VSEBINA – CONTENTS
PREGLEDNI ^LANKI – REVIEW ARTICLES
Characterization and wear performance of CrAgN thin films deposited on Cr-V ledeburitic tool steel
Ocena lastnosti in vedenje pri obrabi tankih plasti CrAgN, nanesenih na ledeburitno orodno jeklo Cr-V
P. Jur~i, J. Bohovi~ová, M. Hudáková, P. Bílek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 159
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Synthesizing Si3N4 from a mixture of SiO2-CaO
Sinteza Si3N4 in me{anice SiO2-CaO
N. Karakuþ, H. Ö. Toplan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 171
Cast cellular metals with regular and irregular structures
Ulite kovine s pravilno in nepravilno celi~no strukturo
V. Bednáøová, P. Lichý, T. Elbel, A. Hanus . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 175
Spin-coating for optical-oxygen-sensor preparation
Uporaba spinskega nanosa pri izdelavi opti~nih senzorjev za kisik
P. Brglez, A. Holobar, A. Pivec, M. Kolar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 181
Electrochemical synthesis and characterization of poly O-aminophenol – SiO2 nanocomposite
Elektrokemijska sinteza in karakterizacija nanokompozita poli O-aminofenol – SiO2
F. Bagheralhashemi, A. Omrani, A. A. Rostami, A. Emamgholizadeh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 189
Development of low-Si aluminum casting alloys with an improved thermal conductivity
Razvoj aluminijeve livarske zlitine z majhno vsebnostjo Si in izbolj{ano toplotno prevodnostjo
J. Shin, S. Ko, K. Kim . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 195
A new method for estimating the Hurst exponent H for 3D objects
Nova metoda za ocenjevanje Hurstovega eksponenta H za 3D-objekte
M. Babi~, P. Kokol, N. Guid, P. Panjan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 203
Design of a microbial sensor using a conducting polymer of polyaniline/poly 4,4’-diaminodiphenyl sulphone-silver
nanocomposite films on a carbon paste electrode
Oblikovanje mikrobnega senzorja z uporabo prevodne polimerne polianilinske/poli 4,4’-diaminodifenil sulfonske srebrne
nanokompozitne tanke plasti na elektrodi z ogljikovo pasto
M. Sharifirad, F. Kiani, F. Koohyar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 209
A new generalized algebra for the balancing of  chemical reactions
Nova posplo{ena algebra za uravnote`enje kemijskih reakcij 
I. B. Risteski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 215
Electrodeposition and characterization of Cu-Zn alloy films obtained from a sulfate bath
Elektronanos in karakterizacija plasti zlitin Cu-Zn, nastalih iz sulfatne kopeli
A. Redjechta, K. Loucif, L. Mentar, M. Redha Khelladi, A. Beniaiche . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 221
Experimental design and artificial neural network model for turning the 50CrV4 (SAE 6150) alloy using coated
carbide/cermet cutting tools
Eksperimentalna zasnova in model umetne nevronske mre`e za stru`enje jekla 50CrV4 (SAE 6150) z uporabo orodij s karbidno ali
kermetno prevleko
M. T. Ozkan, H. B. Ulas, M. Bilgin. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 227
Estimation of the boron diffusion coefficients in FeB and Fe2B layers during the pack-boriding of a high-alloy steel
Dolo~anje koeficienta difuzije bora v plasteh FeB in Fe2B med boriranjem visoko legiranega jekla v skrinji
Z. Nait Abdellah, M. Keddam . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 237
Finite element modelling of submerged arc welding process for a symmetric T-beam
Modeliranje postopka oblo~nega varjenja pod pra{kom simetri~nega T-nosilca z metodo kon~nih elementov
O. Culha . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 243
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 48(2)157–312(2014)
MATER.
TEHNOL.
LETNIK
VOLUME 48
[TEV.
NO. 2
STR.
P. 157–312
LJUBLJANA
SLOVENIJA
MAR.–APR.
2014
Optimization of the turning parameters for the cutting forces in the Hastelloy X superalloy based on the Taguchi method
Optimiranje sil odrezavanja s Taguchijevo metodo pri stru`enju superzlitine Hastelloy X
A. Altin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 249
Thermocyclic- and static-failure criteria for single-crystal superalloys of gas-turbine blades
Termocikli~na in stati~na merila za poru{itve lopatic plinskih turbin iz monokristalnih superzlitin
L. B. Getsov, A. S. Semenov, E. A. Tikhomirova, A. I. Rybnikov . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 255
UHF RFID tags with printed antennas on recycled papers and cardboards
UHF RFID-zna~ke z natisnjenimi antenami na recikliranem papirju in kartonu
U. Kav~i~, M. Pivar, M. \oki}, D. Gregor Svetec, L. Pavlovi~, T. Muck . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 261
Topmost steel production design based on through process modelling with artificial neural networks
Projektiranje proizvodnje vrhunskih jekel na podlagi modeliranja skozi proces z umetnimi nevronskimi mre`ami
T. Kodelja, I. Gre{ovnik, R. Vertnik, B. [arler . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 269
The preparation of magnetic nanoparticles based on cobalt ferrite or magnetite
Priprava magnetnih nanodelcev na osnovi kobaltovega ferita ali magnetita
A. Stamboli}, M. Marin{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 275
Simulation of continuous casting of steel under the influence of magnetic field using the local-radial basis-function collocation
method
Simulacija kontinuirnega ulivanja jekla pod vplivom magnetnega polja na podlagi metode kolokacije z radialnimi baznimi funkcijami
K. Mramor, R. Vertnik, B. [arler . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 281
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
The applicability of sol-gel oxide films and their characterisation on a magnesium alloy
Uporabnost sol-gel oksidnih tankih plasti in njihova karakterizacija na magnezijevi zlitini
E. Altuncu, H. Alanyali . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289
Testing the tribological characteristics of nodular cast iron austempered by a conventional and an isothermal procedure
Preizku{anje tribolo{kih lastnosti nodularne litine, medfazno kaljene po konvencionalnem in izotermi~nem postopku
D. Golubovi}, P. Kova~, B. Savkovi}, D. Je{i}, M. Gostimirovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 293
Quantification of the copper phase(S) in Al-5Si-(1–4)Cu alloys using a cooling curve analysis
Uporaba analize ohlajevalne krivulje za oceno koli~ine bakrovih faz v zlitinah Al-5Si-(1-4)Cu
M. B. Djurdjevic, S. Manasijevic, Z. Odanovic, N. Dolic, R. Radisa . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 299
Investigation of induction and classical-sintering effects on powder-metal parts with the finite-element method
Primerjava vpliva indukcijskega in konvencionalnega sintranja na delce kovinskega prahu z uporabo metode kon~nih elementov
G. Akpýnar, C. Çivi, E. Atik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 305
P. JUR^I et al.: CHARACTERIZATION AND WEAR PERFORMANCE OF CrAgN THIN FILMS ...
CHARACTERIZATION AND WEAR PERFORMANCE OF
CrAgN THIN FILMS DEPOSITED ON Cr-V
LEDEBURITIC TOOL STEEL
OCENA LASTNOSTI IN VEDENJE PRI OBRABI TANKIH PLASTI
CrAgN, NANESENIH NA LEDEBURITNO ORODNO JEKLO Cr-V
Peter Jur~i, Jana Bohovi~ová, Mária Hudáková, Pavel Bílek
Department of Materials, Faculty of Materials and Technology of the STU, Trnava, Paulínská 16, 917 24 Trnava, Slovak Republic
p.jurci@seznam.cz
Prejem rokopisa – received: 2013-02-14; sprejem za objavo – accepted for publication: 2013-05-28
Specimens made from Vanadis 6 cold-work tool steel were machined, ground, heat processed using a standard regime and
finally mirror polished. After that they were layered with CrAgN. The Ag-content in the layers was chosen to be in mass
fractions w = 3 % and 15 %. Microstructural analyses revealed that the CrN grew in a columnar manner. The addition of 3 % Ag
did not influence the manner of growth of the films, but the addition of 15 % Ag made considerable changes to the film growth.
Both the hardness and the Young’s modulus were not influenced by the incorporation of 3 % Ag, but the addition of 15 % Ag
reduced them. The layers with 3 % Ag had excellent adhesion on the steel substrate. On the other hand, the addition of 15 % Ag
had a very negative impact on the coating adhesion. The films with the addition of 3 % Ag had superior tribological properties
against hard material (alumina) as well as against a soft counterpart (CuSn6 as-cast bronze), while those with 15 % Ag were too
soft and they underwent intensive wear.
Keywords: Vanadis 6 cold-work steel, PVD, chromium nitride, silver addition, tribological investigations
Vzorci, izdelani iz orodnega jekla Vanadis 6 za delo v hladnem, so bili obdelani, bru{eni, toplotno obdelani po navadnem
postopku in na koncu polirani do zrcalne povr{ine. Nato je bil nanesen CrAgN. Koli~ina v masnih dele`ih Ag v nanosu je bila
izbrana 3 % in 15 %. Analize mikrostrukture so odkrile, da raste nanos v obliki stebrastih zrn. Dodatek 3 % Ag ni vplival na
na~in rasti nanosa, dodatek 15 % Ag pa je povzro~il ob~utne razlike v rasti nanosa. Oba, trdota in Youngov modul, nista bila
odvisna od vnosa 3 % Ag, dodatek 15 % pa ju je zmanj{al. Nanosi s 3 % Ag so imeli dobro oprijemljivost na podlagi iz jekla. Po
drugi strani pa je dodatek 15 % Ag mo~no poslab{al oprijemljivost nanosa. Nanosi s 3 % Ag so imeli bolj{e tribolo{ke lastnosti
pri trdem materialu (aluminijev oksid), kot tudi pri mehkem (liti bron CuSn), medtem ko so bili tisti s 15 % Ag premehki in so
izkazovali veliko obrabo.
Klju~ne besede: jeklo za delo v hladnem Vanadis 6, PVD, kromov nitrid, dodatek srebra, tribolo{ke preiskave
1 INTRODUCTION
Thin CrN-based films have been developed over the
past three decades. They quickly gained a great deal of
interest and popularity in a variety of industrial applica-
tions due to their wear- and corrosion resistance and
good cutting properties.1–10
CrN films can be synthesized with a wide range of
chemistries, phase constitutions and properties. A wide
range of microhardness values of CrN-coatings, from
1500 HV for arc-vapour-deposited films8,11–13 up to
2450–2600 HV for those prepared by magnetron sputter-
ing techniques14–17 can be obtained. In selected cases, for
ultra-fine-grained films in particular, the microhardness
of CrN-films can reach very high values, exceeding 3000
HV, as reported by Mayrhofer et al.18 It should be noted
that the phase constitution of the films also plays a role
in their microhardness – the majority of authors
established a slightly higher hardness for the Cr2N than
for the CrN,19–23 while it was only Nouveau et al.9 and
Beger at al.24 who reported the opposite tendency. The
films prepared by high-power pulsed-magnetron sputter-
ing (HIPIMS) also have a high hardness.25 The Young’s
modulus E of CrN-based films can also be varied over a
wide range, from 188 GPa to 400 GPa, depending on the
negative substrate bias, the substrate temperature and the
substrate nature (chemistry, phase constitution, hard-
ness).14,16,17,26 The formation of high internal stresses,
mostly compressive, is one of the key problems connec-
ted with film growth on substrates. The stress intensity
depends mainly on the substrate bias voltage – the
application of low bias voltage produces internal stresses
of around 2 GPa,19,25 while the stresses can reach over
4 GPa when a negative substrate bias higher than 75 V is
used for the deposition.17,27 In any case, too high internal
stresses influence the adhesion of the films negatively.
The reduction of the stresses is possible through the
post-deposition annealing of the films, which led to
improved adhesion.13,28 The difference in the mechanical
(hardness) and physical (Young’s modulus, thermal
expansion coefficient, etc.) properties can cause serious
problems with respect to the adhesion of the coatings to
the substrate. Generally, a good adhesion can be
achieved when ledeburitic steels or cemented carbides
are used as a substrate,10,13,27,29,30 due to their high hard-
ness. The adhesion also increases as the portion of the
phase CrN decreases and Cr2N and/or (Cr) solid solution
increases.12,17,28
Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170 159
UDK 669.14.018.252:621.793 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 48(2)159(2014)
The tribological properties of CrN-based films,
however, cannot be changed over a sufficiently wide
range. External lubrication is one of the possible ways
how to improve the tribological behaviour of the films,
but commercially available lubricants (oxides, molybde-
num disulfide, graphite) exhibit considerable shortcom-
ings. Graphite and molybdenum disulfide, for instance,
undergo oxidation above 300 °C and, hence they degrade
rapidly. Furthermore, metallic oxides, on the other hand,
cannot be used at low temperatures since they exhibit an
abrasive behaviour.
This is why self-lubricating composite films have
been a subject of scientific interest in the past few years.
These films combine a hard wear-resistant matrix
(mostly formed by nitrides or carbo-nitrides) with soft
lubricious phases that provide lubricious layer at room
and elevated temperatures.
Silver is the most common addition used with tran-
sition-metal (TM) nitride thin films. It possesses a stable
chemical behaviour and can exhibit self-lubricating
properties due to its low shear strength. In addition it is
known that silver is capable of migrating to the free
surface, providing lubrication above 300 °C. Several
investigations were focused on the TM-based ceramic
films with the addition of silver, deposited on various
substrates. The findings on the effect of Ag addition can
be summarized as follows. The addition of silver can
lead to an alteration in the growth of CrN-based films, as
reported for instance by Mulligan et al.31 There is a
relatively good consensus on the lowering of the friction
coefficient at operating temperatures in the range
300–500 °C.31–37 The nature of this phenomenon is also
well known. Silver is almost completely insoluble in
CrN and forms nanoparticles in the basic CrN-com-
pound. The Ag atoms can easily migrate to the free
surface at elevated temperatures, form lubricious grains
there, and thereby reduce the friction force significantly.
On the other hand, the tribological properties of
Ag-containing films differ significantly above 500 °C,
whereas the nature of the TM nitride matrix is a key
factor influencing their improvement34 or deteriora-
tion.32,33,35 The choice of an "optimal" silver addition,
from point of view of tribological performance, into the
basic film is not in doubt – many authors evaluated films
with very high silver contents, e.g., above the mole
fraction 20 % 31–33,35–37 although it seems that a much
lower silver addition can lead to superior tribological
parameters.38,39 One can assume that silver, as a naturally
soft metal, should reduce both the hardness and the
Young’s modulus of the films. However, some expe-
rimental works established either "no effect" of silver or
a slight increase of the hardness.38,39 The information on
the effect of silver on the adhesion of films is lacking –
one of the possible reasons is that various materials with
completely different properties were used as substrates.
Kostenbauer et al.40 made very important observations.
They investigated the stress development in multilayered
TiN/Ag films with a different modulation period. The
principal finding was that the stresses in all the multi-
layers decreased up to a temperature of 380 °C, followed
by a plateau above this temperature. The authors
attributed this behaviour to the fact that further stresses
were relieved by the plastic deformation of silver
interlayers. Based on these investigations, one can also
assume that soft silver particles can also effectively
relieve the internal residual stresses in the CrN-films,
which can make a substantial contribution to the better
adhesion of CrN-films and their improved wear beha-
viour.
Various materials (silicon wafers, carbon steels,
nickel superalloys and stainless steels) have been used as
substrates for CrAgN films, but not including ledeburitic
steels. These materials, however, belong to the most
important tool steels, which are used in many industrial
operations, due to their good combination of structure
and mechanical properties. Moreover, they can be
changed using variations of heat-treatment parameters
across a wide range.41–43 The main challenge to perform
the experimental works was thus to minimize the gap in
knowledge, e.g., to evaluate what happens when P/M
ledeburitic steels are coated with CrAgN. The current
paper attempts to put together the experimental results
on the development of adaptive nanocomposite CrAgN
coatings on the Vanadis 6 Cr-V ledeburitic tool steel and
to make a comprehensive report on the basic coating
characteristics, like microstructure, hardness, Young’s
modulus, wear resistance, friction coefficient, as a
function of the silver content and deposition temperature,
with a small look into the near future at the end of the
paper.
2 EXPERIMENTAL
2.1 Material and processing
The experimental material was the powder metallur-
gical ledeburitic steel Vanadis 6 with nominally mass
fractions: w(C) = 2.1 %, w(Si) 1.0 %, w(Mn) = 0.4 %,
w(Cr) = 6.8 %, w(Mo) = 1.5 %, w(V) = 5.4 % and Fe as
balance. Two types of samples were made from the expe-
rimental material. The first ones were plates of 40 mm ×
20 mm × 5 mm intended for the wear testing. The second
samples were un-notched Charpy impact-test specimens
(55 mm × 10 mm × 10 mm). After rough machining to
the semi-final dimensions, the samples were subjected to
a standard heat-treatment procedure consisting of the
following steps: vacuum austenitizing up to the final
temperature of 1050 °C, followed by nitrogen gas
quenching (5 bar) and twice tempering. Each tempering
cycle was conducted at 530 °C for 2 h. The resulting
as-tempered hardness of the material was 724 HV (10).
After the heat treatment, all the samples were fine
ground and polished with a diamond suspension to a
mirror finish.
