lower parts,
are under stress from increased moisture. Increased
moisture has an effect both during the construction and
during the use of a building. Building constructions are
protected from increased moisture by horizontal, conti-
nuous, waterproof sheets and by the surface finish on the
outside.6
In the case of increased moisture during the con-
struction, the moisture mainly penetrates into the struc-
ture prior to the roof construction and surface finishing.
A moisture increase during the construction does not
necessarily mean that it has a decisive influence on the
building. However, the building must not be sealed or
used immediately after the finishing if it is completed in
autumn or winter. If we fail to do so, a deterioration of
the thermal properties often occurs due to the moisture in
the lower parts of the building, often leading to
hygrothermal failures such as damp areas and moulds.
The stress caused by increased moisture during the
use of a building can have a number of causes. It may be
a failure of the horizontal/vertical waterproofing as a
Materiali in tehnologije / Materials and technology 49 (2015) 5, 817–820 817
UDK 69:69.07:691.4:699.82 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(5)817(2015)
result of damage or a faulty application. Furthermore,
waterproofing can be damaged during additional build-
ing operations, such as vents or wall repairs, etc. Water-
proofing is also ineffective if the masonry reaches over
the foundation plate, into the exterior. In the case of in-
creased moisture during the use of the building, the
problem is considerably more serious and the degree of
masonry degradation is determined by the moisture
stress.
To prevent hygric defects, hydrophobised masonry
units can be used in the lower part of a masonry con-
struction. This paper deals with the possibilities of addi-
tional hydrophobisation of ceramic bricks using chemi-
cal hydrophobic agents available in the EU market.
2 SPECIMENS
The specimens for testing the hydrophobisation of
ceramic bricks were intended for the plinth masonry.
They were wall bricks with a thickness of 250 mm. The
bricks were chosen because of their low mass, small
dimensions and the water absorption equal to that of the
bricks used in thicker walls.
3 HYDROPHOBIC AGENTS
A total of five specimens from different producers
and of different chemical compositions were chosen for
the laboratory experiments. All the tested hydrophobic
agents were silicon-based. They were organosilicate
compounds exhibiting a high thermal resistance between
–50 °C and 200 °C, the resistance to sudden temperature
change, weather, sunlight and a high chemical resistance.
A hydrophobic film was formed thanks to the intermole-
cular condensation of low-molecular-weight methylsilo-
xanes through the effect of atmospheric carbon dioxide.
The main condition of the hydrophobising effect was the
orientation of silicone molecules towards the surface of
the masonry. There was a high amount of short-chain
hydrocarbon residue, e.g., methyl, ethyl group (CH3–,
C2H5–) (Figure 1). This weak, often macromolecular,
layer increases the contact angle to such an extent that
the capillary elevation is reduced to zero or even that a
capillary depression occurs. Thanks to this principle, the
water-vapour and gas permeability of the building
material is retained. These hydrophobic agents cannot be
used in the cases where water under pressure acts upon
the structure.7
Some the most frequently used waterproofing prepa-
rations are silicone resin solutions formed in organic
solvents, most often in white spirit or benzine. The resin
content usually ranges from 3 % to 8 %. These are
single-component solutions, mostly colourless or with a
slight yellow hue and transparent; the hydrophobic effect
occurs immediately after the solvent evaporates. These
preparations cannot be used on wet surfaces; they can be
used for hydrophobising previously hydrophobised mate-
rials. The presence of the solvent is disadvantageous in
terms of fire, ecological and possible health hazard.
The other silicone hydrophobic agents are low-mole-
cular-weight compounds – oligomers, which, due to air
humidity, spontaneously polymerise upon application,
producing the desired polymer. They are usually single-
compound, translucent and colourless. Their advantage
over the high-molecular-weight resins lies in their ability
to chemically bond with the surfaces of quartz grains (if
present) of the material being treated, thus increasing the
abrasion resistance. They are sometimes supplied as a
concentrate and the preparation is performed by mixing
it with an appropriate solvent. Another advantage is their
better ability to penetrate into a porous material.
Silicone microemulsions represent a special group.
They are again silicone compounds with a relatively low
molecular weight which, thanks to their molecular struc-
ture, are able to produce an emulsion with very small
particles after being mixed with water. A slow reaction
with water then produces the desired hydrophobic poly-
mer. Their advantage lies in the possibility of mixing
with water which facilitates the avoidance of organic sol-
vents and a good penetration capability.
Methylsilanolates (methylsiliconates) – sodium or
potassium – also belong to silicone hydrophobic agents.
These substances are hydrophobising varieties of soluble
glass. Similarly to soluble glass, when in contact with
carbon dioxide, they are transformed into a modified gel
of silicic acid that is hydrophobic in this case. Unfortu-
nately, this chemical reaction also produces a hydroxide
of the respective metal (sodium or potassium), later
transformed into a carbonate. Thus, undesirable water-
soluble efflorescent salts enter the structure of the
material being treated. Their presence represents the risk
of the formation of white efflorescence and even the
degradation of the treated material due to the crystallisa-
tion pressure. The advantages of siliconates are their
good water solubility (i.e., the possibility of application
onto wet surfaces) and substantially lower prices in com-
J. ZACH et al.: CERAMIC MASONRY UNITS INTENDED FOR THE MASONRY RESISTANT ...
818 Materiali in tehnologije / Materials and technology 49 (2015) 5, 817–820
Figure 1: Orientation of silicone molecules7
Slika 1: Orientacija silikonskih molekul7
Table 1: Used hydrophobic agents
Tabela 1: Uporabljena hidrofobna sredstva
No. Description
1 Silane-siloxane microemulsion cream
2 Potassium methylsilanolate preparation
3 Hydrophobised siliconate solution
4 Solvent-free siloxane emulsion
5 Siloxane base with aromatic-free organic solvents
parison with the other silicone preparations. The result-
ing gel is non-soluble and thus effectively unremovable.
Silanolates are recommended for the protection of histo-
rical buildings thanks to their good solubility and low
price8,9.
Table 1 contains an overview of the tested hydropho-
bic agents.
The soaking in hydrophobising solutions was always
performed for (10, 20 and 30) s; only non-water-soluble
preparations were applied by spraying them in a given
amount. Next, the individual specimens were stored after
the application of the hydrophobic agent in the labora-
tory environment for a week in order for the excess
moisture to evaporate. The following tests were per-
formed with the specimens thus prepared.
The bricks were put in a plastic box onto a wooden
grid. Subsequently, they were flooded with water at a
temperature of (20 ± 2) °C, always up to 50 mm of the
height of the bricks. The bricks were weighed at regular
intervals: after (1, 2, 3, 6, 24, 48, 72 and 96) h.
4 RESULTS
Hydrophobic preparations Nos. 1–4 were applied by
soaking for (10, 20 and 30) s. Hydrophobic agent No. 5
was applied by being sprayed in the amount of approxi-
mately 250 g per brick. The measurement was also
performed on the non-hydrophobised bricks. The testing
was always carried out on two specimen bricks with a
given hydrophobic-agent concentration and an immer-
sion time. Hydrophobic agents Nos. 1, 2 and 4 were
always prepared in two concentrations. The overview of
the tested concentrations is summarized below:
• solution No. 1 was tested at concentrations of 1 : 20
and 1 : 30,
• solution No. 2 was tested at concentrations of 1 : 10
and 1 : 25,
• solution No. 3 was tested at a concentration of 1 : 2,
• solution No. 4 was tested at concentrations of 1 : 15
and 1 : 30.
The quantity being measured was the brick mass at
the end of each interval (1, 2, 3, 6, 24, 48, 72 and 96) h.
The amount of the absorbed water was calculated from
the ratio of the measured mass to the brick surface area.
Its capillary absorption was determined with the
following formula:
A
m
Aw
w= (1)
Aw – capillary absorption in time (kg m–2)
mw – mass of the brick soaked with water (kg)
A – area of the brick (m2)
The values of the capillary absorption can be com-
pared with each other despite the different methods of
the application of the hydrophobic agents which is the
validity of their efficiency. The results of the individual
measurements are in Table 2:
Table 2: Overview of the obtained absorption values of the specimens
Tabela 2: Pregled dobljenih vrednosti absorpcije vzorcev
Hydro-
phobic
agent
number
Con-
centra-
tion
Dipping
time (s)
Absorption
(%)
Capillary
absorption
(kg m–2)
24 h 96 h 24 h 96 h
1
1 : 30
10 0.42 1.60 0.79 3.04
20 0.41 1.39 0.77 2.64
30 0.24 0.93 0.45 1.77
1 : 20
10 0.35 0.76 0.65 1.44
20 0.36 0.72 0.63 1.36
30 0.38 0.67 0.60 1.27
2
1 : 25
10 16.38 22.05 32.98 42.31
20 15.52 21.51 29.78 40.56
30 11.08 20.14 20.65 40.07
1 : 10
10 10.27 19.77 20.97 40.03
20 8.47 19.61 17.26 38.19
30 6.54 17.65 12.63 35.98
3 1 : 2
10 1.36 2.91 2.50 5.33
20 1.26 2.15 2.44 4.16
30 1.21 2.11 2.13 3.89
4
1 : 30
10 8.58 16.21 15.69 29.51
20 7.51 15.42 11.74 28.19
30 6.25 15.11 10.85 26.52
1 : 15
10 6.01 12.65 12.35 24.64
20 6.34 12.35 11.32 23.25
30 5.52 11.01 6.28 19.97
5 – – 9.42 17.26 18.39 33.70
Refe-
rence – – 25.70 27.22 47.65 50.47
The table shows that the effectiveness of individual
hydrophobic agents varies significantly. The effective-
ness of the hydrophobic agents is influenced by the
concentration and the soaking time. The following chart
compares the effectiveness values for the hydrophobic
agents after 96 h of soaking in water by means of
capillary absorption (Figure 2).
The values show that the most effective hydrophobi-
sation was achieved by applying hydrophobic agent No.
1, specifically at a concentration of 1 : 20. The following
chart on Figure 3 shows the progression of the capillary
J. ZACH et al.: CERAMIC MASONRY UNITS INTENDED FOR THE MASONRY RESISTANT ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 817–820 819
Figure 2: Capillary absorption of individual hydrophobic agents
Slika 2: Kapilarna absorpcija pri uporabi posameznih hidrofobnih
snovi
absorption over time of this hydrophobic agent applied
by soaking for (10, 20 and 30) s, in comparison with the
reference capillary absorption.
5 CONCLUSION
Table 2 lists all the experimentally obtained values,
which show, when compared to the reference values, that
the capillary absorption was reduced in all the instances
by applying the tested hydrophobic agents. The effec-
tiveness of the hydrophobic agents is influenced by the
type of the preparation, its concentration and the time of
the treatment with the agent.
The best results were reached using hydrophobic
agent No. 1, i.e., the agent based on the silane-siloxane
microemulsion cream with a concentration a 1 : 20. With
this hydrophobic agent, as with as all the others, an
increasing soaking time caused an increase in the
effectiveness of the preparation at a reduced capillary
absorption.
It is necessary to note that during a long-term expo-
sure to moisture, capillarity is always at work. These
tested hydrophobic agents cannot be considered as com-
plete substitutes for sheet waterproofing; however, they
can be used for masonry protection during construction.
Also, in the case of an occasional moisture stress on the
lower part of the masonry (spraying precipitation moi-
sture), hydrophobic agents provide a sufficient protec-
tion.
Acknowledgements
This paper was elaborated with the financial support
of the Technology Agency of the Czech Republic within
the project under ref. n. TA04020920 and the project no.
L01408 "ADMAS Up – Advanced Materials, Structures
and Technologies".
6 REFERENCES
1 Czech standard CSN P 73 0600, Waterproofing of buildings – Basic
provisions, 2000
2 Czech standard CSN P 73 0610, Waterproofing of buildings – The
rehabilitation of damp masonry and additional protection of build-
ings against ground moisture and against atmospheric water – Basic
provision, 2000
3 D. Beben, Z. Manko, Influence of selected hydrophobic agents on
some properties of autoclaving cellular concrete (ACC), Construc-
tion and Building Materials, 25 (2011), 282–287, doi:10.1016/
j.conbuildmat.2010.06.028
4 M. Medeiros, P. Helene, Efficacy of surface hydrophobic agents in
reducing water and chloride ion penetration in concrete, Mater.