The CrN- and CrAgN coatings were deposited in a
magnetron sputter deposition system, in a pulse regime
P. JUR^I et al.: CHARACTERIZATION AND WEAR PERFORMANCE OF CrAgN THIN FILMS ...
160 Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170
with a frequency of 40 kHz. It should be noted that no
adhesion inter-layer (like pure CrN) was deposited prior
to the formation of either the CrN or CrAgN films, in
order to highlight the differences in the adhesion of the
films themselves. Two targets, positioned opposite to
each other, were used. For the deposition of CrN, two
targets from pure chromium (99.9 % Cr) were used. The
target output power was adjusted to 2.9 kW for each
cathode. For the deposition of the silver-containing
films, one silver cathode (99.98 %) was inserted into the
processing chamber instead of one chromium target. In
these trials, the cathode output power was 5.8 kW on the
chromium cathode. On the silver cathode, the output
powers were 0.1 kW and 0.45 kW in order to produce
the silver contents in the coating of 3 % and 15 % (mole
fractions 1.3 % and 6 %), respectively. Two deposition
temperatures were used. The first one was 250 °C. To
achieve that, the samples were heated using resistive
heaters during the sputter-cleaning step. Afterwards, the
substrate temperature of 250 °C was kept constant by ion
bombardment only. The second deposition temperature
was 500 °C. It was achieved using resistive heaters
placed on the internal walls of the processing chamber.
The processes were carried out in a low-pressure atmo-
sphere (0.15 mbar), containing pure nitrogen and argon
(both of 99.999 % of purity), in a ratio of 1 : 4.5.
The specimens were ultrasonically degreased in
acetone and loaded into the processing chamber. They
were placed between the targets on rotating holders, with
a rotation speed of 3 r/min. Just prior to the deposition,
the substrates were sputter cleaned in an argon low-pres-
sure atmosphere for 15 min. The substrate temperature
was 250 °C for the cleaning. A negative substrate bias of
200 V was used for the sputter cleaning and that of 100
V for the deposition. The total deposition time was 6 h.
2.2 Investigation methods
The microstructure of the substrate material was
documented using a light microscope ZEISS NEOPHOT
32 and a field-emission scanning electron microscope
(SEM) JEOL JSM-7600F operated at an accelerating
voltage of 15 kV. Metallographic samples were prepared
using a standard preparation technique (rough and fine
grinding, polishing by diamond suspension) and finally
etched with Villela-Bain reagent.
The microstructure of the counterface materials was
recorded using a light microscope ZEISS NEOPHOT 32
after a standard metallographic preparation.
Microstructural analyses of coatings were completed
on the fracture surfaces of coated Charpy impact speci-
mens. The material was immersed into liquid nitrogen
for 20 min. and then broken down. The fracture surfaces
were cleaned ultrasonically with acetone before observa-
tion, in order to remove any possible contamination of
the material. A scanning electron microscope JEOL
JSM-7600F operating with the same parameters was
used for the analyses.
The nanohardness and the Young’s modulus (E) of
the coatings were determined using a nano-indenter
under a normal load of 20 mN, for a NanoTest (Micro
Materials Ltd) nanohardness tester equipped with a
Berkovich indenter. Ten indentations were made for each
coating applied and the mean value and the standard
deviation were calculated from the measured values. The
penetration depth (loading) of the indenter was chosen so
as to not exceed one tenth of the total coating thickness,
in order to avoid the influence of the substrate on the
measured nanohardness.
The adhesion of the coatings on the substrate was
evaluated using a CSM Revetest scratch-tester. The
scratches were made under progressively increasing
loads from 1 N to 100 N, with a loading rate of 50
N/min. A standard Rockwell diamond indenter with a tip
radius of 200 μm was used. Five measurements were
made on each specimen and the mean value of the
adhesion, represented by the Lc1 and Lc2 critical loads,
respectively, was calculated. The critical loads were
determined by the recording of an acoustic emission
signal as well as by viewing the scratches on the light
micrographs. The Lc1 critical load corresponded to the
occurrence of the first inhomogeneities in the coating
and the Lc2 critical load was determined as the load when
50 % of the coating was removed from the substrate,
unless otherwise discussed in the text.
The tribological properties of the coatings were
measured using the CSM Pin-on-disc tribometer at
ambient and elevated temperatures up to 500 °C. Balls of
6 mm in diameter, made from sintered alumina and
CuSn6 bronze (as-cast structure, hardness of 149 HV
(10)) were used for testing at all the temperatures. The
balls made from heat-processed 100Cr6 ball-bearing
steel (hardness of 700 HV (10)) were used, but for the
testing at ambient temperature only. The experiments
were conducted in laboratory air, at a relative humidity
of 40–50 %. No external lubricant was added during the
measurements. The normal loading used for the inve-
stigations was 1 N. For each measurement, the number
of cycles was 5100, e.g., the total sliding distance was
100 m at the sliding radius of 5 mm. After the testing,
the wear-track widths were measured on a light micro-
scope ZEISS NEOPHOT 32 at a magnification of
50-times. Ten measurements were made on each track
and the mean value, which was used for further evalu-
ation, was calculated. The volume loss of the coated
samples was calculated from the width of the wear tracks
using the formula:44
V R
r
d
r
d
r dl = ⎛
⎝
⎜ ⎞
⎠
⎟
⎡
⎣
⎢
⎢
⎢
⎤
⎦
⎥
⎥
⎥
− −2
2
4
4
2
2 2π
sin
where R is the wear-scar radius, d is the mean value of
the wear-track width, and r is the radius of the ball
counterface.
P. JUR^I et al.: CHARACTERIZATION AND WEAR PERFORMANCE OF CrAgN THIN FILMS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170 161
The wear rate was then derived from the volume-loss
calculation, using the formula:44
W
V
l F
=
×
l
n
where W is the wear rate, l is the sliding distance (m),
Fn is the applied normal load and Vl is the volume loss.
3 RESULTS AND DISCUSSION
3.1 Substrate characterization
The microstructure of the substrate material after the
heat treatment is shown in Figures 1 and 2, respectively.
The light micrograph (Figure 1) shows that the material
consists of a matrix, formed with tempered martensite
and fine carbides, uniformly distributed throughout the
matrix. The carbides are of two types (Figure 2). The
MC-phase (medium-sized particles) mostly forms the
eutectic part of the carbides. The second carbide type is
the M7C3. The M7C3-phase underwent dissolution in the
austenite during the heat processing, being responsible
for the saturation of the austenite with the carbon and
alloying elements.45 The other part of the M7C3-carbides
(designated as large secondary carbides (LSCs) and
small secondary carbides (SSCs), and an almost
complete amount of MC phase remained undissolved.
After the heat treatment, the average hardness of the
material was 724 HV (10).
3.2 Microstructure of the films
Figure 3 shows cross-sectional secondary-electron
micrographs from all the developed films. The thickness
of the CrN film without any silver addition was 4.3 μm
(Figure 3a). This corresponds to a growth rate of 720
nm/h. The addition of 3 % of Ag did not change the
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162 Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170
Figure 3: SEM micrographs showing the microstructure of developed films: a) CrN, deposition at 250 °C, b) CrAg3N, deposition temperature of
250 °C, c) CrAg3N, deposition temperature of 500 °C and d) CrAg15N, deposition temperature of 500 °C47,48
Slika 3: SEM-posnetki mikrostrukture razvoja plasti: a) nanos CrN pri 250 °C, b) CrAg3N, temperatura nana{anja 250 °C, c) CrAg3N, tempe-
ratura nana{anja 500 °C in d) CrAg15N, temperatura nana{anja 500 °C47,48
Figure 2: SEM micrograph showing detailed microstructure of PM
ledeburitic steel Vanadis 6 substrate in the as-quenched and tempered
state
Slika 2: SEM-posnetek detajla mikrostrukture kaljene in popu{~ene
podlage iz PM ledeburitnega jekla Vanadis 6
Figure 1: Light micrograph showing the microstructure of the PM
ledeburitic steel Vanadis 6 substrate in the as-quenched and tempered
state
Slika 1: Posnetek mikrostrukture kaljene in popu{~ene podlage iz PM
ledeburitnega jekla Vanadis 6
thickness (and the growth rate) of the films (Figures 3b
and 3c). On the other hand, the addition of 15 % Ag
accelerated the growth rate of the film to 1050 nm/h and,
as a result, this film had a thickness of 6.3 μm (Figure
3d). It should be noted that these results are inconsistent
with other literature data. Yao et al.,46 for instance,
reported a decreased film thickness with a Ag addition.
However, they have used completely different experi-
mental setup for the coating deposition, which makes the
results only roughly comparable. Our results, on the
other hand, indicate that small addition of silver does not
influence the growth rate of the films, which was con-
firmed recently.47,48
The pure CrN-based film, formed at 250 °C, grew in
a columnar manner with clearly visible individual crys-
tallites (Figure 3a). This type of layer growth is typical
for magnetron sputtered CrN-films formed over a wide
range of deposition parameters, as reported previously.49
The addition of 3 % Ag into the CrN, formed at the same
temperature, did not change the growth mechanism of
the layer significantly (Figure 3b). The temperature effect
on the layer growth for the films with 3 % Ag addition is
visible on the micrograph in Figure 3c. It is clearly
visible that the higher deposition temperature does not
influence the growth manner. Figure 3d shows the
microstructure of the film with 15 % Ag addition. The
secondary-electron-mode detection yields clearly visible
contrast between high atomic Ag (bright) and the
surrounding CrN base. It is also clear that no individual
crystals of CrN were formed. Therefore, it is clear that a
small Ag addition does not change the growth manner of
the film, while a higher Ag content incorporated into the
CrN matrix has a considerable impact on that.
Figure 4 brings representative SEM micrographs of
the deposited films and corresponding EDS maps of
silver. These pictures indicate that the surface micro-
structure is strongly influenced by introducing silver into
the basic CrN-film. The surface of pure CrN (Figure 4a)
exhibits a clearly visible, non-uniform structure, made
up of two types of features. The first type of feature is
the semi-equiaxed grains (SEG) with a size of 0.4–0.7
μm. This feature makes up around 90 % the surface and
can be referred to as a "matrix". In this "matrix", several
formations that can be described as a cauliflower-like
(CFL) structure are embedded. Similar surface structure
of chromium nitride films has already been reported and
discussed by Zhao et al.50 They established that the com-
bination of "equiaxially grained" (or facetted) and CFL
structures is typical for two-phase (CrN and Cr2N) films.
The addition of 3 % Ag into the chromium nitride
basic film causes a significant refinement in the grain
size of the individual crystals (Figure 4b). No individual
silver grains but inhomogeneities in the chemical com-
position throughout the micrograph became visible
(Figure 4c).
The SEM micrograph in Figure 4d and the corres-
ponding EDS mapping of Ag (Figure 4e), show that
silver forms agglomerates (appearing in bright contrast
due to their higher secondary-electron yield versus CrN)
on the surface at a concentration of 15 %. The size of the
silver agglomerates is well below 1 μm.
3.3 Mechanical properties of the films
The microhardness of pure CrN was (16.79 ± 1.49)
GPa (Table 1). The microhardness of films containing
3 % Ag was only very slightly lower than that of the film
that does not contain silver. Furthermore, the microhard-
ness of these films was almost the same, e.g., the depo-
sition temperature plays only a very minor role with
respect to the coating hardness. The addition of 15 %
Ag, on the contrary, led to a substantial hardness reduc-
tion, i.e., (11.43 ± 0.61) GPa. This may be explained by
P. JUR^I et al.: CHARACTERIZATION AND WEAR PERFORMANCE OF CrAgN THIN FILMS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170 163
Figure 4: Plan-view SEM micrographs showing the surface micro-
structure of deposited films: a) CrN, deposition at 250 °C, b) CrAg3N,
deposition temperature of 250 °C, c) corresponding EDS-map of
silver, d) CrAg15N, deposition temperature of 500 °C and e) corres-
ponding EDS-map of silver
Slika 4: SEM-posnetki povr{ine, ki prikazuje mikrostrukturo nanes-
ene plasti: a) CrN, nane{en pri 250 °C, b) CrAg3N, temperatura
nana{anja 250 °C, c) EDS-razporeditev srebra, d) CrAg15N, tempera-
tura nana{anja 500 °C in e) EDS-razporeditev srebra
Table 1: Mechanical properties of investigated films. The circled
values () were obtained in a previous study47
Tabela 1: Mehanske lastnosti preiskovanih nanosov. Vrednosti,
ozna~ene z (), so iz vira47
Coating/deposition
temperature Hardness (GPa)
Young’s modulus
E/GPa
CrN/250 °C  16.79 ± 1.49 244 ± 15
CrAg3N/250 °C  15.97 ± 1.44 241 ± 9
CrAg3N/500 °C 16.13 ± 1.83 246 ± 17
CrAg15N/500 °C 11.43 ± 0.61 204 ± 6
the fact that silver is a very soft metal and its agglome-
rates embedded in the CrN matrix cause softening of the
film. These observations are very consistent with other
experimental works. Yao et al.46 reported only a very
slight Knoop hardness decrease for magnetron sputtered
nanocomposite coatings with very small silver additions.
Mulligan et al.31 established a much more remarkable
hardness decrease, but they added the mole fraction of
Ag 22 % (e.g., a much larger amount than that used in
this study).
The Young’s modulus, E, of the CrN and CrAg3N
films was of about 240 GPa (Table 1). The E value
ranges, in addition, overlap considerably. The addition of
15 % of silver to the basic film, on the other hand, tends
towards a decrease of the Young’s modulus. The
information on the impact of the silver addition on the
Young’s modulus is lacking in the literature. Only
Aouadi and his co-workers38 reported a decreased
Young’s modulus with an increased amount of silver in
magnetron-sputtered YSZ-based films. The fact that the
Young’s modulus reduces with the addition of silver can
be considered as natural, since silver had a much lower E
(around 79 GPa) than chromium nitride (over 200 GPa in
most cases14,16,17). However, at very small silver addi-
tions, there is almost no impact of silver on the Young’s
modulus, as indicated in Table 1. It seems that there is a
threshold below which no impact of the silver on the
Young’s modulus can be expected and, above which the
addition of silver leads to a substantial lowering of the E.
One can believe that the effect of the size and the
distribution of silver particles can also play a role in the
mechanical behaviour of the coatings. However, there is
no relevant information on the impact of these structural
parameters on the mechanical properties of the films yet.
These investigations can be considered as a challenging
factor for further investigations.
3.4 Adhesion of the films
After the scratch-test, the failure of the pure chro-
mium nitride film begins with semi-circular tensile
cracking (Figure 5a). The first cracks were observed at a
normal load of around 24 N (Lc1). The "total" failure of
the chromium nitride film is shown in Figure 5b. It is
manifested by many parallel cracks visible in the scratch,
where about 50 % of the coating is removed from the
substrate. The typical load range when this phenomenon
occurred was 40–45 N.
The slightly softer CrAg3N film deposited at 250 °C
also failed due to the presence of semi-circular tensile
cracks. However, the distance between the first cracks is
larger than that in the pure chromium nitride and some of
the cracks stopped their propagation through the scratch
(Figure 5c). The critical load at which these phenomena
first occurred was around 23 N (Table 2). Figure 5d
shows the total failure of the CrAgN film grown at 250
°C. It is evident that some of the parallel cracks stopped
their propagation through the film, which suggests that
the coating can store a larger amount of plastic defor-
mation energy preceding the failure. This assumption is
supported by the fact that the "total" failure of the film
was detected at a load higher than that of pure chromium
nitride (Table 2).47
Table 2: Critical loads for a defined degree of coatings failure47
Tabela 2: Kriti~ne obremenitve za opredeljeno stopnjo po{kodbe
nanosa47
Coating/deposition
temperature Lc1/N Lc2/N
CrN/250 °C 24.5 ± 1.7 42.7 ± 4.4
CrAg3N/250 °C 23.4 ± 5.9 52.2 ± 5.9
CrAg3N/500 °C 46.9 ± 8.1 82.6 ± 8.4
CrAg15N/500 °C 6.4 ± 0.6 44.1 ± 6.3
For the film with 3 % Ag addition, grown at a sub-
strate temperature of 500 °C, the first indication of
P. JUR^I et al.: CHARACTERIZATION AND WEAR PERFORMANCE OF CrAgN THIN FILMS ...
164 Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170
Figure 5: Light micrographs showing the failures after scratch testing:
a) CrN, deposition temperature of 250 °C, Lc1, b) Lc2, c) CrAg3N,
deposition temperature of 250 °C, Lc1, d) Lc2, e) CrAg3N, deposition
temperature of 500 °C, Lc1, f) Lc2, g) CrAg15N, deposition tempe-
rature of 500 °C, Lc1, h) Lc2 47,48
Slika 5: Svetlobni posnetki napak, nastalih pri preizkusu razenja:
a) CrN, temperatura nana{anja 250 °C, Lc1, b) Lc2, c) CrAg3N, tem-
peratura nana{anja 250 °C, Lc1, d) Lc2, e) CrAg3N, temperatura nana-
{anja 500 °C, Lc1, f) Lc2, g) CrAg15N, temperatura nana{anja 500 °C,
Lc1, h) Lc2 47,48
coating damage occurred at an average loading of around
47 N (Lc1). Coating damage begins with the appearance
of semi-circular tensile cracks, e.g., it looks to be similar
to that of the coating deposited at a lower temperature
(Figure 5e). The "total" failure of the film is in Figure
5f. It is typical, with the occurrence of many parallel
micro-cracks inside the scratch track. Some of them
stopped their propagation during the testing. Also, it is
shown that part of the track (the left-hand side of the
micrograph) does not exhibit visible cracks. Here, no
typical "coating damage" was observed (50 % of the
coating removed). The typical load when this symptom
occurred ranged between 74 N and 88 N (Lc2).