Struct., 41 (2008), 59–71, doi:10.1617/S11527-006-9218-5
5 L. Wolff, M. Raupach, Hydrophobieren von Stahlbetonoberflächen,
Betonwerk und Fertigteil-Technik, 69 (2003) 10, 12–21
6 Czech standard CSN P 73 0606, Waterproofing of buildings – Con-
tinuous sheet water proofing – Basic provisions, 2000
7 R. Drochytka, Plastic substances, Study support, VUT Brno, Brno,
2007
8 J. de Vries, R. B. Polder, Hydrophobic Treatment of Concrete, Con-
struction and Building Materials, 11 (1997) 4, 259–265, doi:10.1016/
S0950-0618(97)00046-9
9 P. Kotlík, Stavební materiály historických objektù, 1. vydání, Praha
1999
J. ZACH et al.: CERAMIC MASONRY UNITS INTENDED FOR THE MASONRY RESISTANT ...
820 Materiali in tehnologije / Materials and technology 49 (2015) 5, 817–820
Figure 3: Process of capillary absorption over time for hydrophobic
agent No. 1 (1 : 20)
Slika 3: Potek kapilarne absorpcije v odvisnosti od ~asa pri hidrofobni
snovi {t. 1 (1 : 20)
V. RUSNÁK et al.: FLAME RESISTANCE AND MECHANICAL PROPERTIES OF COMPOSITES ...
821–824
FLAME RESISTANCE AND MECHANICAL PROPERTIES OF
COMPOSITES BASED ON NEW ADVANCED RESIN SYSTEM
FR4/12
NEGORLJIVOST IN MEHANSKE LASTNOSTI KOMPOZITOV NA
OSNOVI NOVIH NAPREDNIH SISTEMOV SMOL FR4/12
Vladimír Rusnák1, Soòa Rusnáková2, Ladislav Fojtl2, Milan @aludek2,
Alexander ^apka2
1Faculty of Metallurgy and Materials Engineering, V[B-Technical University of Ostrava, 17. listopadu 15, 708 33 Ostrava-Poruba, Czech
Republic
2Faculty of Technology, TBU in Zlín, Nad Stranemi 4511, 760 05 Zlín, Czech Republic
vladimir.rusnak@form-composite.com
Prejem rokopisa – received: 2014-09-10; sprejem za objavo – accepted for publication: 2014-11-24
doi:10.17222/mit.2014.223
Composite materials used in the transport industry and also in other sectors must have a certain degree of flame resistance. For
this purpose, commonly used flame retardants are based on halogen compounds in the liquid state or aluminum hydroxide in the
solid state. Solid flame retardants have a negative effect on the processing and mechanical properties. Low viscosity and rapid
wettability of fibers are very important, especially in an resin transfer molding (RTM) process.
Therefore, a new advanced matrix system based on phosphorus flame retardants was developed. The flame resistance and
mechanical properties of the composite materials produced from the new resin system were tested. Furthermore, the processing
parameters and tests are described in the article.
Keywords: flame-retardant composites, phosphorus, RTM, vacuum infusion
Kompozitni materiali, ki se uporabljajo v industriji transporta in tudi v drugih sektorjih, morajo ustrezati dolo~eni stopnji
negorljivosti. Za ta namen so splo{no uporabljani zaviralci gorenja, ki temeljijo na osnovi spojin halogenov v teko~em stanju ali
na aluminijevem hidroksidu v trdnem stanju. Zaviralci gorenja v trdnem stanju imajo negativen vpliv na izdelavo in mehanske
lastnosti. Posebno majhna viskoznost in hitra omo~ljivost vlaken sta pomembni, posebno {e pri RTM-procesu.
Zato je bila razvita nova napredna matri~na zasnova na podlagi fosfornega zaviralca gorenja. Preizku{ene so bile odpornost proti
v`igu in mehanske lastnosti kompozitnega materiala, izdelanega iz novega sistema smol. V ~lanku so opisani parametri izdelave
in preizkusi.
Klju~ne besede: negorljivi kompoziti, fosfor, RTM, vakuumska infuzija
1 INTRODUCTION
Composite materials based on a polymer matrix and
glass reinforcement are often used in transport vehicles.
The main reasons for their utilization are their low mass,
excellent formability and relatively low production costs
for low and medium manufacturing series.
When composite and sandwich structures are used in
train applications, they have to fulfill different require-
ments. The first (or the main) requirement is the fire
safety.1 In addition to the fire safety of these materials,
relatively high mechanical properties, dimensional and
thermal stabilities and health safety are required. For
example the French standard NF F01-28: Fiber rein-
forced plastic in railway rolling stock requires a mini-
mum bending strength of reinforced plastics of up to 150
MPa, while maintaining the self-extinguishing properties
according to another French standard, NF F16-101.
Polyester matrices reinforced with fiber reinforce-
ments in various forms are primarily used for these
applications. In these composite systems self-extinguish-
ing properties are primarily achieved with additions of
flame retardants to the synthetic matrices.
Flame retardants are the materials that prevent or
retard the burning of a material and are indispensable for
plastic products, electrical applications, construction
materials, textiles, etc. Generally, they can be divided
into two main groups: the reactive organic and non-reac-
tive inorganic compounds (additives). The reactive orga-
nic compounds include molecules of halogens (bromine
and chlorine), melamine compounds, phosphate com-
pounds and many others.2 Their main advantage is that
they do not change the physical properties of the resin.
Generally, the amount of the flame retardant in a poly-
mer matrix depends on the retarding effect. However, it
is necessary to take into account the possibility of a
negative effect influencing the mechanical and pro-
cessing properties of the polymer.
The reactive inorganic compounds include aluminum
hydroxide (ATH) and magnesium hydroxide, red phos-
phorus, ammonium compounds, antimony compounds of
antimony or boron, graphite, etc.3,4 These compounds are
insoluble in a polymer matrix and need to be added in
higher concentrations (30–150 %) to achieve the right
fire-retardancy level. They have a very negative influence
Materiali in tehnologije / Materials and technology 49 (2015) 5, 821–824 821
UDK 678.7 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(5)821(2015)
on the viscosity of the resulting mixture, thus limiting
their use in the preparation of a composite with the
vacuum infusion or RTM technology. And, last but not
least, they negatively influence the mechanical properties
– the composites become less stiff and brittle.
They slow down the combustion process in different
ways, physically or chemically. The physical processes
include cooling the material, creating a protective layer
(that continues to protect the remaining material, block-
ing the combustion by limiting the access of air) or
diluting the gas flame with retardant components.
With respect to the chemical processes, flame
retardancy is a reaction in the gaseous phase, leading to a
release of significant amounts of water or carbon dio-
xide. Thus, a reaction at this stage limits and slows down
a burning process. This group includes the flame
retardants based on halogens. These compounds create
carbon ash on the polymer surface or a swelling under
the protective layer of the ash to provide a better thermal
insulation.5
Some flame retardants based on halogens, such as
polychlorinated biphenyls (PCB), or various brominated
flame retardants (polybrominated diphenyl ethers, poly-
brominated biphenyls) may pose a health risk. The fire
retardancy of a halogenated flame retardant is based on
the formation of a volatile-combustion hydrobromic or
hydrochloric acid in the presence of hydrocarbons (e.g.,
toxic dioxins) at a temperature of around 350 °C.
Flame retardants based on phosphorus are the com-
pounds that are more environmentally friendly and their
consumption is slowly growing. Phosphor flame retard-
ants are already applied in matrices with phosphorus
amounts of about 5–8 %. They work very well for the
polymers containing oxygen in the chain.
2 EXPERIMENTAL WORK
2.1 Materials
2.1.1 Matrix
An unsaturated polyester resin diluted with styrene,
containing a phosphorus flame retardant incorporated
into the chain molecules was developed under the work-
ing title FR4/12 (Table 1). The goal was to prepare a
resin, which enables, due to its viscosity, a preparation of
the composite parts for rail vehicles using the RTM
method or the vacuum infusion method. Pure resin
would also meet the requirement for a flame resistance
without the additives – the limited oxygen index (LOI)
must be higher than 30. The requirement for the maxi-
mum viscosity was 600 mPa s.6 The amount of phospho-
rus in the polyester chain is naturally limited since it
greatly affects the mechanical properties. The resulting
formulation is designed in such a way that the mecha-
nical properties of pure resin match the mechanical
properties of the halogenated resins, normally used for
the production of the composite parts for rail vehicles.
Table 1: Mechanical and processing properties of FR4/12 matrix
Tabela 1: Mehanski in procesni parametri osnove FR4/12
Material description MatrixFR4/12
Viscosity (Brookfield LV 2/12, 20 °C) (mPa s) 6 475
Tg/°C 105
HDT/°C 68
Acid number, the KOH content (mg g–1) 2.5
Hydroxyl number, the KOH content (mg g–1) 50
Tensile strength (MPa) 63
Tensile modulus (MPa) 3400
Elongation (%) 2.2
Flexural strength (MPa) 3600
Flexural modulus (MPa) 130
Oxygen index 32–34
2.1.2 Reinforcements, coatings
A multiaxial nonwoven fabric (bidiagonal and tri-
axial fabric) was selected as the reinforcement. A
nonwoven fabric guarantees a high proportion of glass in
a composite (hence, better mechanical properties). It also
has a suitable drapability. In line with the requirements
of the process technology, mats with the trade name
UNIFILLO ® were used for the RTM technology.
To determine the effect of the coating on the com-
bustibility, two samples of the BÜFA gelcoat layer were
provided.
2.2 Production methods and testing
The testing of the resin system was carried out at two
levels. One level was the production of the test samples.
These samples were produced on a simple plate or in a
simple RTM mold. The second level of the matrix testing
consisted of a production of real products (prototypes).
Production methods were chosen with respect to the
environmental, health and economic aspects (producti-
vity), including two closed-production technologies –
VIP (vacuum infusion process) and RTM (resin transfer
molding).
In addition to the newly produced matrix, a total of
five composite materials (Table 2) of various com-
positions as well as other polymer systems were
produced (Sample D was made from a phenolic prepreg,
Sample E was based on a commercial matrix). Samples
B and E were provided with a gelcoat. Firstly, the
flexural test was realized according to ISO 178.7 With
respect to the fire retardancy the oxygen index was
recorded, and the tests of the flame spread and the
(partial) heat release were performed using the cone
calorimeter method.8,9
3 RESULTS AND DISCUSSION
The measured values of the mechanical properties
and fire behavior are shown in Figure 1. As can be seen,
the highest values for both flexural properties of the
samples prepared with the wet technology were obtained
V. RUSNÁK et al.: FLAME RESISTANCE AND MECHANICAL PROPERTIES OF COMPOSITES ...
822 Materiali in tehnologije / Materials and technology 49 (2015) 5, 821–824
for Sample C. Sample D prepared from prepregs showed
excellent flexural performance and fire properties
compared to the samples prepared with RTM and VIP;
however, these materials are very expensive.
Moreover, comparing FR4/12 to the commercial ma-
trix, it is obvious that the obtained mechanical properties
are almost twice as high as for the samples prepared
from FR4/12.
The tests of the fire properties (Figure 2) showed that
the gelcoat layer also influences the flame spread and the
heat release. Further, the oxygen index of Samples B
(matrix FR4/12) and E (the commercial matrix) was the
same for both samples. Sample C exhibited the lowest
oxygen index of all the produced composites. Regarding
the requirement for the oxygen index (higher than 30),
all the samples passed the test. The heat-release
parameter (Figure 3) was measured only for Samples A,
B and C, where individual values were 211.6, 154.7 and
37.7. The curves in Figure 3 show the heat release
during the combustion of the samples. The flame spread
(Figure 4) for Samples A and C is the same (due to the
same material composition).
4 CONCLUSION
The main objective of the mechanical testing was to
determine whether the strengths of the material combi-
nations reached 150 MPa (this value is the requirement
V. RUSNÁK et al.: FLAME RESISTANCE AND MECHANICAL PROPERTIES OF COMPOSITES ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 821–824 823
Figure 3: Heat release during the flame test of composite materials: a)
sample A, b) sample B
Slika 3: Spro{~anje toplote pri preizkusu kompozitnega materiala v
plamenu: a) vzorec A, b) vzorec B
Figure 1: Flexural properties of prepared samples
Slika 1: Upogibne lastnosti pripravljenih vzorcev
Figure 2: Fire properties of prepared samples
Slika 2: Po`arne lastnosti pripravljenih vzorcev
Table 2: Material characterization of composite materials
Tabela 2: Karakterizacija materialov v kompozitu
Material description Sample A Sample B Sample C Sample D Sample E
Reinforcement lay-up
1 × BIDI glass
fabric*
1 × CGM **
1 × BIDI glass
fabric*
1 × BIDI glass
fabric*
1 × CGM**
1 × BIDI glass
fabric*
1 × BIDl glass
fabric*
1 × CGM**
1 × BIDI glass
fabric*
10 ×
phenolic/glass*
prepreg
1 × BIDI glass
fabric*
1 × CGM**
1 × BIDI glass
fabric*
Matrix matrix FR4/12 matrix FR4/12 matrix FR4/12 phenolic matrix Norester 880
Additive – – – – ATH(100 % to matrix)
Coating – gelcoat BÜFA – – gelcoat BÜFA
Production technology RTM RTM VIP prepreg RTM
(*bidiagonal glass fabric, **continuous-filament glass mat)
of French standard NF F01-281). All the material com-
binations met this requirement with a certain safety
margin, except for Sample E. The result was expressed
with a high amount of the filler in the resin (the ratio of
1 : 1). It is obvious that after adding the additives the
value of the bending strength decreases.
The objective of testing the fire resistance of the
materials was to determine the level of the flame
retardancy of the composites made from the new matrix
system. Using a simple test – determination of the oxy-
gen index – the materials could be classified into the
mid-level category of flame-retardant resins. It was
assumed that with the addition of non-reactive additives
the oxygen index, together with the overall resistance to
fire, would increase. According to all the flame-retard-
ancy tests, the flame-retardant gelcoat plays an important
role in improving the fire resistance.
The oxygen index (Table 3) for various composites
based on the new matrix is very similar to the oxygen
index of pure resin. Increasing the glass amount has a
minimum impact on the oxygen index. However, it is
assumed that with the additions of non-reactive additives
to a composite system, the oxygen index will increase.
The values of the heat release for the phenolic systems
are relatively large in comparison with the values for the
composite systems based on the new matrix.
The developed matrix exhibits excellent processing
properties, good and fast fiber wetting, adjustable reac-
tivity and curing times. Another advantage is that the
used promoter is based on iron compounds.
From the overview of the properties of all the sys-
tems, it is clear that the polyester systems do not reach
the same fire-resistance values as the phenolic systems.
On the other hand, due to other advantages (surface
quality, price and processing methods) the development
of the composite systems based on polyester systems
will dominate.
Acknowledgement
This work and the project were realized with the
financial support of the Ministry of Industry and Trade
of the Czech Republic, project number FR-T13/433.
Further, this study was also supported by an internal
grant of TBU in Zlín, No. IGA/FT/2015/001, funded
from the resources for specific university research.
5 REFERENCES
1 Fire retardant polyester systems, Product guide of the Reichhold
company
2 N. Sain, S. H. Park, F. Suhara, S. Law, Flame retardant and mecha-
nical properties of natural fibre-PP composites containing magne-
sium hydroxide, Polymer Degradation and Stability, 83 (2004),
363–367, doi:10.1016/S0141-3910(03)00280-5
3 A. Genovese, R. A. Shanks, Structural and thermal interpretation of
the synergy and interactions between the fire retardants magnesium
hydroxide and zinc borate, Polymer Degradation and Stability, 92
(2007) 1, 2–13, doi:10.1016/j.polymdegradstab.2006.10.006
4 A. P. Mouritz, A. G. Gibson, Fire properties of polymer composite
materials, vol. 143, Springer, Netherlands 2006, 401, doi:10.1007/
978-1-4020-5356-6
5 G. E. Zaikov, S. M. Lomakin, S. V. Usachev, E. V. Koverzanova, N.
G. Shilkina, L. V. Ruban, New Approaches to Reduce Plastic Com-
bustibility, In: A. K. Kulshreshtha, C. Vasile (Eds.), Handbook of
Polymer Blends and Composites, Volume 2, Smithers Rapra Tech-
nology, UK 2002, 165–199
6 EN ISO 2555:1989, Plastics, Resins in the liquid state or as emul-
sions or dispersions, Determination of apparent viscosity by the
Brookfield Test method, 1989
7 ISO 178:2010, Plastics, Determination of flexural properties, 2010
8 EN ISO 4589-2:1999, Plastics, Determination of burning behaviour
by oxygen index, Ambient-temperature test, 1999
9 ISO 5658-2:2006, Reaction to fire tests, Spread of flame – Part 2:
Lateral spread on building and transport products in vertical confi-
guration, 2006
V. RUSNÁK et al.: FLAME RESISTANCE AND MECHANICAL PROPERTIES OF COMPOSITES ...
824 Materiali in tehnologije / Materials and technology 49 (2015) 5, 821–824
Figure 4: Composite samples after the flame-spread test
Slika 4: Kompozitni vzorci po preizkusu {irjenja ognja
Table 3: Table guide to selecting the test classification1
Tabela 3: Tabelarni napotek za izbiro klasifikacije preizkusa1
ASTM 2863: LIMITED
OXYGEN INDEX (LOI)
UK:
BS 476
PART 6,7
GERMANY:
4102 5510
FRANCE:
NF P 92-501
USA: ASTM E84
UL94
EUROCLASS: SINGLE
BURNING ITEM (SBI)
41– 50 Class 0 Class A2 S4 M1 25 5V 0 B
34.5–41 Class 1 Class B1 S4 M2 50 0 1 B/C
28.5– 34.5 Class 2 Class B2 S3 M3 100 0 2 C/D
22–28.5 Class 3 Class B3 S2 M4 >100 1 3 D/E
Ý. KIRIK, N. ÖZDEMÝR: EFFECT OF PROCESS PARAMETERS ON THE MICROSTRUCTURE ...
825–832
EFFECT OF PROCESS PARAMETERS ON THE
MICROSTRUCTURE AND MECHANICAL PROPERTIES OF
FRICTION-WELDED JOINTS OF AISI 1040/AISI 304L STEELS
VPLIV PROCESNIH PARAMETROV NA MIKROSTRUKTURO IN
MEHANSKE LASTNOSTI TORNO VARJENIH SPOJEV JEKEL
AISI 1040/AISI 304L
Ýhsan Kirik1, Niyazi Özdemýr2
1Batman University, Department of Metallurgy and Material Engineering, 72060 Batman, Turkey
2Fýrat University, Department of Metallurgy and Materials Engineering, 23119 Elazýð, Turkey
alihsankirik@gmail.com
Prejem rokopisa – received: 2014-09-17; sprejem za objavo – accepted for publication: 2014-10-17
doi:10.17222/mit.2014.235
Couples of AISI 304L and AISI 1040 steels were welded using the continuous-drive friction-welding process. The welded joints
were manufactured using three different rotational speeds ((1300, 1500 and 1700) r/min), three different frictional pressures and
three different frictional times. To determine microstructural changes, the interface regions of the welded samples were
examined by scanning electron microscopy (SEM), X-ray diffraction (XRD) and energy dispersive spectrometry (EDS).
Microhardness and tensile tests of the welded samples were conducted. The results show that the properties of the
microstructures are significantly changed and the interface temperature increased with the frictional time. An excellent tensile
strength was observed for the joint made at a rotational speed of 1700 r/min and a frictional time of 4 s.
Keywords: friction welding, AISI 1040 steel, AISI 304L steel
Jekli AISI 304L in AISI 1040 sta bili zvarjeni z uporabo kontinuirnega tornega procesa varjenja. Zvarjeni spoji so bili izdelani
pri treh razli~nih hitrostih vrtenja ((1300, 1500 in 1700) r/min), treh razli~nih tornih tlakih in treh razli~nih tornih ~asih. Za
dolo~itev mikrostrukturnih sprememb je bilo sti~no podro~je zvarjenega vzorca preiskano z vrsti~nim elektronskim
mikroskopom (SEM), z rentgensko difrakcijo (XRD) in z energijsko disperzijsko rentgensko spektrometrijo (EDS). Izvr{eni so
bili preizkusi mikrotrdote in natezni preizkusi varjenih vzorcev. Rezultati ka`ejo, da se mikrostruktura omembe vredno
spremeni in da se temperatura stika povi{uje s ~asom trenja. Odli~na natezna trdnost stika je bila ugotovljena pri spoju s
hitrostjo vrtenja 1700 r/min in pri tornem ~asu 4 s.
Klju~ne besede: torno varjenje, jeklo AISI 1040, jeklo AISI 304L
1 INTRODUCTION
Stainless steels are iron-based alloys containing 8–25 %
nickel and 12–30 % chromium, resisting both corrosion
and high temperature. Generally, stainless steels can be
classified as martensitic, ferritic and austenitic. Auste-
nitic stainless steels represent the largest group of
stainless steels and are produced in higher tonnages than
any other group. They have a good corrosion resistance
in most environments. Austenitic stainless steel can be
strengthened significantly by cold working and are often
used in the applications requiring a good atmospheric or
elevated-temperature corrosion resistance.
In spite of a wide variety of austenitic stainless steel,
the 300-series alloys, based on the 18Cr-8Ni system, are
the oldest and most commonly used.1–5 Type 304 is the
foundation of this alloy series and, along with 304L, it
represents the most commonly selected austenitic grade.
Although austenitic alloys are generally considered to be
very weldable, they are subjected to a number of weld-
ability problems if proper precautions are not taken.
Liquation cracking and weld solidification may occur
depending on the composition of the base, the filler
metal and the level of impurities, particularly S and P. In
spite of the good general corrosion resistance of auste-
nitic stainless steels, they may be subject to localized
forms of corrosion at grain boundaries and in the heat-
affected zone. For joining dissimilar materials, conven-
tional fusion welding is used. However, this method has
a lower efficiency when compared to the friction-weld-
ing method.5–7
The friction-welding process is a solid-state welding
process producing welds due to a compressive-force con-
tact of two workpieces: one is rotating and the other is
stationary.8,9 The heat is generated at the weld interface
due to a continuous rubbing of the contact surfaces,
while the temperature rises and softens the material.
Eventually, the material at the interface starts to flow pla-
stically and forms an upsetting. When a certain amount
of upsetting occurs, the rotation is stopped and the
compressive force is maintained or slightly increased to
consolidate the joint. The fundamental parameters of the
friction-welding (FW) process are the rotational speed,
the frictional pressure and time, and the forging pressure
and time. The advantages of friction welding are a high
material saving, a short production time and the possi-
Materiali in tehnologije / Materials and technology 49 (2015) 5, 825–832 825
UDK 621.791/.792:620.17:620.187:669.14 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(5)825(2015)
bility of joining different metals or alloys. This method
can also be used to join the components that have circu-
lar or non-circular cross-sections.10–17
In this study, parts of the AISI 304L austenitic stain-
less steel and AISI 1040 carbon steel with equal dia-
meters were welded using the friction-welding method.
The strengths of the joints were determined with tensile
tests and compared with those of the base materials.
Then, the hardness variations and microstructures of the
welding zones were obtained and investigated using
macro- and microphotographs.
2 EXPERIMENTAL PROCEDURES
The workpieces used in this work were Ø 12 mm
rods of AISI 1040 and AISI 304L austenitic stainless
steels. The nominal chemical compositions and mechani-
cal properties of the materials used are shown in Tables
1 and 2, respectively. The AISI 1040 medium-carbon
steel and the AISI 304L austenitic stainless steel were
joined with a PLC-controlled continuous-drive friction-
welding machine using different rotational speeds,
frictional times and frictional pressures, as indicated in
Table 3. The temperature of the welded zone during the
Ý. KIRIK, N. ÖZDEMÝR: EFFECT OF PROCESS PARAMETERS ON THE MICROSTRUCTURE ...
826 Materiali in tehnologije / Materials and technology 49 (2015) 5, 825–832
Table 1: Chemical compositions of test materials in mass fractions
Tabela 1: Kemijska sestava materialov za preizkuse v masnih dele`ih
Materials
Alloying elements (w/%)
C Mn P S Cr Mo Ni Fe
AISI 1040 0.37–0.44 0.60–0.90 0.040 0.050 – – – Balance
AISI 304L 0.042 1.47 0.032 0.032 18.25 0.30 8.09 Balance
Table 2: Mechanical properties of parent materials
Tabela 2: Mehanske lastnosti izhodnih materialov
Material Hardness (HV) Yield strength (MPa) Tensile strength (MPa)
AISI 1040 165 215 630
AISI 304L 210 380 700
Table 3: Process parameters used in friction welding
Tabela 3: Procesni parametri, uporabljeni pri tornem varjenju
Samp. No
Welding parameters
Rotational
speed
(r/min)
Frictional
time
(s)
Frictional
pressure
(MPa)
Forging
pressure
(MPa)
Forging
time
(s)
Axial short.
(mm)
Max. temp.
(oC)
Max.
hardness
(HV)
Max. tensile
stren.