The beginning of the failure of the film containing
15 % Ag cannot easily be found. Figure 5g shows a
scratch track at a relatively low loading range (around
6.4 N), where the critical load Lc1 was determined.
However, neither tensile cracking nor spallation of the
film has been observed at low loading. What was deter-
mined was only that the film underwent a local plastic
deformation with clearly visible, semi-circular deforma-
tion zones (arrow designated). These zones are widely
spaced, which suggests that the film is capable of storing
a relatively large amount of plastic energy before failing
cohesively. Typical symptoms for the "total" failure of
the coating have not been detected, in a similar way to
the film with 3 % Ag (Figure 5f). The scratch track
contained many parallel cracks and microcracks when
subjected to higher loads (Figure 5h). Moreover, the first
symptoms of chipping were detected as being adjacent to
the scratch track at a normal load of 44.1 N.
The fact that a high silver content tends to worse
adhesion of the film is consistent with Yao’s findings,46
where a very slight adhesion decrease with increasing
silver content has been reported. However, the experi-
mental setup used was different in this work, i.e., an
M2-type high-speed steel with unknown hardness and
heat treatment state was used as a substrate instead of
Vanadis 6, another deposition system was applied for the
coating growth and, finally, no details of the scratch
testing were reported. All these facts make the results
almost incomparable. What can be suggested is that the
high silver content makes the coating too soft and very
sensitive to the failure at higher loading.
Moreover, the addition of 3 % silver increased the
adhesion of the films in our work, particularly when a
temperature of 500 °C was used for the deposition. Here,
Kostenbauer’s findings,40 that the intrinsic stresses in
silver-containing multilayer coatings are relieved above
380 °C, can give a correct explanation. It is known that
the deposition of chromium nitride films, which do not
contain silver, leads to the formation of high intrinsic
stresses, exceeding 4 GPa.17,19,25,27 They can be lowered
by several methods. One of them is a so-called "post-
deposition annealing”, which can result in the increased
adhesion performance of the films.13 The second suitable
method is the incorporation of a soft and insoluble phase
(pure silver for instance) into the chromium nitride. It
can be assumed that such a phase would be capable of
relaxing the stresses during the deposition of the films.
Based on this assumption, a distinctively higher adhesion
strength of the films that contain 3 % Ag seems to be
logical.
3.5 Tribological investigations
Figure 6 gives an overview of the friction coeffi-
cients μ resulting from testing at room temperature for
all the used counterparts. The pure CrN film has a μ =
0.378 when tested against alumina. The CrAg3N films
formed at 250 °C and 500 °C had average friction
coefficients of 0.389 and 0.373, respectively. The lowest
average μ was recorded for the film containing 15 % Ag,
i.e., 0.365. Therefore, one can conclude that almost no
positive effect of the silver addition can be found when
an alumina ball is used.
Generally, the testing against 100 Cr6 ball-bearing
steel gave a slightly higher μ compared to the alumina.
The pure CrN film had average of μ = 0.425. The silver
containing-films had a slightly lowered friction coeffi-
cient, e.g., around 0.4, but it should be noted that the
friction-coefficient value ranges overlap. Thus, one
would conclude that no positive impact of the silver on
the friction coefficient of CrN against 100 Cr6 steel has
been found, in a similar way to the case of alumina ball
counterpart.
Testing against a CuSn6 bronze ball gave a lower
friction coefficient. The pure CrN film had a μ = 0.332.
A silver addition of 3 % tended to lower the friction
coefficient and the lowering of the friction coefficient
became even more significant for the composite Ag-con-
taining films grown at a temperature of 500 °C. The
impact of silver incorporation into the basic CrN com-
pound can be thus considered as slightly positive, when
tested against bronze at room temperature.
Figure 7 shows an overview of the friction coeffi-
cients μ recorded while testing against alumina at a room
and elevated temperatures. The normal load applied was
P. JUR^I et al.: CHARACTERIZATION AND WEAR PERFORMANCE OF CrAgN THIN FILMS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170 165
Figure 6: Average friction coefficient of investigated films against
various counterpart materials, testing at room temperature, normal
load applied of 1 N
Slika 6: Povpre~ni koeficient trenja za preiskovane nanose pri raz-
li~nih materialih v paru, preizku{eno pri sobni temperaturi, normalna
obremenitev 1 N
1 N. As stated above, almost no differences in μ were
found when tested at room temperature. Testing at 300
°C, however, yields to a different behaviour of the coat-
ings. The friction coefficients for the pure CrN, CrAg3N
grown at 250 °C, CrAg3N grown at 500 °C and
CrAg15N formed at 500 °C were 0.357, 0.304, 0.238
and 0.110, respectively. Higher testing temperatures
lowered the difference in the μ for different coatings, for
instance, the friction coefficients recorded by the testing
at 400 °C were 0.256, 0.194, 0.160 and 0.139 for the
pure CrN, CrAg3N formed at 250 °C, CrAg3N formed at
500 °C and CrAg15N formed at 500 °C, respectively.
The measurement at 500 °C gave rather similar results.
Table 3 shows the wear rates calculated from the
widths of the wear scares produced by the sliding of an
alumina counterpart on the samples’ surfaces. At room
temperature the beneficial effect of the silver addition is
evident only in the case of the CrAg3N film grown at
500 °C. Compared to pure CrN, the wear rate is lowered
by two times. For the film containing 3 % Ag grown at
250 °C, a slight worsening of the wear rate was recorded
and, for the film containing 15 % Ag the wear rate was
higher by an order of magnitude.
The testing at elevated temperature led to a dramatic
increase of the wear rate for the pure CrN film. The film
with a 3 % Ag addition grown at 250 °C behaved in a
very similar way from the qualitative point of view. The
wear rate was minimal after the testing at an ambient
temperature, which was followed by a steep increase at
300 °C and a decrease at 400 °C and 500 °C, respec-
tively. However, it is clearly evident that the wear rates
after the testing at these temperatures are considerably
lower. This makes a distinct difference compared to the
wear behaviour of the pure CrN film. The film with the
same silver content, but grown at 500 °C, had lower wear
rates than that grown at 250 °C at room temperature and
300 °C, respectively. However, the testing at higher tem-
peratures brought a higher wear rate. The film that
contains 15 % Ag had the lowest wear rate when tested
at 300 °C. However, the wear rate of the film increased
as the testing temperature increased, and the increase of
the wear rate was found to be greater than those of other
investigated films.
The results of wear-rate measurements at room tem-
perature are very consistent with the obtained values of
the friction coefficients – the wear rate decreased with a
decrease of the μ.
In order to explain the temperature behaviour of both
coatings, it should be noted that there was no lubricant
added to the experimental setup. The fact that the wear
rate was lower or higher can thus be attributed only to
the self-lubricating effect of the silver. This effect is,
however, prevalent only above the testing temperature of
400 °C, when the Ag-atoms are capable of being trans-
ported to the surface.
Below this temperature, but above the ambient tem-
perature, only the effect of the softening of the coatings
can be expected. It was reflected by a much higher wear
rate at 300 °C than that at ambient temperature. Finally,
it should also be noted that the softening of alumina
could take place during the testing.51 The behaviour of
the film that contains 15 % Ag can be characterized as a
special case. As we determined, the incorporation of
15 % Ag makes the film soft. This gives a natural
explanation for the high wear rate measured at ambient
temperature. At a temperature of 300 °C, the friction
coefficient of the coating became extremely low and
thereby the wear rate was lowered, also. Above this
temperature, however, a softening of the coating can be
expected, as reported for instance by Kostenbauer et al.40
which gives a natural explanation for the high wear rate
measured at 400 °C and 500 °C, respectively.
Figure 8 demonstrates the wear scars obtained by
testing the CrAg3N and CrAg15N films developed at
500 °C, at ambient temperature. The testing gave very
narrow tracks in the cases of CrN and CrAgN, respec-
tively, see the representative micrograph in Figure 8a.
The track developed on the CrAg15N film looks rather
wider (Figure 8b), as a result of the lower hardness of
P. JUR^I et al.: CHARACTERIZATION AND WEAR PERFORMANCE OF CrAgN THIN FILMS ...
166 Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170
Figure 7: Average friction coefficient of investigated films against
alumina at room (RT) and elevated temperatures, normal load applied
of 1 N
Slika 7: Povpre~ni koeficient trenja za preiskovane nanose pri alumi-
nijevem oksidu pri sobni temperaturi (RT) in povi{anih temperaturah,
normalna obremenitev 1 N
Table 3: Wear rate (mm3/(N m)) at ambient and elevated temperatures, alumina used as a counterpart
Tabela 3: Hitrost obrabe (mm3/(N m)) pri sobni in povi{anih temperaturah; aluminijev oksid uporabljen kot par
Testing temperature
coating (°C) CrN CrAg3N/ 250 °C CrAg3N/ 500 °C CrAg15N/ 500 °C
Room temperature 6.947 × 10–13 7.399 × 10–13 3.65 × 10–13 1.031 × 10–12
300 2.926 × 10–11 2.894 × 10–11 8.405 × 10–12 7.053 × 10–12
400 1.162 × 10–11 4.329 × 10–12 9.927 × 10–12 1.925 × 10–11
500 1.927× 10–11 6.208 × 10–12 1.522 × 10–11 4.802 × 10–11
the film. These results correspond well with the
measured wear rates. The wear rates of the films formed
with pure CrN and CrAg3N, respectively, were very low,
while that of the CrAg15N film was considerably higher
(Table 3).
Figure 9 shows representative optical images of the
wear scars developed while testing the CrAg3N and
CrAg15N films, respectively, against the alumina
counterpart at various temperatures. Compared to the
wear tracks obtained at ambient temperature, there is
significant broadening visible. This can be due to the fact
that the elevated temperature tends towards softening of
the films. It should be noted that the softening of alumina
could also be expected. However, the softening of the
films probably became more remarkable.
The testing at 300 °C (Figures 9a and 9d), gave wear
tracks of similar width, e.g., 0.267 mm and 0.251 mm for
the CrAg3N and CrAg15N films, respectively. This can
be considered rather surprising because the CrAg15N
film is much softer than the CrAg3N and, further soften-
ing can be expected at higher temperature. However, the
friction coefficient was measured to be much lower for
the CrAg15N; this can be referred to the higher silver
content in the CrN and its better capability to facilitate
lubrication at higher temperatures, in good agreement
with previous investigations.32,33,37
The testing at 400 °C (Figures 9b and 9e) induced
significant broadening of the wear tracks. This is high-
lighted for the CrAg15N film much more than that for
CrAg3N, compare Figures 9b and 9e. Here, the effect of
softening of CrAg15N became prevalent and the
self-lubrication by Ag particles was insufficient to
compensate for the softening of the film. It is worth
noting that the opinions on the optimal silver content
from the point of view of the self-lubricating effect
facilitation are divergent. Hu et al.,33 for instance, have
reported that the minimal Ag amount in the coating was
the mole fraction x = 24 % to ensure lubrication at higher
temperatures. Mulligan et al.52 have established that the
optimal amount of silver incorporated into CrN is x = 22
%. This is quite enough for the formation of a lubricious
film inside the wear tracks and for the achievement of an
excellent tribological performance. In the current
experiment, much lower silver additions were used. It is
thus logical that this is the softening of the CrN through
the Ag, which plays a dominant role in the tribological
behaviour of the films.
The testing at 500 °C highlighted the differences
between the tribological behaviour of the films contain-
ing 3 % and 15 % Ag, respectively (Figures 9c and 9f).
The width of the wear track on the CrAg3N film
P. JUR^I et al.: CHARACTERIZATION AND WEAR PERFORMANCE OF CrAgN THIN FILMS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170 167
Figure 8: Light micrographs showing the wear tracks of: a) CrAg3N
film grown at 500 °C, b) CrAg15N film grown at 500 °C, after sliding
at room temperature, alumina used as a counterpart
Slika 8: Svetlobni posnetek prikazuje sledi obrabe: a) CrAg3N-nanos,
nanesen pri 500 °C, b) CrAg15N-nanos, nanesen pri 500 °C, po drse-
nju pri sobni temperaturi, za par je bil uporabljen aluminijev oksid
Figure 9: Plan-view optical images showing the wear tracks of the
CrAg3N film grown at 500 °C: a) testing temperature of 300 °C, b)
testing temperature of 400 °C, c) testing temperature of 500 °C and
CrAg15N film grown at 500 °C, d) testing temperature of 300 °C, e)
testing temperature of 400 °C, f) testing temperature of 500 °C
Slika 9: Svetlobni posnetki sledi obrabe na CrAg3N-nanosu, nane-
senem pri 500 °C: a) temperatura preizkusa 300 °C, b) temperatura
preizkusa 400 °C, c) temperatura preizkusa 500 °C in CrAg15N-
nanos, nanesen pri 500 °C, d) temperatura preizkusa 300 °C, e) tem-
peratura preizkusa 400 °C, f) temperatura preizkusa 500 °C
increased to 0.324 mm. Here, no indications of total
failure of the film were observed. The width of the films
that contain 15 % Ag was 0.477 mm. There are indica-
tions of total film failure clearly visible – the free sub-
strate surface is shown inside the wear scars.
Figure 10 shows representative plan-view micro-
graphs of the wear tracks obtained by testing of
CrAg3N-film, grown at 500 °C, against the CuSn6 ball.
The testing at ambient temperature gave wear tracks with
a width of 0.89 mm (Figure 10a). There are two typical
areas inside the wear track. The first one can be clas-
sified as almost unaffected by the sliding (1). The second
part of the wear track exhibits clearly visible symptoms
of scraping (2). However, no evidence of the counter-
part’s adhesion was recorded.
The testing at a temperature of 400 °C produced wear
tracks with a width of 1.02 mm (Figure 10b). Here, two
typical areas inside the track can also be found. The first
one is, similar to the case of the testing at a room tem-
perature, almost unaffected by the sliding (3). The domi-
nant part of the wear track, however, shows indications
of considerable counterpart material transfer (4). The
counterpart’s material exhibits evidence of surface oxida-
tion, because of the elevated temperature used for the
testing.
To explain the behaviour of the sliding couple
CrAg3N deposited at 500 °C vs. CuSn6 bronze, various
aspects should be considered. First of all it should be
noted that the behaviour of the sliding couple is deter-
mined by the fact that the CuSn6 is soft in nature (149
HV 10) and it has a low shear strength, so that it can be
easily smeared on the surface. This tendency is high-
lighted when a higher testing temperature is used. More-
over, the bronze does not contain any hard phases.
Figure 11 shows the microstructure of the CuSn6 ball.
The bright particles are the -phase formations with an
average microhardness of 156 HV (0.1) and the dark
places are formed by the eutectoid  + , having an
average microhardness of 196 HV (0.1). This is why no
abrasive character of the interaction sample/counterpart
has been detected. Moreover, the softness of the CuSn6
has made it impossible to determine the wear rate using
the method.44 The measurement of wear profiles using
profilometers was not possible, also. Therefore it can
only be concluded that the CrAg3N film exhibits good
anti-sticking properties at low temperature, but that these
properties are worsened at elevated temperatures.
4 CONCLUSIONS
Investigations of magnetron-sputtered CrN films with
various Ag additions have brought the following
findings:
The pure chromium nitride film grew in a typical
columnar manner with clearly visible individual crystals.
Its surface structure is a mixture of semi-equiaxed grains
and cauliflower-like formations.
The addition of 3 % Ag tends mainly towards a
refinement in the scale of the microstructural features,
while the incorporation of 15 % Ag induced considerable
changes in the growth manner, whereas no individual
crystals of the CrN are more visible and nano-sized
silver agglomerates are formed.
The pure CrN film as well as those with a 3 % Ag
addition grew at a deposition rate of 720 nm/h, while the
growth rate of the film with 15 % Ag was accelerated to
1050 nm/h. It is clear that the deposition temperature did
not affect both the deposition rate and the final coating
P. JUR^I et al.: CHARACTERIZATION AND WEAR PERFORMANCE OF CrAgN THIN FILMS ...
168 Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170
Figure 11: Light micrograph showing the microstructure of a CuSn6
ball
Slika 11: Svetlobni posnetek mikrostrukture kroglice CuSn6
Figure 10: Plan view optical images showing the wear tracks of the
CrAg3N film grown at 500 °C, developed during contact with CuSn6:
a) testing at ambient temperature, b) testing temperature of 400 °C
Slika 10: Svetlobni posnetek sledov obrabe na CrAg3N-nanosu, nane-
senem pri 500 °C, nastal med kontaktom s CuSn6: a) preizkus pri
sobni temperaturi, b) temperatura preizkusa 400 °C
thickness, but a higher Ag content led to a greater
thickness of the film.
The nanohardness and the Young modulus of the pure
CrN film were 16.79 GPa and 244 GPa, respectively. A
small addition of silver has almost no impact on both the
nanohardness and the Young’s modulus. The incorpora-
tion of 15 % Ag into the film induced a reduction of both
the nanohardness and the Young’s modulus. There is
probably a threshold of Ag content below which the
impact of the silver on the mechanical properties of
chromium nitride can be considered as negligible.
The addition of 3 % Ag improved the adhesion of the
CrN film, whereas the improvement was considerably
higher in the case of the film grown at 500 °C. The
adhesion of the film with 15 % Ag was very poor.
It has been established that there is almost no effect
of silver addition on the friction coefficient when tested
at room temperature against alumina, but the testing
against the same counterpart at higher temperature gave
a positive effect of the silver addition on the μ.
The testing against 100Cr6 steel gave a higher fric-
tion coefficient than that against the alumina, while the
testing against the CuSn6 bronze led to lower μ.
The addition of 3 % Ag to the CrN increased the
wear performance at elevated temperatures, while the
addition of 15 % Ag made the film too soft and sensitive
to wear, which resulted in a partial removal of the film
from the substrate inside the wear tracks.
Based on the obtained results, it seems that the
optimal silver addition into the CrN film lies between
3 % and 15 %. Hence, silver concentrations of 7 % and
11 %, respectively, are going to be investigated (determi-
nation of the most important coating characteristics, i.e.,
microstructure, hardness, Young’s modulus, wear resis-
tance, friction coefficient) in the near future.
Acknowledgements
This paper is the result of the project implementation:
CE for the development and application of diagnostic
methods in the processing of metallic and non-metallic
materials, ITMS: 26220120048.
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170 Materiali in tehnologije / Materials and technology 48 (2014) 2, 159–170
N. KARAKUÞ, H. Ö. TOPLAN: SYNTHESIZING Si3N4 FROM A MIXTURE OF SiO2-CaO
SYNTHESIZING Si3N4 FROM A MIXTURE OF SiO2-CaO
SINTEZA Si3N4 IN ME[ANICE SiO2-CaO
Nuray Karakuþ, Hüseyin Özkan Toplan
Sakarya University, Engineering Faculty, Dept. of Metallurgy and Materials Engineering, 54187 Sakarya, Turkey
nurayc@sakarya.edu.tr
Prejem rokopisa – received: 2012-08-16; sprejem za objavo – accepted for publication: 2013-07-05
In this study -phase-rich Si3N4 powders were synthesized containing the sintering additive (CaO) using carbothermal reduction
and nitridation. The starting agent for the silicon source was high-purity (99 %) synthetic silica. Carbon was added to the
high-purity SiO2 above the stoichiometric amount of oxygen. CaCO3 (for the silicon nitride containing the mass fraction of CaO
w = 10 %) was premixed in the starting reactants depending on the final powder composition and the type and amount of the
secondary phases required for sintering. The synthesis was carried out in a tube furnace in different temperature ranges (1400
°C, 1450 °C and 1475 °C for 3h) under a nitrogen-gas atmosphere. Having completed the synthesis process, the powder proper-
ties were examined using standard characterization techniques (XRD, SEM, etc.).
Keywords: silicon nitride, carbothermal reduction, CaO, SiO2
V tej {tudiji so bili sintetizirani s karbotermi~no redukcijo in nitridacijo z -fazo bogati prahovi Si3N4 z dodatkom za sintranje
(CaO). Za~etni izvir silicija je bil zelo ~isti (99 %) sinteti~ni SiO2. Ogljik je bil dodan zelo ~istemu SiO2 nad stehiometri~no
koli~ino kisika. CaCO3 (za silicijev nitrid je vseboval 10 % CaO) je bil prime{an za~etnim reaktantom, odvisno od kon~ne
sestave prahu in vrste ter koli~ine sekundarnih faz, potrebne za sintranje. Sinteza je bila izvr{ena v cevni pe~i pri razli~nih
temperaturah (1400 °C, 1450 °C in 1475 °C, 3 h) v atmosferi du{ika. Po sintranju so bile lastnosti prahu ugotovljene z uporabo
navadnih tehnik za karakterizacijo (XRD, SEM itd.).
Klju~ne besede: silicijev nitrid, karbotermi~na redukcija, CaO, SiO2
1 INTRODUCTION
Silicon nitride (Si3N4) ceramics have a range of struc-
tural applications, such as engine components, heat
exchangers, pump-seal materials, ball bearings, cutting
tools, etc., owing to their excellent mechanical properties
at both room and elevated temperatures.1 The material
properties of Si3N4 have led to speculation that it may
also have a role in biomedical fields, since it is biocom-
patible and is visible on plain radiographs as a partially
radiolucent material.2
The most common methods for silicon nitride prepa-
ration are the carbothermal reduction and nitridation
(CRN) of silica, the direct nitridation of silicon and the
thermal decomposition of silicon diimide. The carbother-
mal reduction used in this study takes place according to
the following overall reaction:3
3SiO2 + 6C + 2N2  Si3N4 + 6CO (1)
One of the difficulties found in the fabrication pro-
cess is the sintering applied to attain high relative den-
sities. Therefore, the use of additives to form a liquid
phase is required.4 Sintering aids such as MgO, Y2O3 and
Al2O3 added to -Si3N4 powders must be homogeneously
distributed and possess the desired powder composition
before shaping and sintering.5 In this study -phase-rich
Si3N4 powders were synthesized containing a sintering
additive using carbothermal reduction and nitridation.
Thus, Si3N4 powder is ready for sintering and in addition
is produced in a single step. Calcium oxide (CaO) was
chosen as a metal oxide additive and it was mixed before
synthesizing the Si3N4. Silicon nitride ceramic powders
synthesized using this method might therefore be readily
sintered because homogeneously distributed sintering
additives were present in the starting materials. For this
reason, the processing parameters are described in terms
of the powder-synthesis conditions.
2 EXPERIMENTAL
For the CRN process, the raw material was high-pu-
rity synthetic silica of nearly colloidal range and it was
supplied from EGE Kimya A. S. Activated charcoal was
used as a reducing agent and it was supplied by TÜ-
PRAÞ (Turkish Petroleum Refineries Co). The properties
of the silica are given in Table 1. CaCO3 was used as the
CaO source; it was supplied by Çelvit Company. The
CaCO3 (for the silicon nitride containing w = 10 % CaO)
was premixed in the starting reactants, depending on the
final powder composition and the type and amount of the
secondary phases desired for sintering. Carbon was
added to the high-purity SiO2 above the stoichiometric
amount of oxygen (w(C)/w(SiO2) ratio of 3). Dry mixing
was performed by ball milling for 10 h with alumina
Materiali in tehnologije / Materials and technology 48 (2014) 2, 171–173 171
UDK 621.762.5:666.3/.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)171(2014)
Table 1: Properties of silica (with firm data)
Tabela 1: Lastnosti SiO2 (podatki dobavitelja)
Properties SiO2
Purity (%)
Grain size (ìm)
Specific surface area (BET) (m2 g–1)
99
14
139
balls for the composition. The nitrogen gas (99.99 % in
purity) that was used for the nitridation process was
supplied by BOS.
The carbothermal reduction and the nitridation pro-
cess (CRN) were carried out in an atmosphere-controlled
tube furnace. The synthesis was performed under a
nitrogen gas flow (1 L/m) at 1400 °C, 1450 °C and 1475
°C for 3 h. After holding at these temperatures and this
time, the furnace was allowed to cool to room tempe-
rature. The gas flow was stopped during cooling when
the temperature reached 900 °C. After the CRN process,
the products were heated in air for 1 h at 900 °C for resi-
dual carbon burning.
After the CRN process at different temperatures, the
powder properties were examined by using standard cha-
racterization techniques (XRD, SEM, EDS).
3 RESULTS AND DISCUSSION
The silica was mixed with w(CaO) = 10 % (based on
the final product) and carbon black for the CRN process.
The prepared mixture was subjected to the CRN process
for 3 h at 1400 °C, 1450 °C and 1475 °C. The XRD ana-
lyses of the produced powders are shown in Figure 1.
The -Si3N4 was formed at a temperature of 1400 °C
after 3h of reaction. In addition, small amounts of uncon-
verted SiO2 and CaN2O3 phase were found in the product
powders. The temperature was increased to 1450 °C, but
the -Si3N4 formation was not completed at this tempe-
rature, which indicates that the holding time and/or tem-
perature were not enough to complete the reaction. The
best result was obtained at 1475 °C after a 3 h reaction.
At this temperature (1475 °C), the produced powder was
converted to -Si3N4 and -Si3N4, and in addition a
minor amount of CaSi2N2O2 was present. The formation
of the CaSi2O2N2 phase by a solid-state reaction was
completed at 1300–1400 °C for 10 h.6,7
The complete formation of -Si3N4 without using an
additive (without CaO) was obtained at 1475 °C for a 6 h
reaction, but the -phase was not dedicated.8 As seen
from the XRD data, the carbothermal reduction and nitri-
dation took place at 1475 °C. Comparing this with our
previous work, the optimum temperature is the same for
the product powders that used Y2O3 and Y2O3-Li2O and
25 °C higher than the products that used MgO.9 The
powder produced at 1475 °C was fully converted to 
and -Si3N4, and a minor amount of CaSi2N2O2 phase, as
expected, were present. Different types of additives have
been employed for the pressureless and pressure-assisted
sintering of Si3N4 ceramics.
The melting temperature of the CaO-SiO2 system is
approximately 1436 °C, as indicated by the binary pha-
se-equilibrium diagram.10 However, the presence of N
lowers these eutectic temperatures further. The alkali and
alkaline-earth oxides have a low melting point and the
viscosity of the resulting liquid is also low. The solu-
tion-diffusion-precipitation processes are enhanced.11
Therefore, the CaSi2N2O2 phase obtained after the CRN
process can be desirable as the sintering aids for later use
in sintering. Studies on the sintering of these synthesiz-
ing powders will be the focus of future work.
SEM micrographs of the powder synthesized from
SiO2 – w(CaO) = 10 % at 1400 °C and 1475 °C for 3 h
are given in Figure 2. The micron-sized Si3N4 grains
N. KARAKUÞ, H. Ö. TOPLAN: SYNTHESIZING Si3N4 FROM A MIXTURE OF SiO2-CaO
172 Materiali in tehnologije / Materials and technology 48 (2014) 2, 171–173
Figure 2: SEM micrographs of the product powder synthesized at: a)
1400 °C, 3 h and b) 1475 °C, 3 h
Slika 2: SEM-posnetka pra{natega produkta, sintetiziranega pri: a)
1400 °C, 3 h in b) 1475 °C, 3 h
Figure 1: Phases formed at different temperatures (1400–1475 °C)
after CRN of the SiO2 – w(CaO) = 10 % mixture
Slika 1: Faze, ki nastajajo pri razli~nih temperaturah (1400–1475 °C)
po CRN me{anice SiO2 in w(CaO) = 10 %
were formed after the CRN. SEM micrographs of the
powder produced at 1400 °C for 3 h revealed different
morphologies, ranging from irregular-shaped small par-
ticles to equiaxed small grains and long whiskers. The
long whiskers had a cross-section of approximately 300
nm. SEM micrographs of the produced powder at 1475
°C for 3 h revealed the same morphologies as small par-
ticles and long whiskers.
It was clear that the grain size was increased with in-
creasing temperature from 1400 °C to 1475 °C. The
whiskers had a thickness of approximately 1 μm after the
CRN process at 1475 °C. The EDS analysis of the pow-
ders synthesized from SiO2 – w(CaO) = 10 % at 1475 °C
for 3 h is showed in Figure 3. According to the EDS
analysis, the Si, N, O, and Ca elements were detected
from the particles in a manner consistent with the XRD
analysis (Figure 1).
4 CONCLUSIONS
-Si3N4 powders containing CaO as an oxide additive
were successfully synthesized by carbothermal reduction
and nitridation. The best result was obtained at 1475 °C
after a 3 h reaction. The Si3N4 powders showed two
major grain morphologies: submicron equiaxed and long
whiskers. The advantage of using pre-additive oxide
(CaO) in the CRN process is in terms of the reaction
temperature and time.
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schung, Stuttgart, 2003
N. KARAKUÞ, H. Ö. TOPLAN: SYNTHESIZING Si3N4 FROM A MIXTURE OF SiO2-CaO
Materiali in tehnologije / Materials and technology 48 (2014) 2, 171–173 173
Figure 3: EDS analysis of the powders synthesized from SiO2 –
w(CaO) = 10 % at 1475 °C for 3 h
Slika 3: EDS-analiza prahov, sintetiziranih iz SiO2 – w(CaO) = 10 %
pri 1475 °C po 3 h

V. BEDNÁØOVÁ et al.: CAST CELLULAR METALS WITH REGULAR AND IRREGULAR STRUCTURES
CAST CELLULAR METALS WITH REGULAR AND
IRREGULAR STRUCTURES
ULITE KOVINE S PRAVILNO IN NEPRAVILNO CELI^NO
STRUKTURO
Vlasta Bednáøová, Petr Lichý, Tomá{ Elbel, Ale{ Hanus
Department of Metallurgy and Foundry Engineering, FMMI, V[B - Technical University of Ostrava, 17. listopadu 15/2172, Ostrava – Poruba,
Czech Republic
vlasta.bednarova@vsb.cz
Prejem rokopisa – received: 2012-09-03; sprejem za objavo – accepted for publication: 2013-06-10
An appropriate way to reduce the weight of manufactured parts without adversely affecting their strength is to use porous
metallic materials with different internal arrangements of the intentionally created cavities. Porous metallic materials can be
made from liquid metal, from powdered metal, metal vapours, or from metal ions. The aim of this research was to verify the
possibilities of producing metallic foams by conventional foundry processes, to study the process conditions as well as the
physical and mechanical properties of the metal foams produced. The experiments demonstrated the possibility of
manufacturing castings with both a regular cellular structure and a solid skin in a single casting operation using different kinds
of preform made from commonly used resin-bonded core mixtures. From the perspective of the need to destroy the preforms
after the metal’s solidification it seems a very interesting prospect to produce preforms from different salts.
Keywords: cellular metals, metal foams, casting
Primerna pot za zmanj{anje mase izdelanih delov, ne da bi ob~utno vplivali na njihovo trdnost, je uporaba poroznih kovinskih
materialov z razli~no notranjo razporeditvijo namerno povzro~enih praznih jamic. Porozne kovinske materiale se lahko izdela iz
staljene kovine, kovinskih prahov, kovinskih par ali iz kovinskih ionov. Namen raziskave je bil preveriti mo`nost izdelave
kovinskih pen po navadnem livarskem postopku, {tudij pogojev procesa, fizikalne in mehanske lastnosti izdelane kovinske pene.
Izvr{eni poskusi so pokazali mo`nost izdelave ulitkov z urejeno celi~no strukturo in ~vrsto skorjo z eno livarsko operacijo z
uporabo razli~nih predoblik, izdelanih iz navadnih me{anic za jedra. S stali{~a potrebne razgradnje predoblike, potem ko se
kovina strdi, se zdi konkuren~na izdelava predoblike iz razli~nih soli.
Klju~ne besede: celi~ne kovine, kovinske pene, ulivanje
1 INTRODUCTION
Cellular metals and metallic foams are metallic
materials containing pores in their structure that are
created intentionally. To properly identify the material,
according to J. Banhart,1 one has to distinguish the
following: the term cellular metal is a general term
describing a material in which any kind of gaseous voids
are dispersed, in a porous metal the pores are usually
round and isolated from each other, while the terms foam
metal and metal foam are used for a porous metal
produced by foaming a melt in which the pores are not
interconnected ("structure with closed pores"). In addi-
tion, we have the term metal sponge, which is used for
highly porous materials, in which the pores are connec-
ted in a complicated manner and the structure cannot be
divided into individual cavities ("structure with open
pores"). The term metallic foam is, however, very often
used, even in the professional literature, as a general
designation of porous materials.
Since the discovery of porous metallic materials
numerous methods of production have been developed.
According to the state in which the metal is processed
the manufacturing processes can be divided into four
groups. Porous metallic materials can be made from:2
• liquid metal (e.g., direct foaming with gas, blowing
agents, powder compact melting, casting, spray form-
ing),
• powdered metal (e.g., sintering of powders, fibres or
hollow spheres, extrusion of polymer/metal mixtures,
reaction sintering),
• metal vapours (vapour deposition),
• metal ions (electrochemical deposition).
Over the past two decades, opportunities for the use
of porous metals have increased in many research and
industrial applications. Growing interest is related to the
specific characteristics of this material. Porous metals
represent a new type of materials that fulfil current eco-
logical requirements (particularly in the area of weight
reduction) and have unique service properties thanks to a
structure with low densities, large specific surfaces, and
a useful combination of physical and mechanical
properties (absorption of energy, absorption of sound and
vibration, thermal insulation, and heat exchange).
Despite the uniqueness of the properties and the wide
range of possibilities of use, the number of examples of
practical, stable, industrial applications is not large.
Nickel foam materials are mass-produced and used as
electrodes in rechargeable batteries for portable devices
(mobile phones, laptops), while titanium foams are used
Materiali in tehnologije / Materials and technology 48 (2014) 2, 175–179 175
UDK 621.74 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)175(2014)
as implants. At present, the most explored porous metal
materials are aluminium foams, which can be documen-
ted by the development of prototypes as well as practical
applications (small series for AUDI and Lamborghini
crash zones).3 The reasons for the limited use of cellular
metals are primarily economic. This is why the current
interest (as reflected in a growing number of new pub-
lished studies and possibilities for industrial applica-
tions4–7) is focused mainly on the development of
technologies enabling the production of foam materials
at costs that would allow their widespread use.
The research work carried out at the Department of
Metallurgy and Foundry Engineering of the VSB – Tech-
nical University aims to study a technique for manufac-
turing porous metals using the conventional principles of
gravity casting in sand moulds. Our intention is to use
procedures for liquid-metal processing, enabling the
production of complex-shaped castings, not only a
simple block of metal foam. The research focuses on two
approaches that use the standard foundry technique of
casting into disposable sand moulds or permanent
metallic moulds with the use of precursors and ceramic
prefabricated parts prepared by known methods for core
making. These approaches bring low overall costs as
well as the advantage of the direct production of parts
from complex shapes without any necessary further
forming, welding, machining and with use of traditional
foundry procedures.