(MPa)
S1 1300 8 30 60 4 6 1204 362 371.414
S2 1300 8 40 80 4 7 1189 311 412.431
S3 1300 8 50 100 4 8 1105 400 483.223
S4 1300 6 30 60 3 5 1176 348 387.026
S5 1300 6 40 80 3 6 1102 317 559.276
S6 1300 6 50 100 3 7 1152 312 600.061
S7 1300 4 30 60 2 4 1214 330 580.338
S8 1300 4 40 80 2 5 1124 277 599.821
S9 1300 4 50 100 2 6 1182 254 617.106
S10 1500 8 30 60 4 7 1175 320 473.096
S11 1500 8 40 80 4 8 1104 376 575.716
S12 1500 8 50 100 4 9 1123 329 596.613
S13 1500 6 40 60 3 6 1140 319 588.216
S14 1500 6 40 80 3 7 1164 337 599.296
S15 1500 6 50 100 3 8 1129 278 611.714
S16 1500 4 30 60 2 4.5 1104 356 586.311
S17 1500 4 40 80 2 5.5 1185 361 603.196
S18 1500 4 50 100 2 6 1075 256 588.807
S19 1700 8 30 60 4 8 1089 321 531.222
S20 1700 8 40 80 4 9 1067 342 563.691
S21 1700 8 50 100 4 10 1075 263 568.780
S22 1700 6 30 60 3 7 1125 349 573.808
S23 1700 6 40 80 3 7.5 1093 322 589.467
S24 1700 6 50 100 3 8.5 1143 245 571.993
S25 1700 4 30 60 2 4.5 1089 364 592.684
S26 1700 4 40 80 2 5.5 1092 300 612.923
S27 1700 4 50 100 2 6.5 1047 210 637.500
welding process was measured with an IGA 15 PLUS
device. After the friction welding, the microstructural
changes in the interface regions of the joints were
examined with scanning electron microscopy (SEM),
X-ray diffraction (XRD) and energy dispersive spectro-
metry (EDS). Microhardness and tensile tests were con-
ducted and the fracture surfaces of the tensile-test
samples were investigated.
3 RESULTS AND DISCUSSIONS
3.1 Evaluation of macro- and microstructural proper-
ties
Figures 1a and 1b show macro-images of the sur-
faces and interfaces of the friction-welded joints with
three different frictional times, respectively. From the
figures, the materials of the overflow out of the inter-
faces are seen to be clean-cut, and the widths of the
flashes are measured to be (5, 7 and 9) mm for the fric-
tional times of (4, 6 and 8) s, respectively. When joining
dissimilar materials by friction welding flash is formed,
depending on the mechanical properties of the materials.
The extent of the flash increased with the increasing
rotational speed, frictional time and frictional pressure.
The flash diameter of the joints was larger for the AISI
1040 steel than for the AISI 304L steel. There are no
cracks, cavities or disconnected regions seen on the sam-
ples joined at a rotational speed of 1700 r/min.
Table 3 shows the axial shortening of the friction-
welded joints. As it can be seen from these results, the
axial-shortening values for the materials are different
owing to the plastic deformation of the materials. The
axial shortening increases with the increasing rotational
speed, frictional time and frictional pressure due to the
increasing heat input and plastic deformation in the
interface of the friction-welded joints. The maximum
axial shortening is seen on S21 (10 mm) welded at the
following process parameters: rotational speed is 1700
Ý. KIRIK, N. ÖZDEMÝR: EFFECT OF PROCESS PARAMETERS ON THE MICROSTRUCTURE ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 825–832 827
Figure 2: SEM micrographs of the weld zones of the S1, S13 and S27 samples
Slika 2: SEM-posnetki podro~ja zvara pri vzorcih S1, S13 in S27
Figure 1: a) Surface and b) cross-sectional macro-images of FW
joints of samples S1, S13 and S27
Slika 1: a) Povr{ina in b) makrovidez pre~nega prereza FW-spoja pri
vzorcih S1, S13 in S27
r/min, frictional pressure 50 MPa, forging pressure 100
MPa, frictional time 8 s, forging time 4 s.
Figure 2 shows the joint microstructures for samples
S1, S13 and S27 obtained at three different rotational
speeds. On Figure 2 distinct regions are seen in the
heat-affected zone and the interface of the S1 joint. As
reported in the references, the evaluation of the micro-
structures of the joints revealed different zones in all the
samples that were identified as the base metal (BM), the
partially deformed zone (PDZ), the deformed zone (DZ)
and the fully plasticized deformed zone (FPDZ)13,15,17. In
S1, some disconnected regions and gaps were observed
with SEM. In addition, the widths of the FPDZ and DZ
expand, and the irregularities observed on all the samples
were due to the decreasing frictional and forging
pressure, and increasing frictional time. In the SEM
micrographs of sample S1, the width of the deformed
zone and the irregularity are increased due to a low fric-
tional pressure and a high frictional time. The pressure
caused a grain refinement in the central region of the S13
weld adjacent to the FPDZ and DZ. It was observed that
the rotational speed and frictional time affected the
weld-region geometry and the width.
In S27, a high rotational speed caused a rapid heating
to high temperatures at the interface. A greater volume
of the viscous material was pushed out of the interface
due to the increased axial shortening and temperature
gradient. In addition, in the areas adjacent to the PDZs
on three SEM images of the samples, overflows and slip
bands due to the effect of the heat and the axial pressure
are clearly seen. Away from the region of deformation
the density of the slip bands decreased due to the axial
pressure.
3.2 Hardness-test results
The hardness tests of the welded samples were
carried out in the lines perpendicular to the weld inter-
faces of the friction-welded joints (Figure 3). Similar
microhardness profiles were observed for all the welds.
The microhardness values for the weld zones were much
higher than those for the parent steels. The microhard-
ness was the highest at the interface for all the joints,
decreasing toward the main material. A partial presence
of martensite was the reason for an increase in the
hardness of the weld zone16 of the AISI 1040 steel, while
a higher hardness of steel AISI 304L was due to the
grain refinement.9,10,17 Among the joints, the highest
hardness value was measured for sample S1, due to a low
rotational speed and a higher frictional time causing a
more severe plastic deformation.
The temperatures of the weld zones, during the
friction welding, were measured with the IGA 15 PLUS
infrared temperature-measurement device (Figure 4) and
they were between 1025 °C and 1214 °C. The tem-
perature of a weld zone was directly affected and varied
by the increasing frictional time. The increasing hardness
of the interface of a friction-welded joint and the degree
of deformation varied in dependence of the decreasing
rotational speed. With a low rotational speed, the mate-
rial of the interface was not sufficiently viscous. How-
Ý. KIRIK, N. ÖZDEMÝR: EFFECT OF PROCESS PARAMETERS ON THE MICROSTRUCTURE ...
828 Materiali in tehnologije / Materials and technology 49 (2015) 5, 825–832
Figure 4: Temperatures of friction-welded joints
Slika 4: Temperature pri tornem varjenju spojev
Figure 3: Typical microhardness distributions perpendicular to the
interfaces of the joints: a) S1-S9, b) S10-S18 and c) S19-S27
Slika 3: Zna~ilen potek trdote pravokotno na ravnino spoja vzorcev:
a) S1-S9, b) S10-S18 in c) S19-S27
ever, looking at the hardness values for all the samples,
the hardness values for the AISI 304L austenitic stainless
steels were higher than for the AISI 1040 medium-
carbon steels due to the deformation hardening. The
main reasons for this difference are the heat-transfer
capacity and the yield strength of the stainless steels that
are 60 % higher than those of the medium-carbon
steels.18,19 The lowest hardness and the lowest tempera-
ture were obtained for specimen S27, at the rotational
speed of 1700 r/min, frictional pressure of 50 MPa and
frictional time of 4 s.
3.3 Tensile-test results
Figure 5 shows typical tensile strengths for the
friction-welded joints consisting of dissimilar materials
of AISI 304L and AISI 1040 tested at a strain rate of
1 · 10–2 s–1. The tensile fracture surfaces of the friction-
welded joints are illustrated in Figure 6. All the samples
failed in the area adjacent to the HAZ on the AISI 1040
side. The failure distance from the welding line increases
depending on the higher rotational speed and frictional
pressure (Table 4). The tensile strength of the welded
joints depends markedly on the rotational speed and
frictional time. The tensile-strength values for the FW
joints were from 371 MPa to 637 MPa, and the highest
tensile strength was exhibited by S27 welded at the rate
of 1700 r/min and frictional time of 4 s. The increase in
the tensile strength was due to the heat input and high
plastic deformation that occurred at the interface of both
steels as a result of the rotational speed and frictional
pressure.
The lowest tensile-strength value was obtained for
sample S1 at 1300 r/min. This decrease in the tensile
strength is associated with the width of the HAZ and the
reactions in the AISI 1040 HAZ17, causing a grain refine-
ment of the microstructure, the -phase and the
-ferrite10,17,20 in the HAZ of AISI 304L. The occurrence
of Cr23C6, Cr7C3 and Fe7C5 was confirmed with an X-ray
analysis and an EDS analysis of the friction-welded
joints and the occurrence of these phases in the HAZs
has an important effect on the tensile strengths of the
joints. Additionally, the formation of these phases could
partly increase the hardness at the interface.21 Due to a
longer frictional time the volume of the viscous material
Ý. KIRIK, N. ÖZDEMÝR: EFFECT OF PROCESS PARAMETERS ON THE MICROSTRUCTURE ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 825–832 829
Figure 6: Typical failure locations on the tensile-test samples of
friction-welded joints
Slika 6: Zna~ilno podro~je poru{itve pri nateznem preizkusu vzorcev
s torno zvarjenim spojem
Figure 5: Tensile-test results for friction-welded joints of dissimilar
materials
Slika 5: Rezultati nateznih preizkusov torno zvarjenih spojev razli~nih
materialov
Table 4: Tensile results for AISI 304L/AISI 1040 friction welds
Tabela 4: Rezultati nateznih preizkusov torno zvarjenih spojev AISI
304L/AISI 1040
Sam-
ple
No
Lo/
mm
Ln/
mm
Do/
mm
AISI 304L AISI 1040 Fracture
distance
x/mm
Ly/
mm
Dy/
mm
Lt/
mm
Dt/
mm
S1 40 42.1 8 22.6 7.80 20.5 7.30 1.60
S2 40 42.4 8 19.6 7.80 21.6 7.20 1,20
S3 40 43.6 8 22.6 7.49 20.8 7.20 1.60
S4 40 42.9 8 22.5 7.60 21.2 7.00 2.50
S5 40 43.8 8 19.6 7.30 24.1 6.20 3.20
S6 40 43.5 8 24.8 7.45 19.4 6.10 3.80
S7 40 43.9 8 24.0 7.33 19.6 6.00 3.55
S8 40 43.1 8 23.4 7.40 18.5 6.80 3.10
S9 40 43.7 8 23.3 7.30 20.3 6.30 3.00
S10 40 42.7 8 23.4 7.20 19.6 6.80 2.10
S11 40 43.4 8 23.9 7.20 19.8 6.30 3.20
S12 40 43.3 8 24.5 7.30 19.6 6.40 3.40
S13 40 43.7 8 24.4 7.20 19.7 6.20 3.00
S14 40 43.8 8 23.7 7.20 20.1 6.00 3.00
S15 40 43.6 8 24.5 7.30 19.4 6.50 2.60
S16 40 44.0 8 24.0 7.30 19.6 6.00 2.80
S17 40 43.0 8 23.9 7.20 19.3 6.20 3.60
S18 40 43.3 8 25.0 7.30 19.1 6.40 3.00
S19 40 43.7 8 24.0 7.30 19.8 6.50 3.50
S20 40 44.0 8 24.2 7.30 19.6 6.00 4.20
S21 40 44.0 8 24.8 7.20 20.6 6.40 3.00
S22 40 42.5 8 23.7 7.20 18.2 6.30 3.10
S23 40 43.6 8 25.2 7.20 20.1 6.00 4.00
S24 40 43.7 8 24.1 7.30 19.4 6.20 3.90
S25 40 43.9 8 23.5 7.30 19.9 6.40 3.40
S26 40 44.0 8 23.7 7.30 19.4 6.00 3.75
S27 40 44.1 8 23.6 7.20 20.4 5.80 3.50
transferred at the weld interface decreased and the tensile
strength was lowered (Figure 5). On the other hand, the
tensile strength increased with a higher rotational speed
and a lower frictional time. This shows that the forma-
tion and the width of the FPDZ, which is a result of the
heat input and plastic deformation at the interface, have
detrimental effects on the mechanical strength and,
consequently, on the process parameters of dissimilar
friction-welded joints.
3.4 X-ray and EDS analyses
The results of the XRD analysis for friction-welded
joint S13 are shown in Figure 7. The results confirm the
presence of the Cr23C6, Cr7C3 and Fe5C2 compounds at
the interface of the welded joints. SEM micrographs and
EDS-analysis points of friction-welded joints S1, S13
and S27 are presented in Figure 2. In all the SEM
images, the following EDS-analysis points are marked:
point 1 (the AISI 1040 side), point 2 (the deformed
zone), point 3 (the fully plasticized deformed zone),
point 4 (the deformed zone of AISI 304L) and point 5
(on the AISI 304L side). The EDS-analysis results of
samples of S1, S13 and S27 are also given in Table 5.