The production of metallic foams using conventional
gravity casting into foundry moulds presupposes a more
cost-efficient process than other known technologies
such as powder metallurgy, metal evaporation and ioni-
zation. Cast metallic foams are not yet produced in the
Czech Republic.
2 EXPERIMENTAL
The subject of the research was the testing of
infiltration techniques for the manufacture of porous
metals with both regular and irregular structures. The
foundry methods used can be divided into two basic
groups:
• two-stage investment casting
• infiltration of liquid metal into the mould cavity filled
with different types of filling materials
2.1 Two-stage investment casting
According to this process, a polymer foam, e.g.
polyurethane foam, is used as a "lost foam pattern". The
pattern is first infiltrated by a slurry of the plaster, to
create a plaster "investment". This then undergoes a heat
treatment that solidifies the "investment" and burns the
polymer foam. Molten metal is poured into the prepared
investment, and after the metal’s solidification the invest-
ment must be removed to leave the internal cavities.
Plaster slurry composition:
100 mass parts of Goldstarpowders MO28 plaster
50 mass parts of water
The material used for the casting was the
AlSi10MgMn alloy.
2.2 Infiltration techniques
Infiltration techniques are based on the infiltration of
a liquid metal between various filler materials (called a
preform or precursor) placed in the mould cavity. Since
they do not occupy the entire volume, these precursors
form a network of interconnected porosities. The pre-
form must be made of a material that retains its shape
during the liquid metal’s infiltration (sufficient strength,
low abrasion) and can be destroyed after the casting to
leave the cavities. The preform must not contain a dis-
connected island of material so that it can be completely
eliminated from the solidified metal.
2.2.1 Use of a precursor as a filler material
Material of precursors that were inserted into the
green sand mould cavity:
• granules made of ceramic material with fractions of
8–16 mm and 3–4 mm (Figure 1).
• granules made of resin-bonded moulding mixtures
(shell-mould, CO2 resol), fractions 8–16 mm (Figure
2).
Apart from the different granulometry, several de-
signs of gating system were used, and the temperature of
the filler material was changed (in order to positively
affect the castabily of the molten metal).
V. BEDNÁØOVÁ et al.: CAST CELLULAR METALS WITH REGULAR AND IRREGULAR STRUCTURES
176 Materiali in tehnologije / Materials and technology 48 (2014) 2, 175–179
Figure 2: Shell-mould particles inserted into mould cavity
Slika 2: Lupinasti delci, vlo`eni v livno votlino
Figure 1: Ceramic particles with fractions 8–18 mm
Slika 1: Kerami~ni delci zrnatosti 8–18 mm
The materials used for the casting were the
AlSi10MgMn alloy and cast iron with lamellar graphite
– in accordance with EN GJL-200.
2.2.2 Use of preform as a filler material
A regular cellular structure can be achieved using
different types of preforms that fill the mould cavity.
Using a preform like a core not filling the whole mould
cavity can make it possible to manufacture a casting with
a solid surface layer and an internal porous structure
with defined cell dimensions.
For the manufacturing of a sand preform/core it was
necessary to prepare a core box. Fused Deposition
Modelling technology was used to manufacture the core
box. Due to the complex lattice shape of the core, a
sectional core box was designed consisting of five parts,
which allowed the easy removal of the core.
The materials used for the preform/core manufac-
turing were:
• CO2 hardened alkaline phenolic process (the resin is
an alkaline phenolic one, containing a linking sub-
stance stabilised at a high pH, curing occurs by
gassing with carbon dioxide, which dissolves in the
water solvent of the resin, so lowering its pH and
activating the linking substance.)
• and KCl salt (Figure 3).
3 RESULTS AND DISCUSSION
3.1. Two-stage investment casting
The pattern was made of polyurethane foam, the
cavities of which was first filled by plaster (100 mass
parts of plaster, 50 mass parts of water). The plaster
investment was then allowed to dry completely, by
subsequently annealing at (100, 500, 640) °C for drying
and evaporation of a polyurethane foam. Melting and
pouring were made with an INDUTHERM MC 15
device. The equipment enables a combination of vacuum
and high pressure to ensure full infiltration. Pouring off
takes place using a 90° rotation of the casting unit. The
AlSi10MgMn alloy was poured at a temperature 750 °C
under a reduced pressure of 1 bar. After pouring, the
INDUTHERM MC 15 device automatically switches to
an overpressure of 2 bar in order to optimize the mould
filling, even for delicate parts. The plaster was removed
by dissolving out in water. In the case of a very fine me-
tal foam there are sometimes problems with removing
the ceramic without damaging the metal foam. The ob-
tained castings are the exact replicates of the original poly-
mer foam (Figure 4) and exhibit the highest porosities
(80–97 %). This type of foam provides a very promising
application, for example, in the filtration of liquid
metals.
Investment casting can also be used to obtain castings
with both high porosities and a regular cellular structure.
The pattern was made from extruded polystyrene
foam, forming regular structures made from elements
bonded layer by layer, as shown in Figure 5. The pattern
cavities were filled with a plaster slurry (100 mass parts
of plaster, 35 mass parts of water). After solidification
and sufficient drying of the matrix prepared in this way
the pattern/investment was subjected to 8 h of gradual
annealing (Figure 6) in order to remove the polystyrene
V. BEDNÁØOVÁ et al.: CAST CELLULAR METALS WITH REGULAR AND IRREGULAR STRUCTURES
Materiali in tehnologije / Materials and technology 48 (2014) 2, 175–179 177
Figure 5: Polymer foam pattern with regular internal structure
Slika 5: Predoblika s pravilno notranjo strukturo iz polimerne pene
Figure 3: Test samples of KCl salt
Slika 3: Preizku{anci iz soli KCl
Figure 4: AlSiMg alloy near-net-shape metal foam, porosity 90 %
Slika 4: Drobna mre`a kovinske pene iz zlitine AlSiMg s poroznostjo
90 %
foam and to achieve sufficient strength as the plaster
matrix is subjected to large loads, both thermal and
mechanical, during the pouring off.
The plaster pattern was inserted into the cavity of the
commonly used green sand mould (bentonite bonded
mixture) and then the AlSi10MgMn alloy was poured at
a temperature of 750 °C. The plaster was removed by
dissolving into a water bath and its residues were
removed by a mechanical force in an ultrasonic bath.
The manufacture of castings with a regular cellular
structure (Figure 7) by investment casting makes it
possible to achieve internal cavities (cells) with a preci-
sely defined shape, and thus predicable mechanical pro-
perties and reproducible results.8,9 The main disadvan-
tages of the production process are a high complexity
and rather high costs.
3.2. Infiltration techniques
3.2.1 Use of a precursor as a filler material
The performed experiments using different types of
precursors10 proved the feasibility of this method for
manufacturing porous metals. The condition of extrac-
tion of the precursors after solidification of the metal is
an obstacle for the use of ceramic particles because it is
generally difficult to remove the ceramic to leave the
cavities. More advantageous is the use of granules
manufactured from the used resin-bonded core mixtures,
which can be easily removed, depending on the pouring
temperature and the core mixture collapsibility. In the
case of the worse collapsibility the filler material may
remain in the casting (Figure 8) and its removal would
require an additional annealing of the castings (to
destroy the resin), which would slightly increase the
energy consumption of such production.
The main characteristic of the method using different
kinds of precursors as a filler material is the irregular
structure and the random distribution of pores throughout
the volume of the casting, and therefore the impossibility
of achieving reproducible results (Figure 9).
3.2.2 Use of a preform as a filler material
The performed experiments have demonstrated the
possibility of manufacturing castings with both a regular
cellular structure and solid skin in a single casting
operation using different kinds of preform made from
commonly used resin-bonded core mixtures. The
specified resin-core mixture11 has a sufficient strength
and abrasion resistance enabling the handling operation
and good collapsibility allows destroying it after
solidification.
From the perspective of the necessity to destroy
preforms after metal solidification (Figures 10 and 11) it
seems a good strategy to produce preforms from
V. BEDNÁØOVÁ et al.: CAST CELLULAR METALS WITH REGULAR AND IRREGULAR STRUCTURES
178 Materiali in tehnologije / Materials and technology 48 (2014) 2, 175–179
Figure 9: Clean cavities of cast iron EN GJL-200
Slika 9: O~i{~ene praznine v litem `elezu EN GJL-200
Figure 7: AlSiMg alloy castings with a regular cellular structure,
porosities 60–68 %
Slika 7: Ulitek iz zlitine AlSiMg z enakomerno celi~no strukturo s
poroznostjo 60–68 %
Figure 8: Casting made of AlSi10MgMn with remaining precursors
Slika 8: Ulitek, izdelan iz AlSi10MgMn s preostalo predobliko
Figure 6: Thermal treatment of plaster investment to burn the polymer
foam
Slika 6: Toplotna obdelava forme iz mavca, za odstranitev polimerne
pene
different salts. The tested salt was 100 % KCl, and the
cores were prepared by stamping with a pressing force to
10 t.
The use of the KCl made destroying the preform very
easy by dissolving out the water. But industrial imple-
mentation would require an extra waste-water treatment
circuit, which would make the application more compli-
cated.
4 CONCLUSIONS
Investment casting makes it possible to achieve the
highest porosity (97 %) in both types of porous castings.
In the case of very fine metal foams, problems may
occur with removing the ceramic without the metal foam
being damaged. Near-net-shape cast-metal foams pro-
vide a promising implementation, e.g., liquid-metal
filtration. The manufacturing of castings with a regular
cellular structure, using investment casting, makes it pos-
sible to achieve internal cavities (cells) with a precisely
defined shape, and thus predicable mechanical properties
and reproducible results. The main disadvantages of the
investment-casting production process are its high com-
plexity and thus rather high production costs. Using
different types of precursors as a filler material results in
porous castings with open pores in a stochastic arrange-
ment. The main drawback of the method is the irregular
structure and random distribution of the pores throughout
the volume of the casting. The performed experiments
have demonstrated the possibility of manufacturing
castings with both a regular cellular structure and a solid
skin in a single casting operation using different kinds of
the preform made from commonly used resin-bonded
core mixtures.
Mastering the production of metallic foams with a
defined structure and properties using gravity casting
into sand foundry moulds will contribute to an expansion
of the assortment produced in foundries by a completely
new type of material, which has unique service proper-
ties thanks to its structure, and which fulfils the current
ecological requirements. The manufacture of foams with
the aid of gravity casting in conventional foundry moulds
is a cost-advantageous process that can be industrially
used in foundries without high investment demands.
Metal foams are progressive materials with continuously
expanding use. Cast metallic foams are not yet produced
in the Czech Republic.
Acknowledgements
This work was supported by the Technology Agency
of the CR within the frame of the research project
TA02011333.
5 REFERENCES
1 J. Banhart, Manufacture, characterisation and application of cellular
metals and metal foams, Progress in Materials Science, 46 (2001),
559–632
2 J. Banhart, Manufacturing routes for metallic foams, Journal of
Minerals, Metals and Materials, 52 (2000) 12, 22–27
3 Y. Gaillard, J. Dairon, M. Fleuriot, B. W. Corson, Les materiaux
cellulaires: une innovation aux applications multiples, Fonderie
magazine, (2010) 1, 21–33
4 Cellmet News: http://www.metalfoam.net/cellmet-news_2006-1_
net.pdf
5 I. Paulin, et. al., Synthesis of aluminium foams by the powder
metallurgy process: compacting of precursors, Mater. Tehnol., 45
(2011) 1, 13–19
6 MetFoam2009: http://www.metfoam2009.sav.sk/index.php?ID=2571
7 J. Dairon, et al., Mousses métalliques: CTIF innove dans les maté-
riaux cellulaires, Fonderie – Fondeur d’aujourd’hui, (2009) 295,
12–19
8 M. Cholewa, M. Dziuba-Kalu`a, Analysis of structural properties of
skeleton castings regarding the crystallization kinetics, Archives of
Materials Science and Engineering, 38 (2009) 2, 93–102
9 I. Zyrjanova, Lité kovové pìny z Al slitin, diplomová práce, V[B-TU
Ostrava, 2011
10 A. Hanus, P. Lichý, V. Bednáøová, Production and properties of cast
metals with porous structure, Metal 2012, 21st International Confe-
rence on Metallurgy and Materials, Conference proceedings, 1–6
11 V. Bednáøová, P. Lichý, T. Elbel, Casting routes for porous metals
manufacturing, Proceedings book, 12th International Foundrymen
Conference, Sustainable Development in Foundry Materials and
Technologies, Opatija, 2012, 16–23
V. BEDNÁØOVÁ et al.: CAST CELLULAR METALS WITH REGULAR AND IRREGULAR STRUCTURES
Materiali in tehnologije / Materials and technology 48 (2014) 2, 175–179 179
Figure 11: Detail of AlSiMg alloy casting with regular cellular
structure
Slika 11: Detajl ulitka iz zlitine AlSiMg z enakomerno celi~no
zgradbo
Figure 10: AlSiMg alloy castings with a regular cellular structure
Slika 10: Ulitki iz zlitine AlSiMg z enakomerno celi~no zgradbo

P. BRGLEZ et al.: SPIN-COATING FOR OPTICAL-OXYGEN-SENSOR PREPARATION
SPIN-COATING FOR OPTICAL-OXYGEN-SENSOR
PREPARATION
UPORABA SPINSKEGA NANOSA PRI IZDELAVI OPTI^NIH
SENZORJEV ZA KISIK
Polonca Brglez1,2, Andrej Holobar1, Aleksandra Pivec3, Mitja Kolar2,4
1ECHO, d. o. o., Stari trg 37, 3210 Slovenske Konjice, Slovenia
2University of Maribor, Faculty of Chemistry and Chemical Engineering, Smetanova 17, 2000 Maribor, Slovenia
3ZRS Bistra Ptuj, Slovenski trg 6, 2250 Ptuj, Slovenia
4Centre of Excellence PoliMaT, Tehnolo{ki park 24, 1000 Ljubljana, Slovenia
mitja.kolar@um.si
Prejem rokopisa – received: 2012-09-27; sprejem za objavo – accepted for publication: 2013-06-18
Thin-film oxygen sensors were prepared using the spin-coating technique, where a tris (4,7-diphenyl-1,10-phenanthroline)
ruthenium(II) dichloride complex (RuDPP) in various solvents and silicones deposited on different substrates was used for the
sensor production. By changing the spin-coating set-up parameters, homogeneous sensor coatings and the optimum sensor
response to oxygen were studied – the sensors were exposed to various concentrations of oxygen within the range from 0 % to
100 %. During the presented study, the optimum results were obtained when a 150 μL of sensor solution was applied to a
Dataline foil using silicone E4 and a chloroform solvent. A spin coater with the following three rotation stages was used:
750/700 r/min for 3 s, 300 r/min for 3 s and 150 r/min for 4 s. The spin-coating technique has several benefits: it is fast, easy to
use and appropriate for low-volume operations. It allows modifications and preparations of several sensor series using the
minimum reagent consumption. However, the disadvantage of this technique also has to be mentioned, namely, an uneven film
thickness in the radial direction. The film thickness mainly depends on the experimental set-up (volume, rotation time and
speed, solvent viscosity and evaporation). Spin coating as an alternative and very flexible technique for an oxygen-sensor
preparation is suggested for the laboratory-scale work, where the majority of experimental data could be used when other new
coating methods are also researched and implemented.
Keywords: tris (4,7-diphenyl-1,10-phenanthroline) ruthenium(II) dichloride complex, spin coating, optical oxygen sensor,
oxygen
Izdelani so bili tankoplastni opti~ni senzorji za kisik s tehniko spinskega nanosa. Pri tem so bile uporabljene razli~ne kon-
centracije tris (4,7-difenil-1,10-fenantrolin) rutenij(II) diklorid kompleksa (RuDPP), razli~na topila, polimerni nosilci, silikoni
in parametri spinskega prekritja. Na{ namen je bil pripraviti najbolj homogen nanos senzorske raztopine in tako dobiti najbolj
optimalne lastnosti senzorjev. Preu~evali smo tudi vpliv hitrosti in ~asa vrtenja spinske naprave za prekrivanje na odziv
senzorjev, saj so bili le-ti po izdelavi izpostavljeni razli~nim koncentracijam kisika v obmo~ju od 0 % do 100 %. Najbolj{i nanos
senzorske raztopine smo dobili s senzorsko raztopino v kloroformu 150 μL z uporabo silikona E4 z nanosom na folijo Dataline.
Pri tem smo uporabili tri razli~ne stopnje vrtenja: 3 s pri 750/700 r/min, 3 s pri 300 r/min in 4 s pri 150 r/min. Prednost uporabe
spinskega prekrivanja je, da je ta tehnika zelo hitra, enostavna za uporabo in je primerna za nanos majhnih prostornin. Omogo~a
izdelavo ve~ serij senzorjev z razli~nimi lastnostmi ob minimalni porabi reagentov. Nanos senzorske raztopine na polimernem
nosilcu v radialni smeri je v veliki meri odvisen od eksperimentalnih razmer: prostornine nanosa, hitrosti vrtenja, viskoznosti in
hlapnosti topil. Metoda spinskega prekritja se je izkazala kot u~inkovita metoda za nanos senzorskih raztopin v laboratorijskem
merilu, vendar je po celotni senzorski povr{ini te`ko pripraviti popolnoma homogen nanos, zato je za pripravo ve~jih koli~in
identi~nih senzorjev – po optimiranju vseh drugih eksperimentalnih razmer – smiselno preu~iti {e alternativne metode
nana{anja.