The results of the elemental analysis and microstruc-
tural examination in the interface regions of the
friction-welded joints clearly demonstrated that different
amounts of C, Fe, Cr, Si, Ni and Mn were obtained.
Additionally, the EDS analysis of the interfaces of the
friction-welded samples revealed that there was a tran-
sition of Cr and Ni from AISI 304L to AISI 1040.
Furthermore, there was a Fe transition from AISI 1040 to
Ý. KIRIK, N. ÖZDEMÝR: EFFECT OF PROCESS PARAMETERS ON THE MICROSTRUCTURE ...
830 Materiali in tehnologije / Materials and technology 49 (2015) 5, 825–832
Figure 9: SEM micrographs of fracture surfaces after the tensile test of friction-welded samples S1, S13 and S27
Slika 9: SEM-posnetki provr{ine preloma po nateznem preizkusu torno zvarjenih vzorcev S1, S13 in S27
Figure 8: Concentration profiles of EDS analysis across the welding
interface of friction-welded sample
Slika 8: Profil koncentracij pri EDS-analizi prereza stika pri torno
zvarjenem vzorcu
Figure 7: Result of XRD analysis for sample S13
Slika 7: Rentgenogram vzorca S13
Table 5: EDS analyses across the welding interfaces of friction-
welded samples S1, S13 and S27
Tabela 5: EDS-analiza preko ravnine spoja torno varjenih vzorcev S1,
S13 in S27
Sample
No
EDS
Points
Alloying elements (w/%)
C Si Cr Mn Fe Ni
S1
point 1 0.62 0.22 0.54 2.54 92.7 0.68
point 2 0.53 0.19 1.66 3.06 90 0.82
point 3 0.55 0.33 10.06 3.64 79.4 4.89
point 4 0.02 0.34 15.3 4.15 72.3 8.34
point 5 0.37 0.49 16.45 3.74 65.43 8.03
S13
point 1 0.17 0.26 0.48 6.57 86.9 0.74
point 2 0.59 0.31 6.07 6.86 81.17 2.98
point 3 0.29 0.45 17.3 6.60 66.5 8.40
point 4 0.24 0.50 17.46 6.58 62.38 9.61
point 5 0.10 0.59 17.5 6.27 68.6 9.53
S27
point 1 0.04 0.21 0.53 6.38 88.3 0.69
point 2 0.69 0.19 0.68 6.03 90.4 0.94
point 3 0.03 0.45 10.7 6.21 74.3 4.23
point 4 0.45 0.44 12.8 5.95 68.8 4.99
point 5 0.39 0.44 18.1 6.90 62.3 8.72
the AISI 304L austenitic stainless steel, as seen in the
concentration profiles of S13 in Figure 8.
3.5 Fractography
Figure 9 shows images of the fracture surfaces after
the tensile test of friction-welded samples S1, S13 and
S27. Examining the fracture-surface images, it was
found that the fractures resulting from the tensile test
mostly occurred on the AISI 1040 side, and especially
ductile fractures in the form of quasi-cleavages were
observed in the dimples. Moreover, all the samples
showed a ductile fracture behavior; as seen on the SEM
photographs of the fracture surfaces of the FW joints,
these images resemble spider webs.
4 CONCLUSIONS
In this study, the following conclusions were made
after a thorough investigation of the effects of the
process parameters on the microstructure and mechanical
properties of the friction-welded joints of the AISI
1040/AISI 304L steels:
1. The friction-welded AISI 1040 and AISI 304L steels
were free of pores and cracks. A grain-size reduction
occurred in the DZ of both materials. It was observed
with SEM and EDS analyses that a transition of the
materials occurred in the friction-weld zone.
2. Microstructural studies showed a presence of four
different regions at the welding interface: a fully
plasticized deformed zone, a partially deformed zone,
a deformed zone and the base materials. The widths
of the FPDZ and DZ of the joints were mainly
affected by the frictional time and rotational speed.
The use of a higher rotational speed and a lower
frictional time increased the tensile strength of a
friction-welded AISI 1040/AISI 304L steel couple.
This shows that the formation of a FPDZ, which is a
result of the heat input and plastic deformation at the
interface, has a detrimental effect on the mechanical
strength and, consequently, on the process parameters
of dissimilar friction-welded joints, therefore, it
needs to be controlled.
3. The highest tensile strength and a low microhardness
were obtained very close to the AISI 1040 parent
material of friction-welded sample S27 using a
rotational speed of 1700 r/min, a frictional pressure
of 50 MPa and a frictional time of 4 s. Depending on
the formation of the FPDZ and DZ, a decrease in the
tensile strength was detected because of the grain-
refined microstructure, -phase, -ferrite and an
increase in the frictional time.
Acknowledgements
The authors wish to thank the University of Fýrat
Research Fund for the support of this work under the
FUBAP-2054 project.
5 REFERENCES
1 A. Aran, Stainless steel production standards and using, STY,
Istanbul, Turkey, 2003
2 ASM, Metals handbook, vol.8, 1973, 424
3 R. J. Castro, J. J. De Cadenet, Welding metallurgy of stainless steel
and heat- resisting steels, Cambridge University Press, Cambridge
1974
4 J. J. Demo, Structure, Constitution, and General Characteristics of
Wrought Ferritic Stainless Steels, ASTM, Philadelphia 1977,
doi:10.1520/STP619-EB
5 J. C. Lippold, D. J. Kotecki, Welding metallurgy and weldability of
stainless steels, 1st ed., Wiley-Interscience, 2005
6 M. Erdoðan, Materials science and engineering materials 1, Ankara,
Turkey, 2002, 326–331
7 K. G. K. Murti, S. Sundaresan, Parameter optimization in friction
welding dissimilar materials, Metal Construction, (1983), 331–335
8 K. S. Mortensen, C. G. Jensen, L. C. Conrad, F. Losee, Mechanical
properties and microstructures of intertia friction welded 416
stainless steel, Welding Journal, 80 (2001) 11, 268–273
9 M. Sahin, H. E. Akata, An experimental study on friction welding of
medium carbon and austenitic stainless steel components, Industrial
Lubrication and Tribology, 56 (2004) 2, 122–129, doi:10.1108/
00368790410524074
10 N. Özdemir, Investigation of the mechanical properties of friction-
welded joints between AISI 304L and AISI 4340 steel as a function
rotational speed, Materials Letters, 59 (2005) 19–20, 2504–2509,
doi:10.1016/j.matlet.2005.03.034
11 V. V. Satyanarayana, G. Madhusudhan Reddy, T. Mohandas,
Dissimilar metal friction welding of austenitic-ferritic stainless
steels, Journal of Materials Processing Technology, 160 (2005) 2,
128–137, doi:10.1016/j.jmatprotec.2004.05.017
12 S. Celik, I. Ersozlu, Investigation of the mechanical properties and
microstructure of friction welded joints between AISI 4140 and 1050
steels, Materials & Design, 30 (2009) 4, 970–976, doi:10.1016/
j.matdes.2008.06.070
13 I. Kirik, N. Özdemir, Weldability and joining characteristics of AISI
420/AISI 1020 steels using friction welding, International Journal of
Materials Research, 104 (2013) 8, 769–775, doi:10.3139/146.110917
14 P. Sathiya, S. Aravindan, A. Noorul Haq, K. Paneerselvam, Optimi-
zation of friction welding parameters using evolutionary computa-
tional techniques, Journal of Materials Processing Technology, 209
(2009) 5, 2576-2584, doi:10.1016/j.jmatprotec.2008.06.030
15 Ý. Kýrýk, N. Özdemir, F. Sarsilmaz, Microstructure and Mechanical
Behaviour of Friction Welded AISI 2205/AISI 1040 Steel Joints,
Materials Testing, 54 (2012) 10, 683–687, doi:10.3139/120.110379
16 O. Torun, I. Celikyürek, B. Baksan, Friction welding of Fe-28Al
alloy, Intermetallics, 19 (2011) 7, 1076–1079, doi:10.1016/j.intermet.
2011.02.009
17 N. Özdemir, F. Sarsilmaz, A. Hasçalik, Effect of rotational speed on
the interface properties of friction-welded AISI 304L to 4340 steel,
Materials & Design, 28 (2007) 1, 301–307, doi:10.1016/j.matdes.
2005.06.011
18 C. R. G. Ellis, Friction Welding, Some Recent Applications of
Friction Welding, Weld. and Metal Fab., (1977), 207–213
19 H. Ates, M. Turker, A. Kurt, Effect of friction pressure on the pro-
perties of friction welded MA956 iron-based superalloy, Materials &
Design, 28 (2007) 3, 948–953, doi:10.1016/j.matdes.2005.09.015
Ý. KIRIK, N. ÖZDEMÝR: EFFECT OF PROCESS PARAMETERS ON THE MICROSTRUCTURE ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 825–832 831
20 F. D. Duffin, A. S. Bahrani, Frictional Behaviour of Mild Steel in
Friction Welding, Wear, 26 (1973) 1, 53–74, doi:10.1016/0043-1648
(73)90150-6
21 D. Ananthapadmanaban, V. Seshagiri Rao, N. Abraham, K. Prasad
Rao, A study of mechanical properties of friction welded mild steel
to stainless steel joints, Materials & Design, 30 (2009) 7, 2642–2646,
doi:10.1016/j.matdes.2008.10.030
Ý. KIRIK, N. ÖZDEMÝR: EFFECT OF PROCESS PARAMETERS ON THE MICROSTRUCTURE ...
832 Materiali in tehnologije / Materials and technology 49 (2015) 5, 825–832
H. AVDU[INOVI] et al.: INFLUENCE OF INOCULATION METHODS AND THE AMOUNT OF AN ADDED INOCULANT ...
833–836
INFLUENCE OF INOCULATION METHODS AND THE AMOUNT
OF AN ADDED INOCULANT ON THE MECHANICAL
PROPERTIES OF DUCTILE IRON
VPLIV METOD MODIFIKACIJE IN KOLI^INE DODANEGA
MODIFIKATORJA NA MEHANSKE LASTNOSTI DUKTILNEGA
@ELEZA
Hasan Avdu{inovi}1, Almaida Gigovi}-Geki}1, Diana ]ubela1,
Raza Sunulahpa{i}1, Nermin Mujezinovi}2
1University of Zenica, Faculty of Metallurgy and Materials Science, Travni~ka cesta 1, 72000 Zenica, Bosnia and Herzegovina
2CIMOS TMD CASTING, d. o. o., Zenica, Radna zona Zenica-1 bb, 72000 Zenica, Bosnia and Herzegovina
hasan.avdusinovic@famm.unze.ba
Prejem rokopisa – received: 2014-09-30; sprejem za objavo – accepted for publication: 2014-10-13
doi:10.17222/mit.2014.248
In most cases an addition of inoculants to molten cast iron is advisable and even necessary to produce good-quality castings.
The mechanical properties and machinability of cast iron with nodular graphite greatly depend on the formation of graphite and
the matrix microstructure and both are significantly influenced by the inoculation treatment. The mechanism of inoculation, the
influence of the inoculation method and the amount of added inoculants are presented.
Keywords: ductile iron, inoculation, microstructure, graphite, metallic matrix
V ve~ini primerov je dodatek modifikatorjev staljenemu livnemu `elezu priporo~ljiv in celo potreben za izdelavo dobrih ulitkov.
Mehanske lastnosti in obdelovalnost litega `eleza z nodularnim grafitom so mo~no odvisne od oblike grafita in od
mikrostrukture osnove, na oboje pa mo~no vpliva obdelava z modifikatorji. Predstavljen je mehanizem modifikacije, vpliv
metode modifikacije in vpliv koli~ine dodanega modifikatorja.
Klju~ne besede: duktilno `elezo, inokulacija, mikrostruktura, grafit, kovinska osnova
1 INTRODUCTION
Inoculants are the materials added to molten cast
iron, modifying the microstructure and, thereby, chang-
ing the physical and mechanical properties to a degree
not explained on the basis of the change in the compo-
sition. Inoculation is a phase of the technological process
of producing ductile iron that controls and improves the
microstructure and mechanical properties of castings.
Through inoculation, graphite nucleation and eutectic
undercooling of the melt can be controlled, which is
crucial for achieving the required service properties of
the castings.1–3
The inoculant most commonly used in foundries is a
ferrosilicon alloy with precisely defined contents of Ca,
Ba, Sr, Zr, Al and rare-earth elements. The benefits of
inoculations are: an improved machinability, increased
strength and ductility, reduced hardness and fracture sen-
sitivity, a more homogenous microstructure, a reduced
tendency for solidification shrinkage, etc.