Klju~ne besede: tris (4,7-difenil-1,10-fenantrolin) rutenijev(II) diklorid kompleks, spinsko prekritje, opti~ni kisikov senzor,
kisik
1 INTRODUCTION
Oxygen (O2) is considered to be one of the more
important gases in our environment. The determination
of O2 concentrations in the air, especially at low levels,
plays an important role in different areas ranging from
environmental, biological, analytical and industrial
monitoring. These are the reasons why there is still a
growing interest in the construction and development of
oxygen sensors.1–4
There has been a trend in the development of optical
oxygen sensors over the last few decades because these
sensors are more attractive than conventional ampero-
metric sensors. Optical oxygen sensors have a lot of
advantages such as: a faster response time, a high sensiti-
vity and selectivity, no O2 consumption, the inertness
against sample flow rate or stirring speed, absence of
poison, and no need for a reference electrode.5–15 They
are immune to exterior electromagnetic-field interference
and can be produced as disposable sensors.16,17 Optical
oxygen sensors are cheap, easily miniaturized and simple
to use; they mainly operate on the principle of oxygen
quenching those dye molecules that have been entrapped
within a porous support matrix. Ruthenium(II) com-
plexes are, by far, the most widely used oxygen dyes,
because they have relatively long fluorescent lifetimes
determined by the metal-to-ligand charge-transfer
(MLCT) excited state, fast response time, strong visible
absorption, large Stokes shift, and high photochemical
stability.2,3,11,12,14,18–23 Ru(II) complexes exhibit a high
Materiali in tehnologije / Materials and technology 48 (2014) 2, 181–188 181
UDK 543 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)181(2014)
sensibility to luminescence quenching and the positions
of their absorption and emission spectra permit an
application of low-cost, solid-state optoelectronics for
the detection of luminescence intensity. The dyes can be
excited with blue or even blue-green light-emitting
diodes (LEDs) exhibiting a large Stokes shift and result-
ing in the emission of orange-red light.24
The basic operational principle of a fluorescent opti-
cal sensor for measuring oxygen is based on reducing the
intensity of the fluorescence (quenching) due to the
involvement of oxygen within the dye structure. The
calibration of the most luminescence quenching-based
optical sensors relies, in essence, on the Stern-Volmer
equation.
The immobilization of the Ru(II) complexes in sol-
gel matrices has been recently investigated.4,25–27 There
have also been reports on optical oxygen sensors based
on the luminescence change of the ruthenium(II) com-
plex immobilized in organic and inorganic polymers
(polystyrene, silicone polymer, sol-gel glass, etc.) and
zeolite matrix.28–30 The sol-gel process has been so far
the most widely used method for the preparation of
oxygen sensors.5,31,32 It is an efficient immobilization
technique due to its many desirable properties such as
high thermal stability, good photostability and optical
transparency within the visible region.5
The spin-coating technique is also used for the sensor
preparation. Spin coating has been used for several
decades for the application of thin films.33,34 This is a
technique that uses centrifugal forces created by a spin-
ning substrate for spreading a coating solution evenly
over a surface.35 It can be controlled with a few parame-
ters in order to yield a well-defined coating coverage.36,37
The flow is governed by a balance between the centri-
fugal force against the viscosity and surface tension. It
has been shown that the non-uniform distribution in the
initial film profile tends to become uniform during
spinning. Spin coating has been mainly used in the
photoresist coating process because of its simplicity of
operation, its uniformity and the thinness of the coated
layers. The spin-coating process involves depositing a
small puddle of fluid onto the center of a substrate, and
then spinning the substrate at a high speed (typically
around 3000 r/min). The centripetal acceleration then
causes the solution to spread towards, and eventually off,
the edge of the substrate leaving a thin film on the
surface. The coat thickness is controlled with the rotatio-
nal speed of the substrate; faster rotations result in
thinner coating layers.
The spin-coating process needs to be reshaped and
optimized because of the changes in the operational
parameters and the wafer size.38 There is scientific litera-
ture describing the spin-coating process, emphasizing the
importance of rotational speed, time, acceleration,
periods, liquid viscosity, density, polymer, temperature
and humidity for the film thickness.39
There are four distinct stages in the spin-coating pro-
cess (Figure 1):
• Deposition of a coating fluid onto a wafer or
substrate.
• Acceleration of the substrate up to its final desired
rotational speed.
• Spinning of the substrate at a constant rate; fluid
viscous forces dominate the fluid thinning behavior.
• Spinning of the substrate at a constant rate; the
solvent evaporation dominates the coating thinning
behavior.
The final film thickness and other properties depend
on the nature of the used polymer (viscosity, drying rate,
percent of solids, surface tension) and the parameters
chosen for the spin process (final rotational speed,
acceleration). One of the more important factors in spin
coating is the repeatability. Subtle variations in these
parameters defining the spin process can result in drastic
variations in the coated films.
The presented goal was a preparation of thin-film
oxygen sensors using the spin-coating technique. In this
work a spin coater was used for spreading different
sensor solutions onto various polymer substrates. The
substrates (polymer solid layers – foils) were optically
transparent films. The most important function of the
substrate was to act as a strong mechanical carrier with a
high transparency, physical strength, and chemical
resistance. Different amounts of RuDPP in various
solvents were used for the sensor production and various
transparent polymer substrates were used as the carriers.
The goal was to obtain the most homogeneous sensor
coating with the spin coater by changing the set-up
parameters. After the sensor preparation the sensors were
exposed to various concentrations of oxygen, ranging
from 0 % to 100 %.
2 EXPERIMENTAL WORK
2.1 Chemicals and solutions
All the chemicals used were of analytical purity
grade. All the solutions were prepared with deionized
P. BRGLEZ et al.: SPIN-COATING FOR OPTICAL-OXYGEN-SENSOR PREPARATION
182 Materiali in tehnologije / Materials and technology 48 (2014) 2, 181–188
Figure 1: Scheme of the spin-coating sensor preparation
Slika 1: Shemati~en prikaz postopka spinskega nanosa pri izdelavi senzorjev
water. Silicon (Elastosil E4, Elastosil E41, Wacker), a
polymer layer (foil DATALINE 57170, Dataline, EU;
foil PLASTIBOR TOP COD 12530 12950, Lazertechas,
UAB; foil ESSELTE 509700, Esselte, EU), a tris
(4,7-diphenyl-1,10-phenanthroline) ruthenium(II) dichlo-
ride complex (Sigma-Aldrich), toluene (Sigma-Aldrich),
chloroform (CHLO) (Sigma-Aldrich) and methyl ethyl
ketone (ME) (Sigma-Aldrich) were used for the sensor
preparation.
The following gases from Messer, d. o. o., Slovenia,
were used for testing an optical oxygen sensor: nitrogen
(N2, 99.999 %) and oxygen (O2, 99.9999 %).
2.2 Apparatus
Optical measurements were studied using an
EOM-O2 micro electro-optical module (PreSens) con-
trolled by the EOM-O2_v1_3_exe software, a gas-mixing
device (Echo, d. o. o.) and a flow cell (Echo, d. o. o.).
Additional equipment included: an AB54-S balance
(Mettler Toledo), a spin-coater (Polos), a MST digital
magnetic stirrer (Ika) and a SUPRA 35 VP (Carl Zeiss)
scanning-electron microscope (SEM).
2.3 Preparation of RuDPP optical oxygen sensors
Different amounts of RuDPP within the range of 20
mg to 80 mg were diluted using different solvents
(toluene, chloroform and methyl ethyl ketone).
An appropriate amount of RuDPP was weighted in a
10 mL flask and diluted using an appropriate solvent.
The prepared sensor solution was then being stirred with
a magnetic stirrer for 10 min. The sensor solution was
then filtered through filter paper and a 4 mL sensor solu-
tion was added to 2 g of silicone. This sensor solution
was mixed on a magnetic stirrer for about 1 h to become
homogeneous and viscous. The sensor solution was
protected from the external light with an aluminum foil,
and was applied to the solid layers using the spin-coating
technique. Different polymer solid layers (foils) were
used for the substrate (ESSELTE, DATALINE and
PLASTIBOR). Different amounts ((100, 150 and 200)
μL) of the sensor solutions were applied using the
spin-coating technique. The effects of changing the
rotation speeds and times of spinning were also studied;
the details are given in Table 1.
Table 1: Rotation speed and spinning time
Tabela 1: Hitrost vrtenja in ~as vrtenja naprave za spinsko prekrivanje
Stage Number of turns(r/min)
Time of spinning
(s)
1st step 500 to 900 1–10
2nd step 300 to 700 1–10
3rd step 100 to 150 1–10
The sensor solution was mounted on a rotating
platform. The substrate was rotated according to the
selected rotation speed/spinning time and the sensor
solution was dispensed directly onto it. The high-speed
rotation threw off most of the solution, leaving behind a
thin, uniform coating.
The prepared optical sensors were then dried; they
were usually left to dry out for 24–48 h at a room tem-
perature of (20 ± 2) C°. After drying, the optical sensors
were cut to the diameter dimensions of 1.75 cm2/15 mm.
The sensors were stored in a dark and dry place before
use.
2.4 Measurement procedures
The optical oxygen sensors were tested in a flow cell
(Figure 2). They were excited with a blue LED and
measured with an optical detector from PreSens. The gas
mixtures (N2/O2) passed the active sensor surfaces at a
constant flow rate of 1 L/min.
The changes of the signal were measured for diffe-
rent concentrations of oxygen. The gas mixtures were
prepared with a gas mixing device (Echo, d. o. o.).4 Dur-
ing the constant flow of the carrier gas, various concen-
trations of oxygen were added to obtain different con-
centrations within the range of 1 · 10–6 to 1000 · 10–6.
The accuracy of the concentrations was ± 0.7 · 10–6. The
gas-mixing device provided a repeatability of ± 0.15 %
and, within the full-scale mode, the temperature range
was from 15 °C to 25 °C and the pressure varied from 70
kPa to 400 kPa. Figure 3 schematically presents the
system used for the optical measures. The measuring
system consisted of: a gas-mixing device, a flow cell, an
electro-optical module and a computer.
Surface analyses of the sensors were performed with
a scanning electron microscope (SEM), Supra 35 VP
P. BRGLEZ et al.: SPIN-COATING FOR OPTICAL-OXYGEN-SENSOR PREPARATION
Materiali in tehnologije / Materials and technology 48 (2014) 2, 181–188 183
Figure 2: Scheme of the flow cell (left) and sensor positioning (right)
Slika 2: Shema preto~ne celice (levo) in namestitev senzorja (desno)
Carl Zeiss. All the pictures were recorded using a 30 μm
scan window at the 1 kV electronic potential.
3 RESULTS AND DISCUSSION
The influences of the dye concentration, different
polymer solid layers (foils), silicones, film thickness and
different solvents on the sensor sensitivity were studied.
3.1 Influence of the RuDPP concentration vs. the sen-
sor sensitivity
The Stern-Volmer equation describes the fluorescen-
ce intensity versus the measured concentration of oxy-
gen.4 A deviation from the linearity is connected with the
heterogeneity of a polymer matrix; the fluorophore
molecules are usually surrounded by voids and polymer
particles, therefore, all the indicator molecules are non-
equally accessible to oxygen. A decline in fluorescence
is strongly dependent on the diffusion and adsorption of
oxygen and on the dye solubility. The concentration of
the indicator must be appropriately selected in order to
obtain the optimum sensor sensitivity, and the dye con-
centration must be additionally optimized according to
the measure range.
The sensors were prepared according to the proce-
dure described in Section 2.3. The amounts of (20, 40,
60 and 80) mg of RuDPP were used for preparing the
sensor solution (Figure 4), while the linearity (R2) was
tested within the range of 0 % to 100 % concentration of
oxygen.
The linearity of the sensors prepared from 20 mg of
RuDPP was 0.9309, for 40 mg of RuDPP it was 0.9691,
for 60 mg of RuDPP it was 0.9888, and 0.9904 for 80
mg of RuDPP. The optimum sensor was the one with 80
mg of RuDPP, therefore, it can be concluded that the
concentration of RuDPP strongly influences the sensor
response. In general, with higher concentrations of
RuDPP, the sensor linearity, accuracy and precision are
improved. In addition, a strong fluorescence signal was
obtained and no additional amplification of the measur-
ing signal was used. The electronic-optical noise usually
increased with a higher amplification rate, which can
also be a reason for a nonlinear sensor response. On the
other hand, due to the high cost of RuDPP, it is important
to incorporate low dye concentrations. In order to pre-
pare sensors with different properties, typical amounts of
(40, 60 and 80) mg of RuDPP were used for further
studies.
3.2 Sensor preparation – modification of foils and sili-
cones
In the next step different foils and silicones as the
support matrices were tested. In order to optimize the
sensors, different foils (Plastibor, Dataline, Esselte) and
commercially available silicones (E4, E41) were used.
The linearities and sensitivities (k) of different sensors
were tested; Table 2 presents all the major sensor cha-
racteristics. Figure 5a presents the changes in the
measured signal versus the various concentrations of
oxygen with different foils, and Figure 5b shows the
change in the signal with different silicones. When using
the Datalain foil a slightly better linear response was
obtained, especially at low concentration ranges of
oxygen, when compared to the Plastibor or Esselte foils,
but the selection of the solid layers does not significantly
improve or change the sensor properties. It is generally
known that with the increasing roughness of a substrate
foil the adhesion of the coatings on the surface is
improved even in spin coating. Here, it is important to
mention that rough surfaces cause a lower transparency
with a significant back-scattering light effect, and for this
reason we used low-roughness foils. Additionally, a
compromise between the foil transparency and surface
P. BRGLEZ et al.: SPIN-COATING FOR OPTICAL-OXYGEN-SENSOR PREPARATION
184 Materiali in tehnologije / Materials and technology 48 (2014) 2, 181–188
Figure 3: Scheme of the measuring system
Slika 3: Shema merilnega sistema
Figure 4: Influence of RuDPP concentration vs. sensor response
Slika 4: Vpliv koncentracije RuDPP na odziv senzorjev
P. BRGLEZ et al.: SPIN-COATING FOR OPTICAL-OXYGEN-SENSOR PREPARATION
Materiali in tehnologije / Materials and technology 48 (2014) 2, 181–188 185
Figure 5: a) Impact of a solid layer (a foil) on sensor response (signal/a.u. – arbitrary units vs. concentration of O2 in %), b) comparison of
different silicones vs. sensor response
Slika 5: a) Vpliv trdnega nosilca (folije) na odziv senzorjev (izmerjen signal (a.u. enote) vs. koncentracija O2 v %), b) vpliv silikona na ob~utlji-
vost senzorja
Table 2: Influence of the solid layers on the sensor characteristics
Tabela 2: Vpliv trdnih plasti na lastnosti senzorjev
RuDPP/mg FOIL SILICONE SOLVENT V/μL STAGES k R2
60 DATALINE E4 ME 150
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
2.699 0.978
60 PLASTIBOR E4 ME 150
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
2.464 0.960
60 ESSELTE E4 ME 150
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
2.684 0.886
60 DATALINE E41 CHLO 200
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
4.141 0.964
60 PLASTIBOR E41 CHLO 200
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
4.516 0.974
60 ESSELTE E41 CHLO 200
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
4.531 0.957
80 DATALINE E41 ME 200
1300 r/min –› 3 s
2200 r/min –› 5 s
3100 r/min –› 2 s
–2.727 0.361
80 ESSELTE E41 ME 150
1300 r /min –› 3 s
2200 r/ min –› 5 s
3100 r/ min –› 2 s
–8.761 0.992
80 PLASTIBOR E41 ME 150
1300 r/ min –› 3 s
2200 r/min –› 5 s
3100 r/min –› 2 s
–5.709 0.834
80 DATALINE E41 CHLO 150
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
5.106 0.975
80 PLASTIBOR E41 CHLO 150
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
4.073 0.943
80 ESSELTE E41 CHLO 150
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
4.787 0.966
20 DATALINE E4 ME 150
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
4.277 0.735
20 PLASTIBOR E4 ME 150
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
3.693 0.606
P. BRGLEZ et al.: SPIN-COATING FOR OPTICAL-OXYGEN-SENSOR PREPARATION
186 Materiali in tehnologije / Materials and technology 48 (2014) 2, 181–188
20 ESSELTE E4 ME 150
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
3.962 0.635
20 DATALINE E4 CHLO 150
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
2.545 0.946
20 PLASTIBOR E4 CHLO 150
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
1.609 0.585
20 ESSELTE E4 CHLO 150
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
1.94 0.845
40 DATALINE E41 CHLO 200
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
4.531 0.957
40 PLASTIBOR E41 CHLO 200
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
3.857 0.959
40 ESSELTE E41 CHLO 200
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
4.077 0.865
40 DATALINE E4 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
6.468 0.528
40 PLASTIBOR E4 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
6.654 0.524
40 ESSELTE E4 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
7.003 0.524
40 DATALINE E41 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
2.166 0.555
40 PLASTIBOR E41 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
2.207 0.562
40 ESSELTE E41 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
1.806 0.405
Table 3: Variation of silicones vs. sensor sensitivity
Tabela 3: Vpliv silikona na odzivnost senzorjev
RuDPP/mg FOIL SILICONE SOLVENT V/μL STAGES k R2
40 DATALINE E4 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
6.468 0.528
40 DATALINE E41 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
2.166 0.555
40 PLASTIBOR E4 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
6.654 0.524
40 PLASTIBOR E41 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
2.207 0.562
40 ESSELTE E4 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
7.003 0.524
40 ESSELTE E41 CHLO 200
1750 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
1.806 0.405
60 DATALINE E4 ME 150
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
2.699 0.978
60 DATALINE E41 ME 150
1600 r/min –› 3 s
2300 r/min –› 3 s
3150 r/min –› 4 s
2.152 0.991
roughness was made by adding inert silicones to the
sensor solution. Table 3 presents the conditions for pre-
paring the sensor solutions using different silicones. The
optimum linearity (0.9914) was obtained using silicone
E41, foil Dataline and 60 mg of RuDPP. The silicone
adhesion of the sensor solution to the foils was signifi-
cantly improved, while preserving the optimum light
transparency.