During modularizing, numerous inclusions with a
sulfide core and an outer shell containing complex mag-
nesium silicates are formed. Such micro-inclusions do
not provide an effective nucleation for graphite because
the crystal lattice structure of magnesium silicates does
not match sufficiently with the lattice structure of gra-
phite. After the inoculation with a ferrosilicon alloy con-
taining Ca, Ba or Sr, the surfaces of the micro-inclusions
are modified and other complex Ca, Sr, or Ba silicate
layers are obtained. Such silicates have the same hexago-
nal crystal lattice as graphite and act as effective nucle-
ation sites for graphite nodules to grow during solidifica-
tion.4
The required rate of adding inoculants to a liquid de-
pends very much on the place and time of their inocula-
tion.
2 EXPERIMENTAL PROCEDURE
In the investigations of the response of ductile-iron
test melts to different amounts of added inoculants and
different inoculation methods, three types of the inocu-
lation method were applied: the ladle inoculation, the
in-stream inoculation and the in-mold inoculation. The
partner foundry in this research project was "Cimos
TMD Casting" in Zenica.
In the first project step, 15 melts were prepared. Full
lists of the chemical compositions, the amounts of added
inoculants and the places of the inoculant introduction
are included in Table 1. High-purity FeSi with additions
of Al, Ca, and Sr was used for the inoculation. The test
Materiali in tehnologije / Materials and technology 49 (2015) 5, 833–836 833
UDK 669.18:621.74 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(5)833(2015)
iron melts were obtained using an induction furnace and
casting iron into a green-sand mold. After the casts were
cooled down and cleaned, samples for the mechanical-
property examination were prepared.1
The testing of the mechanical properties was per-
formed according to BAS EN 100002:2002 relating to
tensile testing and BAS EN 6507-1/2007 relating to
hardness testing of fractured tensile-test bars.
Based on the summary of the results of testing the
casting properties, conclusions about the optimum
amounts of the added inoculants needed for obtaining the
required values of the mechanical properties, were
drawn.
Casting procedures involving the exact amounts of
the added inoculants required for three different methods
H. AVDU[INOVI] et al.: INFLUENCE OF INOCULATION METHODS AND THE AMOUNT OF AN ADDED INOCULANT ...
834 Materiali in tehnologije / Materials and technology 49 (2015) 5, 833–836
Table 1: Chemical compositions of the experimental melts, places of the inoculation and added amounts of the inoculants (first step of the
project)
Tabela 1: Kemijska sestava eksperimentalnih talin, mesto dodajanja modifikatorja in koli~ina dodanega modifikatorja (prva stopnja projekta)
No. Chemical composition (%) Inoculant (%) Temp.
°CC Si Mn P S Cr Cu Mg stream mold ladle
1 3.52 2.17 0.428 0.016 0.006 0.086 0.314 0.043 0 0 0 1388
2 3.51 2.20 0.433 0.016 0.006 0.090 0.314 0.040 0 0.05 0 1392
3 3.49 2.06 0.401 0.015 0.006 0.089 0.332 0.032 0 0.15 0 1402
4 3.58 2.15 0.403 0.014 0.005 0.076 0.346 0.043 0 0.20 0 1407
5 3.56 2.09 0.482 0.011 0.003 0.111 0.340 0.040 0.05 0.05 0 1395
6 3.44 2.11 0.485 0.012 0.003 0.095 0.356 0.047 0.05 0.10 0 1410
7 3.60 2.14 0.345 0.016 0.006 0.054 0.080 0.053 0.10 0 0 1380
8 3.54 2.10 0.337 0.017 0.009 0.071 0.090 0.055 0.05 0 0 1380
9 3.52 2.15 0.362 0.014 0.007 0.064 0.091 0.045 0.15 0 0 1390
10 3.56 2.11 0.398 0.017 0.008 0.068 0.316 0.047 0.10 0.05 0 1400
11 3.62 2.45 0.353 0.017 0.009 0.049 0.119 0.065 0.20 0 0 1387
12 3.65 2.11 0.392 0.019 0.012 0.054 0.149 0.079 0.15 0.05 0 1405
13 3.62 2.12 0.382 0.019 0.011 0.054 0.144 0.060 0 0.20 0 1390
14 3.61 2.11 0.357 0.018 0.009 0.037 0.187 0.048 0 0 0.05 1398
15 3.64 2.14 0.384 0.016 0.011 0.084 0.212 0.052 0 0 0.20 1402
Table 2: Range of the chemical compositions of the test melts (second step of the project)
Tabela 2: Obseg kemijskih sestav preizkusnih talin (druga stopnja projekta)
Composition C Si Mn S P Mg
w/% 3.4–3.6 2.1–2.3 0.2–0.4 0.003–0.012 0.01–0.02 0.003–0.055
Figure 3: In-mold inoculation, Nital etched
Slika 3: Modifikacija v kokili, jedkano v nitalu
Figure 2: In-stream inoculation, Nital etched
Slika 2: Modifikacija v curek, jedkano v nitalu
Figure 1: Ladle inoculation, Nital etched
Slika 1: Modifikacija v ponvi, jedkano v nitalu
of inoculation were conducted in the second step to find
the most suitable casting technology in the partner
foundry. A total of 60 melts were prepared in the second
step of the project and the average melt chemical com-
position is presented in Table 2.
3 RESULTS AND DISCUSSION
Figures 1 to 3 show the microstructures of the sam-
ples taken from the melts (castings) inoculated with
different methods; the tensile strengths and hardness
values are presented in Figures 4 and 5.
The required hardness was in the range of 180–220
HB and the tensile strength was in the range of 500–550
MPa. According to the results in Figures 1 and 2, the
best result was observed with the samples with 0.1 % of
the inoculants. In the second project step we used 20
melts for each inoculation method, i.e., 60 melts were
treated with 0.1 % of the added inoculants. The mecha-
nical properties are shown in Figures 6 and 7 and in
Table 3.
4 CONCLUSIONS
Experimental materials with quite satisfactory cha-
racteristics were prepared with vacuum induction melt-
ing. The results obtained in the first project step showed
that with 0.1 % of the inoculant the targeted hardness
and tensile strength for the given ductile iron were
achieved.
In the second project step it was found that the
in-stream inoculation ensured the best result as almost
100 % of all the samples achieved the targeted values for
the hardness and tensile strength and the standard devi-
ation was the smallest.
Acknowledgment
The authors would like to thank the Federal Ministry
of Education and Science of the Federation of Bosnia
H. AVDU[INOVI] et al.: INFLUENCE OF INOCULATION METHODS AND THE AMOUNT OF AN ADDED INOCULANT ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 833–836 835
Figure 7: Tensile-strength values of the tested samples (second step of
the project); black horizontal bars – targeted values
Slika 7: Natezna trdnost preizkusnih vzorcev (druga stopnja projekta);
~rne vodoravne ~rte – ciljne vrednosti
Figure 5: Tensile-strength values of the tested samples (first step of
the project)
Slika 5: Natezna trdnost preizkusnih vzorcev (prva stopnja projekta)
Table 3: Average values and standard deviations of the hardness and
tensile strength (second step of the project)
Tabela 3: Povpre~ne vrednosti in standardna deviacija trdote in
natezne trdnosti (druga stopnja projekta)
Hardness (HB) Tensile strength (MPa)
Inoculatio
n method In mold
In
stream Ladle In mold
In
stream Ladle
Average 212.15 203.8 205.3 564.65 546.3 474.95
St.
deviation 7.86 6.88 27.02 29.1 9.41 33.61
Figure 4: Hardness values of the tested samples (first step of the
project)
Slika 4: Trdota preizkusnih vzorcev (prva stopnja projekta) Figure 6: Hardnes values of the tested samples (second step of the
project); black horizontal bars – targeted values
Slika 6: Trdota preizkusnih vzorcev (druga stopnja projekta); ~rne
vodoravne ~rte – ciljne vrednosti
and Herzegovina and Foundry "CIMOS TMD CAST-
ING" in Zenica for their financial aid and cooperation
during the implementation of the research project.
5 REFERENCES
1 H. Bornstein et al., Cast Metals Handbook, American Foundrymen’s
Society, Des Plaines, Illinois, USA 1957, 31
2 E. Fras, M. Gorny, Inoculation Effects of Cast Iron, Archives of
Foundry Engineering, 12 (2012) 4, 39–46
3 J. N. Harvey, G. A. Noble, Inoculation of Cast Irons – An Overview,
55th Indian Foundry Congress, India, 2007, 343–360
4 T. Skaland, A New Approach to Ductile Iron Inoculation, American
Foundry Society, Schaumburg, Illinois, USA 2001
H. AVDU[INOVI] et al.: INFLUENCE OF INOCULATION METHODS AND THE AMOUNT OF AN ADDED INOCULANT ...
836 Materiali in tehnologije / Materials and technology 49 (2015) 5, 833–836
S. S. LOW et al.: ONE-STEP GREEN SYNTHESIS OF GRAPHENE/ZnO NANOCOMPOSITES ...
837–840
ONE-STEP GREEN SYNTHESIS OF GRAPHENE/ZnO
NANOCOMPOSITES FOR NON-ENZYMATIC HYDROGEN
PEROXIDE SENSING
ENOSTOPENJSKA ZELENA SINTEZA NANOKOMPOZITA
GRAFEN-ZnO ZA NEENCIMATSKO DETEKCIJO VODIKOVEGA
PEROKSIDA
Sze Shin Low1, Michelle T. T. Tan1, Poi Sim Khiew1, Hwei-San Loh2,
Wee Siong Chiu3
1Division of Materials, Mechanics and Structures, Center of Nanotechnology and Advanced Materials, Faculty of Engineering, University of
Nottingham Malaysia Campus, Jalan Broga, 43500 Semenyih, Selangor, Malaysia
2School of Biosciences, Faculty of Science, University of Nottingham Malaysia Campus, Jalan Broga, 43500 Semenyih, Selangor, Malaysia
3Low Dimensional Materials Research Center, Department of Physics, Faculty of Science, University Malaya, 50603 Kuala Lumpur, Malaysia
michelle.tan@nottingham.edu.my
Prejem rokopisa – received: 2014-10-14; sprejem za objavo – accepted for publication: 2014-11-17
doi:10.17222/mit.2014.259
In the present study, a disposable electrochemical biosensor for hydrogen peroxide (H2O2) was fabricated using a graphene/ZnO
nanocomposite-modified screen-printed carbon electrode (SPCE). The adopted method is simple, cost feasible and it avoids the
usage of harsh oxidants/acids during the synthesis. Graphite material was subjected to liquid-phase exfoliation with the aid of
ultrasonication, without going through the intermediate graphene-oxide phase that can disrupt the pristine structure of the yield.
The as-prepared graphene/ZnO nanocomposite was then thoroughly characterized to evaluate its morphology, crystallinity, com-
position and product purity. All the results clearly indicate that pristine graphene was successfully produced with the graphite
exfoliation and ZnO nanoparticles were homogeneously distributed on the graphene sheet, without any severe aggregation. The
biosensing capability of the graphene/ZnO nanocomposite-modified SPCE was electrochemically evaluated with cyclic
voltammetry (CV) and an amperometric analysis. The resulting electrode is found to exhibit an excellent electrocatalytic
activity towards the reduction of H2O2. The graphene/ZnO-modified SPCE can detect H2O2 in a linear range of 1 mM to 15 mM
with a correlation coefficient of 0.9859. The electrode is found to have a high sensitivity, selectivity and superior reproducibility
for the non-enzymatic detection of an H2O2 compound.
Keywords: graphene/ZnO nanocomposite, sonochemical facile synthesis, electrochemical sensor, screen-printed carbon
electrode, hydrogen peroxide
V predstavljeni {tudiji je bil izdelan elektrokemijski biosensor za enkratno uporabo za zaznavanje vodikovega peroksida (H2O2)
z uporabo sitotiskane ogljikove elektrode (SPCE), modificirane z nanokompozitom grafen-ZnO. Ta prilagojena metoda je
enostavna, stro{kovno ugodna in omogo~a, da se izognemo uporabi ostrih snovi, oksidanti-kislina, med sintezo. Grafitni
material je bil izpostavljen ultrazvo~ni obdelavi, lu{~enju v teko~i fazi, ne da bi bil {el skozi vmesno fazo grafenovega oksida, ki
lahko moti neokrnjeno zgradbo pridelka. Pripravljeni nanokompozit grafen-ZnO je bil nato skrbno karakteriziran, ocenjena je
bila morfologija, kristalini~nost, sestava in tudi ~istost proizvoda. Rezultati jasno ka`ejo, da je bil neokrnjen grafen izdelan z
glajenjem grafita, ZnO nanodelci pa so bili homogeno razporejeni na plo{~ici iz grafena brez ve~jih segregacij. Biolo{ka
ob~utljivost plo{~ice SPCE, modificirane z nanokompozitom grafen-ZnO, je bila ugotovljena elektrokemijsko s cikli~no
voltametrijo (CV) in amperometri~no analizo. Ugotovljeno je bilo, da ima dobljena elektroda odli~no elektrokataliti~nost za
redukcijo H2O2. SPCE, modificirana z grafen-ZnO, lahko zazna H2O2 v linearnem podro~ju od 1 mM do 15 mM s korelacijskim
koeficientom 0,9859. Ugotovljeno je bilo, da ima elektroda veliko ob~utljivost, selektivnost in bolj{o ponovljivost pri
neencimatski detekciji spojine H2O2.