3.3 Spin coating and sensor thickness
A lot of factors participated in the spin-coating
sensor preparation33,34 – one of these was also the selec-
tion of a suitable solvent. The sensors were prepared
using chloroform, toluene, and methyl ethyl ketone.
Toluene as a solvent proved to be unsuitable because it
partially dissolved the surfaces of the foils. Other
solvents were constantly evaporating, also during the
spin-coating period, but the sensors prepared with
chloroform presented better characteristics (linearity and
sensitivity) than the sensors prepared with methyl ethyl
ketone (Figure 6a). Chloroform had a lower evaporating
rate than methyl ethyl ketone and the sensors prepared
with chloroform had a uniform film thickness. Figure 6b
demonstrates that the sensor prepared with chloroform
also had a substantially higher signal (by approximately
30 %), while the other parameters remained constant.
Our further experimental work used different spin-
coating stages (periods) and accelerations. The optimum
results were obtained when a 150 μL sensor solution in
chloroform containing 80 mg of RuDPP with silicone
E41 (foil Dataline) was applied (Figure 6a) under the
following spin-coating conditions:
1st step: 750/700 r/min  3 s
2nd step: 300 r/min  3 s
3rd step: 150 r/min  4 s
This spin-coating technique had several benefits
including a fast process time (only a few seconds) when
using low volumes of reagents. A modification of the
spin speeds or an increase in the spin time allowed a
thin-film preparation (below 1 μm).
Using scanning electron microscopy, the irregula-
rities in the sensor surface were searched. For the SEM
analysis, the sensors with homogeneous surfaces (an
optical selection) and optimum oxygen responses were
selected. An optical selection means that the sensors
with the most uniform coating and without any visible
solid particles or air bubbles were scanned (Figure 7).
The coatings and thicknesses of the sensors were
incompletely uniform throughout the sensor surfaces
varying within the range of (3.5–5.0 ± 0.5) μm. Air
bubbles were visible on individual parts, probably cap-
tured in the sensors during the polymerization step. This
P. BRGLEZ et al.: SPIN-COATING FOR OPTICAL-OXYGEN-SENSOR PREPARATION
Materiali in tehnologije / Materials and technology 48 (2014) 2, 181–188 187
Figure 6: a) Sensor response under optimum conditions (signal/a.u. – arbitrary units vs. concentration of O2 in %), b) sensitivity of the sensor
using different solvents (chloroform and methyil ethyil ketone)
Slika 6: a) Odziv senzorja pri optimalnih pogojih (izmerjen signal (a.u. enote) vs. koncentracija O2 v %), b) vpliv topil (kloroform, metil-etil
keton) na odziv senzorjev
Figure 7: Optical-oxygen-sensor SEM images at 2500-times
Slika 7: SEM-posnetka povr{ine opti~nih senzorjev za kisik pri
2500-kratni pove~avi
could be avoided, to some degree, by implementing a
vacuum chamber over the treated surface. The main
problem regarding the entrapped air bubbles was the
fluctuation of the scattering light causing a lower fluo-
rescence signal – the light was scattered in all directions.
The entrapped air bubbles could also cause a longer
response time – the measuring oxygen molecules can be
trapped within the presented voids.
4 CONCLUSION
The spin-coating technique was studied while used
for the optical-sensor preparation when different parame-
ters directly affected the film thickness and, therefore,
also the sensor response to oxygen. This paper primarily
focuses on the influences of the rotation speed and spin-
ning time on the film thickness, in addition to the accele-
ration, temperature, humidity, viscosity, solvents, silicones,
foils and RuDPP concentration studied. The optimum
results were obtained when 80 mg of RuDPP was
dissolved in chloroform, silicone E41 was added and a
150 μL of sensor solution was applied to the Dataline
foil under the following spin-coating conditions: 1st step:
750/700 r/min for 3 s, 2nd step: 300 r/min for 3 s, and 3rd
step: 150 r/min for 4 s.
Spin coating is an alternative method for a sensor
preparation. It is a very fast, simple method, consuming
low volumes of reagents, but making it difficult to pre-
pare completely homogeneous layers on the whole sen-
sor surface. Therefore, it is suggested to be used for a
laboratory-scale sensor preparation, where the majority
of experimental data could be used later when new
coating methods are researched.
Acknowledgements
The authors would like to thank the Slovenian
Technology Agency (TIA) for the financial support
through Grant P-MR-08/54. The authors acknowledge
the financial support from the Ministry of Education,
Science, Culture and Sport of the Republic of Slovenia
through the contract No. 3211-10-000057 (Centre of
Excellence for Polymer Materials and Technologies).
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Mass Transfer, 53 (2010) 9–10, 1712–1717
39 P. Yimsiri, M. R. Mackley, Chemical Engineering Science, 61 (2006)
11, 3496–3505
P. BRGLEZ et al.: SPIN-COATING FOR OPTICAL-OXYGEN-SENSOR PREPARATION
188 Materiali in tehnologije / Materials and technology 48 (2014) 2, 181–188
F. BAGHERALHASHEMI et al.: ELECTROCHEMICAL SYNTHESIS AND CHARACTERIZATION OF ...
ELECTROCHEMICAL SYNTHESIS AND
CHARACTERIZATION OF POLY O-AMINOPHENOL –
SiO2 NANOCOMPOSITE
ELEKTROKEMIJSKA SINTEZA IN KARAKTERIZACIJA
NANOKOMPOZITA POLI O-AMINOFENOL – SiO2
Fatemeh Bagheralhashemi, Abdollah Omrani, Abbas Ali Rostami, Abbas Emamgholizadeh
Faculty of Chemistry, University of Mazandaran, P. O. Box 453, Babolsar, Iran
gholizadehabbas612@yahoo.com
Prejem rokopisa – received: 2012-10-08; sprejem za objavo – accepted for publication: 2013-06-07
We report on the influence of SiO2 on the electropolymerization of O-aminophenol (OAP). A poly O-aminophenol (POAP)
nanocomposite with different particle sizes was deposited on a glassy carbon (GC) electrode in a solution having OAP 0.01 M
and sulfuric acid 0.5 M using cyclic voltammetry (CV). The surface morphologies of the POAP films were studied using
scanning electron microscopy (SEM). The results indicated that a modified surface, having a high surface coverage and an
improved specific capacitance with semiconducting properties, was obtained. The cyclic voltammetry and electrochemical
impedance spectroscopy (EIS) studies confirmed that the nanocomposite films have a higher capacitance than the pure POAP
films. The presence of SiO2 led to an obvious improvement in the overall electrochemical performance of the GC surface
covered by POAP films.
Keywords: poly (O-aminophenol), glassy carbon electrode, electropolymerization
Poro~ilo obravnava vpliv SiO2 na elektropolimerizacijo O-aminofenola (OAP). Nanokompozit poli O-aminofenol (POAP) z
razli~no velikostjo delcev je bil nanesen na elektrodo iz svetle~ega ogljika (GC) s cikli~no voltametrijo (CV) v raztopini z OAP
0,01 M in `vepleno kislino 0,5 M. Morfologija povr{ine POAP-nanosa je bila pregledana z vrsti~nim elektronskim mikro-
skopom (SEM). Rezultati so pokazali, da ima spremenjena povr{ina visoko pokritost povr{ine, izbolj{ano specifi~no
kapacitivnost in izkazuje polprevodni{ke lastnosti. [tudije cikli~ne voltametrije in elektrokemijske impedan~ne spektroskopije
(EIS) so potrdile, da ima nanokompozitni nanos vi{jo kapacitivnost kot ~isti POAP-nanos. Prisotnost SiO2 je pokazala ob~utno
izbolj{anje splo{nih elektrokemijskih zmogljivosti povr{ine GC, pokrite s POAP-nanosom.
Klju~ne besede: poli (O-aminofenol), elektroda iz svetle~ega ogljika, elektropolimerizacija
1 INTRODUCTION
Nanotechnology is a rapidly growing area of nano-
structured material science because of its huge number
of applications in various fields, such as catalysis,1,2
sensors,3,4 electronics,5 optics,6,7 and medical sciences.8–10
Nanometer silicon dioxide (nano-SiO2) is one of the
most popular nanomaterials that are being employed in
different areas, like industrial manufacturing, packaging,
composite materials, disease labeling, drug delivery,
cancer therapy, and biosensors. The modified silica is
low-price, esuriently oxide with a high chemical and
thermal stability.11 The development of chemically modi-
fied electrodes (CMEs) is an area of great interest. CMEs
can be broadly divided into two main categories: as (a)
surface-modified and (b) bulk-modified electrodes.
Methods of surface modification include adsorption,
covalent bonding, attachment of polymer films, etc.12
Polymer-coated electrodes can be differentiated using
other modification methods, like adsorption and covalent
grafting on the surface. Depending on the surface pro-
perties and the experimental conditions, films with a
different morphology, such as multilayer or monolayer,
will be deposited. A conducting polymer provides an
excellent opportunity for the preparation of composites
having desirable mechanical properties. These include
their applications in storage batteries,13 electrochemical
devices,14,15 light-emitting diodes,16 corrosion inhibi-
tors,17 and sensors.18,19 Understanding the nature of these
polymers is of great importance for developing electro-
chemical devices. Among the conducting polymers,
polyaniline (PANI) and its derivatives have been studied
extensively due to the commercial availability of the
monomer, an easy protocol for synthesis, a well-behaved
electrochemical response, a high environmental stability,
and a high conductivity. In recent years, various compo-
sites of PANI and inorganic compounds have been
fabricated with improved characteristics.20,21 Although
silica is an insulating material, some of its composites
showed conductivity at the level of conducting PANI.22
However, some of silica composites displayed enhanced
conductivity, which may be due to the change in mor-
phology of the conductive films in the hybrid materials.23
Electrochemical polymerization offers the advantage of a
reproducible deposition in terms of film thickness and
loading level, making the immobilization procedure of a
metal-based electrocatalyst very simple and reliable. By
considering this point that the nature of the work-
ing-electrode surface is a key factor in observing the
electrochemical response of a deposited polymeric film,
Materiali in tehnologije / Materials and technology 48 (2014) 2, 189–193 189
UDK 66.017:669.018.95 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)189(2014)
the electropolymerization of an OAP monomer in a
suspension of silica nanoparticles is investigated in the
present research for the first time. The silica nanoparticle
is selected because of its good stability in common
aqueous acidic solutions such as H2SO4, HCl and HClO4.
Accordingly, a poly (O-aminophenol)/SiO2 nanocom-
posite as a novel organic matrix was fabricated and
characterized.
2 EXPERIMENTAL
2.1 Materials
The solvent used in this work was bi-distilled water.
Sulfuric acid (purity 99 % from Merck) was used as a
supporting electrolyte. The OAP (purity 98 % from
Merck) was utilized as the monomer and distilled prior
to be used. Silica nanoparticles (purity 99 % with ave-
rage size of 10 nm and a surface area of 600 m2 g–1) were
purchased from Aldrich and used as an additive. Buffer
solutions with various pHs between 5 and 12 were pre-
pared using O-phosphoric acid and its salts (purity 99 %
from Fluka).
2.2 Instruments
Cyclic voltammetery experiments were carried out
using a Potentiostat/Galvanostat EG&G Model 263
(USA) that was well controlled and operated using M
270 software. Impedance spectroscopy measurements
were performed using a potentiostat/galvanostat (model
Autolab, PGSTAT30, Eco Chemie, Netherlands) with
FRA software. All the measurements were conducted
using a three-electrode cell configuration system. The
utilized three-electrode system was composed of
Ag/AgCl · KCl (saturated) as the reference electrode, a
platinum wire as the auxiliary electrode and GCE as the
working electrode substrate. The surface morphology of
the deposited films was characterized by scanning
electron microscopy (model VEGA-TESCAN). A pH
meter (model 3030, JENWAY) was used to read the pH
of the buffer solutions. The pH values for the solution of
phosphate were adjusted by either a solution of sulfuric
acid or sodium hydroxide depending on the pH needed.
2.3 Electrode modification
Prior to the modification of the GCE, it was polished
with alumina slurries on a polishing cloth to a mirror
finish and then ultrasonically cleaned for 2.0 min in
ethanol. Then, the electrode was rinsed thoroughly with
distilled water.
3 RESULTS AND DISCUSSION
3.1 Electrochemical polymerization
The POAP film and POAP/SiO2 nanocomposite were
prepared at the surface of the GCE in the absence and
presence of nanoparticles SiO2 (w = 2 %) in 0.5 mol L–1
H2SO4/0.1 M OAP using the potentiostatic method at E =
0.6 V.
3.2 Surface morphology
The morphology of POAP film and its blend with
SiO2 nanoparticles was characterized using scanning
electron microscopy (Figure 1). The SEM image of the
pristine POAP film on the GCE (Figure 1a) shows a
composed morphology of elongated globules (densely
smooth film). Figure 1b shows the structure of the nano-
composite (POAP/SiO2) surface. The surface of the
POAP film prepared in the presence of SiO2 nano-
particles exhibited discretely shaped agglomerates due to
the influence of the SiO2 nanoparticles. The clusters of
small SiO2 nanoparticles are not uniformly adhered on
the surface of POAP globules. This structure allows the
electrolyte constituent better access to the interior of the
nanocomposite. Certainly, the high dispersion of SiO2
nanoparticles on the surface of the POAP spheres might
have improved the surface area and the stability of the
F. BAGHERALHASHEMI et al.: ELECTROCHEMICAL SYNTHESIS AND CHARACTERIZATION OF ...
190 Materiali in tehnologije / Materials and technology 48 (2014) 2, 189–193
Figure 1: SEM images of: a) POAP film and b) POAP/SiO2 synthe-
sized in H2SO4 solution 0.5 mol L–1
Slika 1: SEM-posnetka: a) POAP-nanosa in b) POAP/SiO2, sinteti-
ziran v raztopini H2SO4 0,5 mol L–1
nanocomposite. This may be attributed to possible
interactions via the hydrogen bonding between the imine
(-NH) group of POAP and the hydroxyl (-OH) group on
the surface of the nano-silica.
3.3 Cyclic voltammetry of POAP/SiO2 and POAP films
Figure 2 shows 15 consecutive voltammograms,
from 0 V to 1.3 V, to deposit the POAP/SiO2 film in a
phosphate solution at pH = 3 at a scan speed of 0.2 V/s.
By the 5th cycle the anodic peak could not be observed.
When the number of scans increased over 5, an anodic
peak current (Ia) was observed clearly at +0.65 V.
The charge measured by the integration of the anodic
peak area was found to be 12.02 mC cm–2 for the 15th
cycle, and oxidation of the POAP film occurred at
0.65 V. The voltammetric behavior of the POAP film in
the same electrolyte solution and in the same potential
range is shown in Figure 3.
Although a similar polymerization charge was used
to synthesize both of the films (POAP/SiO2 and POAP),
the voltammetric behavior of the POAP film was
significantly different. During the 1st to the 5th cycle, the
cyclic voltammogram did not show any change. The shift
of the anodic peak potential to higher currents was
continued until the 10th cycle. The reached anodic peak
charge for the POAP/SiO2 film after the 15th cycle was
not similar to the charge measured for the pure POAP
film until the same cycle. It seems that a transition from
the so-called POAP(II) structure to the POAP(I) structure
occurred. The changes in the POAP(II) structure during
the voltammetry traces could be related to the nature of
the transporting species like counter ions and/or nano-
particles.24 The increase in the anodic peak current of the
POAP/SiO2 film can be attributed to the re-structuring of
the POAP during cycling, where the polymer forms a
more incompact structure, which required the lower
overall potential to be flattened.
The silanol groups of SiO2 could be adsorbed onto
the GC surface. This changes the interfacial structure and
the property of the GC/electrolyte solution, which
benefits the electropolymerization process. However, the
organic-inorganic interactions between the POAP and
SiO2 involved in the electropolymerization process
pushed the polymer chains and so facilitated the growth
of the polymer chains. The anodic peak current increased
by increasing the number of cycles for both of the sys-
tems, but its values were higher for the nanocomposite
than for the pure POAP system.
3.4 Effect of the scan rate
The effect of the scan rate () was investigated in a
phosphate solution at pH = 3 in order to compare the
electrochemical behavior of the deposited polymeric
films. The results of this investigation are shown in
Figure 4.
F. BAGHERALHASHEMI et al.: ELECTROCHEMICAL SYNTHESIS AND CHARACTERIZATION OF ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 189–193 191
Figure 4: The anodic peak current (Ia) of POAP/SiO2 and POAP films
at various scan rates  = 0.2–1.2 V s–1
Slika 4: Vrh anodnega toka (Ia) pri POAP/SiO2- in POAP-nanosu pri
razli~nih hitrostih skeniranja  = 0,2–1,2 V s–1
Figure 2: Cyclic voltammograms of the POAP/SiO2 nanocomposite in
a phosphate solution at pH = 3, scan rate:  = 0.2 V s–1. The 15 cycles
are shown (Ei = 0 to Ef = +1.5 V).