Klju~ne besede: nanokompozit grafen-ZnO, sonokemijska sinteza, elektrokemijski senzor, sitotiskarsko natisnjena ogljikova
elektroda, vodikov peroksid
1 INTRODUCTION
Graphene, a 2D one-atom-thick sheet of sp2 carbon
atoms arranged in a honeycomb lattice, was discovered
by Novoselov et al.1 in 2004. Ever since, it has intrigued
enormous scientific activities due to its extraordinary
properties. Graphene-based composites are anticipated to
provide groundbreaking properties in new applications
as they improve those of the host material.
Zinc oxide (ZnO), a II-VI compound semiconductor
has been receiving considerable attention in the scientific
community due to its unique properties. The wide-band-
gap properties of ZnO have triggered tremendous
research attention for its potential applications in
electronic and optoelectronic devices. Besides, ZnO is
also popular in the field of biosensing because it is bio-
compatible, biodegradable, biosafe as it can be dissolved
into mineral ions within a few hours.2
For the purpose of hybridization of graphene with
ZnO, researchers explored unique attractive properties
like enhanced photocatalytic performance,3 energy-sto-
rage property,4 sensing property,5 optoelectronic proper-
ty6 and ultrafast, nonlinear, optical switching property.7
These intriguing properties of the composite are due to
Materiali in tehnologije / Materials and technology 49 (2015) 5, 837–840 837
UDK 620.3:66.017:544.6 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(5)837(2015)
the synergistic effect between graphene and ZnO, in
which the carbon-based material acts as a good electron
conductor and when coupled with a metal oxide, the
charge transfer between the materials is improved.
There are many synthesis routes to combine graphene
sheets and ZnO nanoparticles such as thermal decompo-
sition,8 electrochemical route,9 ultrasonic spray pyro-
lysis,10 electro-hydrodynamic atomization11 and solvo-
thermal process.12 However, most of the attempts in the
synthesis of a graphene/ZnO nanocomposite involve a
reduction of the graphene oxide and an in-situ growth of
ZnO which culminate in the presence of an oxygenated
functional group on the graphene surface and impurities
in the composite due to an incomplete reduction and
purification process.
In this work, a graphene/ZnO nanocomposite was
synthesised with one-step, low-cost, green approach, in
which conductive graphene sheets serve as superior plat-
forms for a deposition of flower-like ZnO nanoparticles.
In comparison with the other work reporting on the syn-
thesis of a graphene/ZnO composite, this synthesis
method is novel and environmentally friendly because it
omits the use of harsh chemicals for the exfoliation of
graphite and growth of ZnO, resulting in a safer synthe-
tic procedure and also preserving the pristine quality of
both materials. From an economical perspective, the
starting materials are inexpensive, while ultrasonication
does not consume much power and can be conducted in a
laboratory or scaled for a mass production. The as-syn-
thesized graphene/ZnO nanocomposite was used to mo-
dify electrodes to demonstrate its viability for biosensing
applications.
2 METHODOLOGY
2.1 Preparation of graphene and a graphene/ZnO
nanocomposite
Graphite flakes (Bay Carbon, USA) were subjected
to the one-step, green liquid-phase exfoliation method to
synthesize graphene.13 Graphite flakes were added into
an optimized mixture of ethanol and deionized (DI)
water (ratio 2 to 3) and sonicated for 3 h. The mixture
was then centrifuged and washed with ethanol and DI
water twice. The sediment was dried at 80 °C overnight
and the graphene was collected.
The as-synthesized graphene was re-dispersed into
ethanol via sonication. The graphene dispersion was then
mixed with the zinc oxide (ZnO) dispersion and further
sonicated to achieve an even mixing. The mixture was
then subjected to mechanical stirring for 4 h for the syn-
thesis of a graphene/ZnO nanocomposite. Next, the
mixture was centrifuged and washed with ethanol and DI
water twice. The graphene/ZnO nanocomposite was
collected after drying at 80 °C overnight.
The as-synthesized graphene and graphene/ZnO
nanocomposite were characterized via scanning electron
microscopy (SEM), transmission electron microscopy
(TEM) and X-ray diffractometry (XRD).
2.2 Electrochemical measurements
The electrochemical performances of the samples in
connection with cyclic voltammetry (CV) and ampero-
metry were analyzed with AUTOLAB PGSTAT302N. A
planar screen-printed carbon electrode (SPCE) was
fabricated (ScrintTechnology, Malaysia) based on the
design described by Chan et al.14 2 μL of the as-prepared
graphene/ZnO nanocomposite dispersion was dropped
onto the surface of the working SPCE and dried at
ambient temperature.
3 RESULTS AND DISCUSSIONS
3.1 Morphological analysis
A detailed morphology of the samples was studied
with SEM and TEM, as shown in Figure 1. It can be
observed in Figure 1a that graphene has a flat lamellar
structure which provides graphene with a large surface to
accommodate the hybridization with ZnO. The transpa-
rent nature of graphene shown in Figure 1b implies that
graphite was fully exfoliated into a few layer sheets.
Figure 1c shows that ZnO nanoparticles are dispersed on
the graphene nanosheets. Figure 1d also demonstrates
that the graphene sheets and ZnO nanoparticles are
closely in contact with each other, which is beneficial for
the electron transfer. Furthermore, no ZnO nanoparticles
are observed outside the graphene sheets, suggesting a
successful hybridization of the graphene/ZnO nanocom-
posite.
S. S. LOW et al.: ONE-STEP GREEN SYNTHESIS OF GRAPHENE/ZnO NANOCOMPOSITES ...
838 Materiali in tehnologije / Materials and technology 49 (2015) 5, 837–840
Figure 1: SEM and TEM images of the: a), b) as-produced graphene
and c), d) as-synthesized graphene/ZnO nanocomposite
Slika 1: SEM- in TEM-posnetka: a), b) izdelanega grafena in c), d)
sintetiziranega nanokompozita grafen-ZnO
3.2 Structural analysis
The crystalline structure of the as-synthesized
graphene/ZnO nanocomposite was corroborated with an
XRD measurement. The XRD pattern of the gra-
phene/ZnO nanocomposite is shown in Figure 2, in
comparison with graphite, graphene and ZnO as the
references. Figure 2 manifests diffraction peaks located
at 2 = (31.88°, 34.50°, 36.35°, 47.60°, 56.67°, 62.94°,
66.45°, 68.01° and 69.18°), which correspond to the
crystal planes (002), (202), (501), (640), (660), (901),
(931), (1000) and (941) of the hexagonal wurtzite
structure of ZnO (ICSD no. 98-006-5172). Additional
peaks at 2 = 26.63° and 54.71° account for the graphitic
reflection of the (002) plane, corresponding to a
d-spacing of 0.334 nm. It is also illustrated that the XRD
pattern of the graphene/ZnO nanocomposite is similar to
those of graphite, graphene and ZnO, indicating that the
synthesis method employed retained the pristine cry-
stalline structures of the materials. This is crucial as
lattice disruptions and crystalline-structure deformations
can introduce defects that adversely affect the desired
properties of the materials. The XRD data of the as-syn-
thesized nanocomposite shows the absence of impurities,
reflecting its high quality.
3.3 Electrochemical measurement
Cyclic voltammetry (CV) was recorded at room tem-
perature for bare, graphene, ZnO and graphene/ZnO
nanocomposite-modified SPCEs. The CV analyses were
performed at a scan rate of 50 mV/s in the potential
range of –0.4 V to 0.4 V. The redox reactions of H2O2
occur at very high switching potentials and therefore, for
non-enzymatic H2O2 sensors, the maximum current is
normally used for determining the sensing ability.15
Figure 3 shows that the graphene/ZnO nanocompo-
site-modified SPCE exhibits the greatest enhancement in
terms of the electrochemical performance, with the
highest peak currents (the maximum currents) being
detected for both forward and reverse scans. There is a
significant increase in the peak current by almost 2.5
times in the CV response of the composite-modified
SPCE if compared to the bare electrode. On the other
hand, the CV response of the graphene/ZnO nanocom-
posite-modified SPCE in different concentrations of
H2O2 is also illustrated in Figure 3. The maximum
current increased gradually upon increasing the concen-
tration of H2O2 from (1 to 3, 5 and 7) μM (Figure 3
S. S. LOW et al.: ONE-STEP GREEN SYNTHESIS OF GRAPHENE/ZnO NANOCOMPOSITES ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 837–840 839
Figure 4: Amperometric i–t response of graphene/ZnO nanocompo-
site-modified SPCE to the additions of different H2O2 concentrations
at an applied potential of 0.4 V. Inset shows a plot of concentration of
H2O2 versus current in the linear range.
Slika 4: Amperometri~en odgovor i–t SPCE, modificiranega z
nanokompozitom grafen-ZnO, pri dodatku razli~nih koncentracij
H2O2 pri uporabljenem potencialu 0,4 V. Vstavljeni diagram prikazuje
koncentracijo H2O2 v odvisnosti od toka v linearnem podro~ju.
Figure 2: XRD patterns of graphite, graphene, ZnO and as-synthe-
sized graphene/ZnO nanocomposite with the standard reference
pattern
Slika 2: Rentgenski posnetki grafita, grafena, ZnO in sintetiziranega
nanokompozita grafen-ZnO s standardnim referen~nim vzorcem
Figure 3: Cyclic voltammograms of: a) bare, b) ZnO, c) graphene, d)
graphene/ZnO nanocomposite-modified SPCE in PBS buffer and
cyclic voltammograms of graphene/ZnO nanocomposite-modified
SPCE in the presence of: e) 1 μM, f) 3 μM, g) 5 μM, h) 7 μM H2O2.
Inset shows the plot of the peak current against H2O2 concentration.
Slika 3: Cikli~ni voltamogrami: a) goli, b) ZnO, c) grafen, d) SPCE,
modificiran z nanokompozitom grafen-ZnO v PBS-pufru. Cikli~ni
voltamogrami modificiranega SPCE nanokompozita grafen-ZnO ob
prisotnosti: e) 1 μM, f) 3 μM, g) 5 μM, h) 7 μM H2O2. Vstavek
prikazuje diagram toka proti koncentraciji H2O2.
inset). The high electrocatalytic activity of the gra-
phene/ZnO nanocomposite can be ascribed to a large
surface-area-to-volume ratio and the synergistic effect
between graphene and ZnO.
The amperometric current – time (i – t) curve is the
most often used method to evaluate the electrocatalytic
activity of electrochemical sensors. The amperometric
response of the graphene/ZnO nanocomposite-modified
SPCE was investigated by successively adding hydrogen
peroxide (H2O2) to a continuously stirred PBS solution.
Figure 4 clearly shows that the modified SPCE exhibits
a rapid and steady response to each H2O2 addition, while
the Figure 4 inset displays the linear relationship bet-
ween the current response and H2O2 concentration in the
range of 1 mM to 15 mM with a sensitivity of 0.055
μA/mM. The linear regression equation is i/μA = 0.0548
C (mM) + 0.0547 with a correlation coefficient (R2) of
0.9859, which permits a reliable quantification of the
amount of H2O2 in a sample.
4 CONCLUSION
In summary, a facile green approach for the synthesis
of a graphene/ZnO nanocomposite at room temperature
was developed. The SEM and TEM images revealed that
flower-like ZnO nanoparticles were homogenously dis-
tributed on the surface of the graphene sheets, while the
XRD result affirms the structural integrity and purity of
the as-synthesized graphene/ZnO nanocomposite. The
graphene/ZnO nanocomposite-modified SPCE displayed
excellent electrocatalytic activity towards H2O2 due to
the synergistic effect between graphene and ZnO. The
results herein suggest promising potentials of the gra-
phene/ZnO nanocomposite for various sensing applica-
tions.
Acknowledgments
The Ministry of Science, Technology and Innovation
research grant (E-Science code: 04-02-12-SF0198), the
Early Career Research and Knowledge Transfer Scheme
Award (ECKRT: A2RHM), HIR-Chancellory UM
(UM.C/625/1/HIR/079), HIR-MOHE (UM.C/625/1/
HIR/MOHE/SC/06) and the University of Nottingham
are gratefully acknowledged. The authors also express
heartfelt gratitude to Professor Dino Isa for access to the
SAHZ pilot plant and its analytical equipment.
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S. S. LOW et al.: ONE-STEP GREEN SYNTHESIS OF GRAPHENE/ZnO NANOCOMPOSITES ...
840 Materiali in tehnologije / Materials and technology 49 (2015) 5, 837–840
T. VERLAK et al.: MATERIAL PARAMETERS FOR A NUMERICAL SIMULATION OF A COMPACTION PROCESS ...
841–844
MATERIAL PARAMETERS FOR A NUMERICAL SIMULATION OF
A COMPACTION PROCESS FOR SINTERED DOUBLE-HEIGHT
GEARS
MATERIALNI PARAMETRI ZA NUMERI^NO SIMULACIJO
POSTOPKA STISKANJA SINTRANIH DVOVI[INSKIH ZOBNIKOV
Toma` Verlak1, Marko [ori2, Sre~ko Glode`1
1University of Maribor, Faculty for Mechanical Engineering, Smetanova ulica 17, 2000 Maribor, Slovenia
2University of Maribor, Faculty of Natural Sciences and Mathematics, Koro{ka cesta 160, 2000 Maribor, Slovenia
tomaz.verlak@gmail.com
Prejem rokopisa – received: 2014-10-16; sprejem za objavo – accepted for publication: 2014-11-06
doi:10.17222/mit.2014.262
This paper presents the initial material parameters for the Ecka Alumix 231 aluminum-powdered metal required for a successful
numerical simulation of a compaction process for sintered components with a proper software package. Experimental work was
used to obtain Drucker-Prager-cap (DPC) model material parameters with the help of a Brazilian disc test and a uniaxial com-
pression test for the linear part of the DPC model in the equivalent pressure stress/deviatoric stress (p-q) plane. The specimens
used for this test were cylindrical probes (greens) that were compacted with a mechanical press and then compressed to the
failure point in the axial and radial directions.
Keywords: aluminum-based powder, compaction process, Drucker-Prager cap model, numerical simulation, Ecka Alumix 231
V ~lanku so predstavljeni za~etni materialni parametri za aluminijev prah Ecka Alumix 231, ki so potrebni za uspe{no izvedbo
numeri~ne simulacije stiskanja sintranih komponent s primerno programsko opremo. S poru{itvenimi preizkusi in pridobljenimi
eksperimentalnimi rezultati je bilo mogo~e dolo~iti potrebne parametre materiala z Drucker-Prager-jevim cap (DPC) modelom,
ki dolo~ajo linearno obmo~je DPC modela v grafu ekvivalentne tla~ne obremenitve v odvisnosti od deviatori~ne obremenitve.
Uporabljeni so bili cilindri~ni preizku{anci (zelenci), ki so bili stisnjeni z mehansko stiskalnico in kasneje obremenjeni do
poru{itve v osni in radialni smeri.
Klju~ne besede: aluminijev prah, stiskanje, Drucker-Prager-jev cap model, numeri~na simulacija, Ecka Alumix 231
1 INTRODUCTION
Sintering is a densification process of powder com-
pacts based on the atomic diffusion of powder compo-
nents. These mechanical compacts (products), especially
gears (Figure 1), are widely used in numerous fields of
mechanical engineering because they can be made into
near-net- shaped objects with minimum or no machining
after the sintering. The equipment for sintering is quite
expensive and, for that reason, the sintering process
required for manufacturing different, small, complex
components is economical for larger production series.
The process of sintering can be divided into multiple
steps, of which one of the most important is the com-
paction process. When a powder is being compacted
with mechanically or hydraulically operated presses, the
density distribution is changing.
The density magnitude and the density distribution
within a component have large impacts on the mecha-
nical properties of finished product1,2 so it is crucial that
the achieved density after the compaction process is as
close as possible to 100 % of the theoretical density of
the material and that the density distribution is uniform.
To achieve a higher density and a uniform density
distribution throughout the whole component after the
compaction process, the press operator must have a great
deal of experience to optimize the compaction parame-
ters (the movement of the press components).
When optimizing compaction parameters based on
experience, there is a great chance that damages occur to
the press tools. To avoid high costs for repairing da-
maged tools and spending the time for optimizing the
compaction parameters on a press, numerical simulations
can be applied to resolve these situations. Using a nume-
rical simulation in ABAQUS with a build-in Drucker-
Prager cap model, it is possible to eliminate the trial-and-
error process and numerically optimize press-compac-
tion parameters, thus saving time and money.2 The
Materiali in tehnologije / Materials and technology 49 (2015) 5, 841–844 841
UDK 519.61/.64:621.762.5:669.71 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(5)841(2015)
Figure 1: Double-height gear
Slika 1: Dvovi{inski zobnik
results of this kind of simulation show the local density
distribution over the entire model. By optimizing the
movement of the punches in a numerical simulation, the
simulation can be restarted again and again with diffe-
rent compaction parameters and this way the results of
the final density distributions can be compared. When
the results of a numerical simulation are satisfactory, the
compaction parameters can be transferred onto the press.
Before a numerical simulation can begin, the material
parameters for the chosen metal powder must be known.
These parameters are hard to come by or, in some cases,
they do not even exist, which is specially the case with
most aluminum powders.3 To obtain the material para-
meters for a successful simulation, suitable tests must be
done. These tests include a Brazilian disc test, a uniaxial
compression test and a triaxial compression test. With
these tests, it is possible to obtain the parameters that are
needed for drawing a Drucker-Prager-cap-model chart
and start a numerical simulation. The Drucker-Prager-
cap-model chart, shown in Figure 2, is presented with
the equivalent pressure stress/deviatoric stress (p-q)
plane. It consists of three surfaces.4,5 The first is the shear
surface, the second is the transition surface and the third
is the cap surface. This paper focuses on gathering the
Drucker-Prager-cap-model parameters for the linear part,
presented with the shear surface, and the tests required to
gather proper information include a Brazilian disc test
and a uniaxial compression test.
2 EXPERIMENTAL WORK
For the experimental work, the Ecka Alumix 231 alu-
minum powder was used. The basic powder specifi-
cations can be found in Table 1 where 
A is the apparent
density, 
G is the green density and CP is the
recommended compaction pressure. The testing can be
divided into two parts. The first part involves the
preparation of the samples and the second part involves
the testing of the samples to the failure point and the
reading of the failure-point load. The specimens used in
this experimental work ware cylindrical discs (Figure 3)
prepared with a DORST TPA 45 mechanical press (Fig-
ure 4) that is capable of a compaction force of 440 kN.
A total of 24 samples were prepared. Each compacted
specimen was measured and the values are presented in
Table 2 where D is the diameter of a specimen, L is the
height of a specimen, m is the mass and 
G is the green
density of a specimen.
Table 1: Basic properties of aluminum powder Ecka Alumix 231
Tabela 1: Osnovne lastnosti aluminijevega prahu Ecka Alumix 231
Physical characteristics Chemical compositions,w/%

A/(g/cm3) 
G/(g/cm3) CP/MPa Si Cu Mg
1.05 – 1.25 2.56 620 14–16 2.4–2.8 0.5–0.8
T. VERLAK et al.: MATERIAL PARAMETERS FOR A NUMERICAL SIMULATION OF A COMPACTION PROCESS ...
842 Materiali in tehnologije / Materials and technology 49 (2015) 5, 841–844
Figure 4: Dorst TPA 45 mechanical press
Slika 4: Mehanska stiskalnica Dorst TPA 45
Figure 3: Specimen – a cylindrical disc
Slika 3: Vzorec – cilindri~ni plo{~ek
Figure 2: Drucker-Prager-cap-model chart
Slika 2: Graf Drucker-Pragerjevega cap modela
Table 2: Average numerical values of specimens
Tabela 2: Povpre~ne numeri~ne vrednosti vzorcev
Specimens D/mm L/mm m/g 
G/(g/cm3)
1–6 24.15 16.19 18.73 2.53
7–12 24.15 16.21 18.79 2.53
13–18 24.14 16.27 17.28 2.32
19–24 24.15 16.28 17.30 2.32
For the second part of the experimental work, spe-
cimens were tested to the failure point with a WOLPERT
TZZ 100 material-testing machine (Figure 5) that is
capable of measuring compression and tension forces of
up to 100 kN. The specimens were placed between two
metal plates and then tested.6 12 cylindrical discs were
tested to the failure point in the axial direction (Figure
6a) using a testing speed of 5 mm/min and another 12 in
the radial direction (Figure 6b) using a testing speed of
2.5 mm/min.
3 RESULTS AND DISCUSSION
In order to obtain the linear part of the Drucker-
Prager-cap model (Figure 7) parameters p and q for the
compression test and Brazilian disc test must be calcu-
lated from the gathered data, where p is the hydrostatic-
pressure stress and q is the deviatoric stress. For the
compression test, parameter p can be calculated using
Equation (1) and parameter q is calculated using Equa-
tion (2). For the Brazilian disc test, parameter p can be
calculated using Equation (3), and parameter q is
calculated using Equation (4):
p =
 c
3
(1)
q =  c (2)
p =
 t
3
(3)
T. VERLAK et al.: MATERIAL PARAMETERS FOR A NUMERICAL SIMULATION OF A COMPACTION PROCESS ...
Materiali in tehnologije / Materials and technology 49 (2015) 5, 841–844 843
Figure 8: Parameters for calculation
Slika 8: Parametri za prera~un
Figure 6: a) Axial and b) radial load directions
Slika 6: a) Osna in b) radialna obremenitev
Figure 5: Wolpert TZZ 100 material-testing machine
Slika 5: Stroj za preizku{anje materialov Wolpert TZZ 100
Figure 7: Linear part of the Drucker-Prager-cap model
Slika 7: Linearni del Drucker-Pragerjevega cap modela
q = ! t (4)
where c represents the compression strength expressed
with Equation (5) and t represents the splitting tensile
strength expressed with Equation (6):
 c =
F
A
(5)
 t =
2P
LDπ
(6)
where F is the maximum load applied during the uni-
axial compression test, A is the surface area of a speci-
men, P is the maximum load applied during the
Brazilian disc test, L is the thickness of a specimen and
D is the diameter of a specimen. All the parameters
mentioned above are presented in Figure 8. Table 3
presents the average values gathered during the experi-
mental work, where F is the maximum load applied in
the axial direction and P is the maximum load applied in
the radial direction. As expected, the values of the
maximum load applied during the Brazilian disc test are
smaller because of a grater notch factor and a smaller
surface. With the help of Equations (1) to (6) and the
added trend line, the linear part of the Drucker-Prager-
cap model was plotted and it is presented in Figure 9.
Table 3: Average values of applied loads
Tabela 3: Povpre~ne vrednosti uporabljene sile
Specimens Load direction F/N P/N
1–6 axial 77145 /
7–12 radial / 8405
13–18 axial 64394 /
19–24 radial / 4275
Using Microsoft Excel and its feature "display equa-
tion", the equation for the linear part of the Drucker-Pra-
ger-cap model of the Ecka Alumix 231 aluminum-
powder metal can be displayed (Equation (7)):
y = 2.65x + 18.43 (7)
In order to successfully complete the numerical simu-
lation in ABAQUS, parameters d and  must be ob-
tained.2,7 Parameter d represents the material cohesion
and parameter  represents the angle of friction. Both of
these parameters can be obtained from Equation (7).
Parameter d is the point on the p-q plane where the line
intersects the y axis (Figure 2) and it has a value of
18.43 MPa. Parameter  represents the slope of the linear
part (Figure 2) and can be calculated using Equation (8):
 = tan–1 2.65 (8)
Using Equation (8), the angle of friction was calcula-
ted, having a value of 69 °.
4 CONCLUSIONS
Both the Brazilian disc test and the compression test
are very popular testing methods for gathering material
data for the linear part of a Drucker-Prager-cap model.2
They are both inexpensive methods because individual
specimens do not have to have exactly the same dimen-
sions and the testing equipment is widely accessible.
Dimensions are then taken into account with the use of
appropriate equations. Lower-density specimens result in
a lower applied failure load due to a lower diffusion of
dust particles and a greater notch factor. On the other
hand, higher-density specimens result in a higher applied
failure load. The present study shows that there is a
linear relationship for the specimens with different den-
sities and that the specimens with higher densities are
much more resistant to higher loads.
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T. VERLAK et al.: MATERIAL PARAMETERS FOR A NUMERICAL SIMULATION OF A COMPACTION PROCESS ...
844 Materiali in tehnologije / Materials and technology 49 (2015) 5, 841–844
Figure 9: Linear part of the Drucker-Prager-cap model for powder
Ecka Alumix 231
Slika 9: Linearni del Drucker-Prager-jevega cap modela za prah Ecka
Alumix 231