Slika 2: Cikli~ni voltamogrami nanokompozita POAP/SiO2 v fosfatni
raztopini pri pH = 3, hitrost skeniranja  = 0,2 V s–1. Prikazanih je 15
ciklov (Ei = 0 do Ef = +1,5 V).
Figure 3: Cyclic voltammograms of the POAP in a phosphate solution
at pH = 3, scan rate:  = 0.2 V s–1. The 15 cycles are shown (Ei = 0 to
Ef = +1.5 V).
Slika 3: Cikli~ni voltamogrami POAP v fosfatni raztopini pri pH = 3,
hitrost skeniranja  = 0,2 V s–1. Prikazanih je 15 ciklov (Ei = 0 do Ef =
+1,5 V).
The oxidation current clearly increased as the scan
rate increased. An increase in the scan rate is likely due
to an enhancement of the electron flow. It appears that
the increased collision of electrons resulted in a
reduction in the velocity of the electrons leading to a
saturation of the current. The oxidation current for the
POAP/SiO2 system increases linearly with the square of
the scan rate. The above result indicated that the redox
process was confined to the surface of the GCE,
confirming the immobilized state of the POAP/SiO2. The
differences in the redox currents reflect the effective
active surface areas that are accessible to the electrolytes
for the POAP/SiO2. It seems that the porous POAP/SiO2
film has a higher effective surface area. This improve-
ment of the doping/undoping rate is a result of the
increase in the surface area and the porous structure,
which are of benefit to the ion diffusion and migration.
3.5 The influence of pH on the electropolymerization
of POAP
It is well known that the electrochemical process
involving aniline-type polymers requires the exchange of
electrons and protons.25 Accordingly, the solution pH has
a significant effect on the electrochemical behavior of
the polymeric films. Figure 5 shows the cyclic voltam-
mograms of POAP films deposited at various pH values.
The results indicated that the current intensity increased
as the pH of the electrolyte solution became more acidic.
From this it can be implied that the doping/undoping
rates in acidic media are more facilitated. In the other
situation the POAP film becomes more electro-inactive
in basic media.
3.6 EIS measurements
EIS experiments were conducted to provide some
insights into the electrode/polymeric film/electrolyte
interface. Figure 6 displays typical impedance spectra of
the POAP/SiO2 and pure POAP in a phosphate solution
0.1 M recorded at a dc potential of 0.65 V for the
frequency range 40 mHz to 60 kHz. The semicircle
obtained from the high-frequency region was ascribed to
the blocking properties of a single electrode, which can
be related to the faradic process of an ion exchange that
is extremely slow at the polymer/electrolyte interface.
The charge-transfer resistance (Rct) for the POAP/
SiO2 was determined to be 2 × 102 k cm–2, which is
smaller than the value found for the pure POAP (2.52 ×
103 k cm–2). In other words, the pure POAP film pre-
sents a higher electrochemical charge-transfer resistance
than the nanocomposite film. This could be assigned to
the compact structure of the chains in POAP film and its
less active sites for faradic reactions. These results were
also confirmed by the cyclic voltammetry observations,
where the peak current values increased in the presence
of silica nanoparticles.
4 CONCLUSIONS
In this work, a new nanocomposite based on OAP
was prepared by the electropolymerization of OAP at the
surface of GCE, as a low-cost substrate, in the presence
of SiO2 nanoparticles. Apart from the higher electropoly-
merization rate, the POAP/SiO2 nanocomposite showed
good electrochemical behavior, which can be due to the
different morphology of the POAP in the presence silica
nanoparticles. A discrete agglomerated morphology of
the prepared composite was observed owing to the influ-
ence of the SiO2 nanoparticles. The nanoparticles can
reduce the possible interactions between the POAP
chains and, hence, have an important role in the
dynamics of the polymeric chain. Impedance spectro-
scopy results confirmed that the POAP film is more
resistive toward charge transfer than the POAP/SiO2
nanocomposite. Therefore, the presence of SiO2 nano-
F. BAGHERALHASHEMI et al.: ELECTROCHEMICAL SYNTHESIS AND CHARACTERIZATION OF ...
192 Materiali in tehnologije / Materials and technology 48 (2014) 2, 189–193
Figure 6: Impedance profiles of the produced films in phosphate
electrolyte solution 0.1 M for: a) the POAP and b) the POAP/SiO2
nanocomposite
Slika 6: Profil impedance nanosov v raztopini fosfata 0,1 M za: a)
POAP- in b) POAP/SiO2-nanokompozit
Figure 5: The anodic current (Ia) of POAP in different buffer solu-
tions of O-phosphoric acid at pH = (5, 7, 9 and 10)
Slika 5: Anodni tok (Ia) POAP v razli~nih puferskih raztopinah O-fos-
forne kisline pri pH = (5, 7, 9 in 10)
particles results in the improved conductivity of the
POAP films.
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F. BAGHERALHASHEMI et al.: ELECTROCHEMICAL SYNTHESIS AND CHARACTERIZATION OF ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 189–193 193

J. SHIN et al.: DEVELOPMENT OF LOW-Si ALUMINUM CASTING ALLOYS ...
DEVELOPMENT OF LOW-Si ALUMINUM CASTING
ALLOYS WITH AN IMPROVED THERMAL
CONDUCTIVITY
RAZVOJ ALUMINIJEVE LIVARSKE ZLITINE Z MAJHNO
VSEBNOSTJO Si IN IZBOLJ[ANO TOPLOTNO PREVODNOSTJO
Jesik Shin, Sehyun Ko, Kitae Kim
Production Technology R/D Div., Korea Institute of Industrial Technology, Incheon, South Korea
jsshin@kitech.re.kr
Prejem rokopisa – received: 2012-10-12; sprejem za objavo – accepted for publication: 2013-05-28
To develop an aluminum alloy that can combine a high thermal conductivity with a good castability and anodizability, low
Si-containing aluminum alloys, Al-(0.5–1.5)Mg-1Fe-0.5Si and Al-(1.0–1.5)Si-1Fe-1Zn alloys were assessed as potential
candidates. The developed alloys exhibited a thermal conductivity of 170–190 % level (160–180 W/(m K)), a fluidity of 60–85
% level, and an equal or higher ultimate tensile strength compared to those of an ADC12 alloy. In each developed alloy system,
the thermal conductivity and the strength decreased and increased, respectively, as the content of the major alloying elements,
Mg and Si, increased. The fluidity was inversely proportional to the Mg content and directly proportional to the Si content. The
Al-(1.0–1.5)Si-1Fe-1Zn alloys showed better thin-wall castability due to their lower surface energy. In the experimental
aluminum alloys with a low Si content, the fluidity was mainly dependent on the melt surface energy, the Al dendrite coherency
point (DCP), and the first intermetallic crystallization point (FICP), rather than on the solidification interval, latent heat, or the
viscosity.
Keywords: alloy design, low-Si aluminum casting alloy, thermal conductivity, castability
Za razvoj aluminijeve zlitine, ki zdru`uje dobro toplotno prevodnost z dobro livnostjo in mo`nostjo eloksiranja, se predvideva,
da sta dobro izhodi{~e aluminijevi zlitini z majhno vsebnostjo Si, kot sta Al-(0,5–1,5)Mg-1Fe-0,5Si in Al-(1,0–1,5)Si-1Fe-1Zn.
Razvite zlitine so pokazale, da je toplotna prevodnost med 170–190 % (160–180 W/(m K)), livnost med 60–85 % in enaka
natezna trdnost, kot jo ima primerjalna zlitina ADC12. V vsakem razvitem sistemu je toplotna prevodnost nara{~ala in trdnost
padala, ko je nara{~ala vsebnost glavnih legirnih elementov Mg in Si. Teko~nost je bila obratno sorazmerna z vsebnostjo Mg in
sorazmerna z vsebnostjo Si. Zlitina Al-(1,0–1,5)Si-1Fe-1Zn je pokazala bolj{o livno sposobnost pri tankih stenah zaradi manj{e
povr{inske energije. Pri preizkusnih aluminijevih zlitinah z majhno vsebnostjo Si je bila teko~nost predvsem odvisna od
povr{inske energije taline, koheren~ne to~ke Al-dendritov (DCP) in prve to~ke strjevanja (FICP) intermetalne zlitine in manj od
intervala strjevanja, latentne toplote ali viskoznosti.
Klju~ne besede: oblikovanje zlitine, aluminijeva livna zlitina z majhno vsebnostjo Si, toplotna prevodnost, livna sposobnost
1 INTRODUCTION
As the amount of heat that needs to be removed from
electric devices such as LED lighting increases rapidly
with the tendency towards higher outputs, the develop-
ment of heat-dissipating components has recently
become a subject of special interest. Aluminum, the most
common heat-sink material, has inherent disadvantages
that need to be overcome. Although high-purity alumi-
num possesses excellent thermal conductivity, it is
extremely difficult to diecast; thus, alloying elements
must be added, despite the thermal conductivity loss that
occurs as a result of adding these alloying elements. The
ADC12 alloy, a commercial Al–Si-based aluminum
alloy, has been the most common aluminum alloy for
heat-sinks. A heat-sink with a three-dimensional com-
plex shape that is favorable to heat dissipation can be
fabricated in a net shape, with a high productivity and
without the cost penalty that comes with using a high-
pressure diecasting process, like with the ADC12 alloy.
However, a low thermal conductivity below 100 W/(m
K) and the poor anodizing characteristics of the ADC12
alloy, caused by its high Si content, are becoming serious
problems with the increasing power requirements of
electric devices. Other commercial aluminum alloys are
also difficult to diecast or exhibit a conductivity that is
too low for them to be used as heat-dissipating compo-
nents for high-power electric devices.1–4
Therefore, the aim of this study was to develop a
novel, low-Si-containing anodizable aluminum alloy that
possesses both a good thermal conductivity and casta-
bility. To achieve this goal, the elements known to have
advantages in improving castability, strengthening the
matrix, and preventing die sticking as well as to have a
minimal effect on the resistivity increment of aluminum,
such as Mg, Si, Zn, and Fe, were chosen and alloyed.
The amounts of the total alloying elements and Si were
kept between mass fractions 2 % and 3.5 %, and below
1.5 %, respectively. The thermal conductivity, fluidity,
and mechanical strength of the newly designed Al-xMg-
1Fe-0.5Si and Al-xSi-1Fe-1Zn alloys were investigated
as functions of the Mg and Si content and compared to
those of the ADC12 alloy.
Materiali in tehnologije / Materials and technology 48 (2014) 2, 195–202 195
UDK 669.715:621.74.046 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)195(2014)
2 EXPERIMENTAL
2.1 Alloy design
To achieve the combination of a high thermal
conductivity and good castability and anodizability, two
low-Si quaternary aluminum alloys were designed as
follows. First, in terms of the effects of the alloying
elements on the electrical resistivity,5 energy release for
solidification,6 and viscosity5,7 of aluminum, as shown in
Table 1, Mg and Si were chosen as the major alloying
elements. The energy release for solidification was
obtained by summing the latent heat and the superheat-
ing energy calculated for a superheat of 100 °C using the
specific heat of each element according to the simple
mixture rule. A low electrical resistivity and viscosity,
and a high-energy release for the solidification are pre-
ferred because the decreasing electrical resistivity tends
to increase the thermal conductivity, while increasing the
energy release for solidification and decreasing the
viscosity, both of which tend to increase the require-
ments for good castability, such as melt fluidity. The
elements marked with an asterisk (*) in Table 1 are those
that are considered favorable to thermal conductivity and
castability. Thus, it can be seen that only two elements,
Mg and Si, meet all the required conditions. Secondly,
Fe was included to prevent mold-sticking problems.
Lastly, Si and Zn were utilized as supplementary alloy-
ing elements, as described below.
In Table 2, the constituent elements of the two low-Si
quaternary aluminum alloy systems, Al-xMg-Fe-Si
(Alloy 1 series) and Al-xSi-Fe-Zn (Alloy 2 series), and
their chemical compositions are summarized. The total
amounts of alloying elements were kept between mass
fractions 2 % and 3.5 % to achieve a good balance of
thermal conductivity and castability. The amounts of Mg
and Si were varied from 0.5 % to 1.5 % in order to syste-
matically investigate the effects of these major alloying
elements on the thermal conductivity and castability. If
the alloying levels are too high or too low, the thermal
conductivity and castability may deteriorate, respecti-
vely. Because the Si particles decrease the anodizability,
the level of Si in particular was kept below 1.5 %, which
is the maximum amount of Si in commercial wrought Al
alloys known to have good anodizability. The amount of
Fe was 1 %, the same as in the ADC12 alloy. In the
Al-xMg-Fe-Si alloys, 0.5 % Si was added as a supple-
mentary element to effectively increase the energy
196 Materiali in tehnologije / Materials and technology 48 (2014) 2, 195–202
J. SHIN et al.: DEVELOPMENT OF LOW-Si ALUMINUM CASTING ALLOYS ...
Table 1: Effects of alloying elements on electrical resistivity,5 energy release for solidification,6 and viscosity of aluminum5,7 (calculation was
made for 	T of 100 °C)
Tabela 1: Vpliv legirnih elementov na elektri~no upornost,5 energijo, spro{~eno pri strjevanju,6 in viskoznost aluminija5,7 (izra~un je bil izdelan
za 	T 100 °C)
Element
Resistivity Energy release for solidification
Viscosity
variation of Al
with alloying
Maximum
solubility in Al
(w/%)
Resistivity increment of Al per
w/% (μ cm) Latent heat,H of pure ele-
ments (kJ/kg)
Specific heat,
c’ of pure ele-
ments (kJ/kg)
H + c’	T in-
crement of Al
per w/%
(kJ/kg)In solution Out of solution
Cr 0.77 4.00 0.180 402 0.66 –0.3 (+)
Cu 5.65 *0.34 0.030 205 0.45 –2.5 (+)
Fe *0.05 2.56 0.058 272 0.78 –1.5 (+)
Li 4.00 3.31 0.680 422 4.46 *3.7
Mg 14.90 *0.54 0.220 362 1.34 *0.0 *(–)
Mn 1.82 2.94 0.340 268 0.70 –1.6 (+)
Ni *0.05 *0.81 0.061 292 0.56 –1.5 (+)
Si 1.65 *1.02 0.088 1804 0.93 *14.0 *(–)
Ti 1.00 2.88 0.120 366 0.68 –0.6 (+)
V 0.50 3.58 0.280 329 0.62 –1.1
Zn 82.80 *0.09 0.023 111 0.48 –3.4 *(0)
Zr 0.28 1.74 0.044 212 0.37 –2.5
Table 2: Chemical composition in mass fractions (w/%) of the developed low-Si aluminum alloys Al-xMg-Fe-Si (alloy 1 series) and
Al-xSi-Fe-Zn (alloy 2 series)
Tabela 2: Kemijska sestava v masnih dele`ih (w/%) razvitih aluminijevih zlitin z majhno vsebnostjo Si, Al-xMg-Fe-Si (zlitina 1. serije) in
Al-xSi-Fe-Zn (zlitina 2. serije)
Alloy
Major Element Anti-die-stickingelement Supplementary element Base element
Thermal
conductivity
(W/(m K))Mg Si Fe Zn Si Al
1
1-1 0.5 – 1.0 – 0.5 98.0 186
1-2 1.0 – 1.0 – 0.5 97.5 175
1-3 1.5 – 1.0 – 0.5 97.0 160
2
2-1 – 1.0 1.0 1.0 – 97.0 171
2-2 – 1.2 1.0 1.0 – 96.8 163
2-3 – 1.5 1.0 1.0 – 96.5 153
release for solidification. Because Si has been reported to
be less effective in strengthening the Al matrix than Mg,5
1 % Zn was added as a supplementary element in the
Al-xSi-Fe-Zn alloys. Although heat-sinks are not struc-
tural components that need a very high strength, proper
strength is imperative in order to eject the castings from
molds safely. Moreover, Zn has the lowest resistivity
increment, as shown in Table 1. The thermal conducti-
vities of the alloys were calculated by a simple rule-of-
mixture and the Wiedemann-Franz law using the data of
Table 1 in order to predict the effects of the elements on
the resistivity of Al.
2.2 Evaluation and analyses
The fluidity test was conducted using a ceramic-
coated steel mold with multiple channels on a low-
pressure casting machine under an inert-gas atmosphere
in order to prevent melt oxidation. Figure 1a shows the
parting plane of the metal mold for the fluidity test. The
flow channels were 100 mm long and open to the air at
the end, and their diameters were (8, 4, 2, and 1) mm.
Figure 1b shows the fluidity test casting after solidifica-
tion. The average flow lengths for each channel diameter
were taken as the fluidity values, and 10 experiments
were performed to confirm the reproducibility. To eva-
luate the fluidity only as a function of alloy composition,
the mold temperature, the superheat temperature, and the
pressure during pouring were kept constant: 190 °C, 100
°C, and 15 kPa, respectively. The melting points of the
alloys were determined by a thermal analyzer TG/DTA,
model SDT Q600, applying a heating rate of 10 °C/min
up to 700 °C in an Ar atmosphere (flow rate: 0.1 L/min).
The thermal conductivities of the alloys were drawn
from their electrical resistivities as measured by an
eddy-current technique, utilizing the Wiedemann-Franz
law to determine the relation between the electrical resi-
stivity and the thermal conductivity. The tensile strength
evaluation was carried out according to ASTM B 557M
using specimens taken from a Y-block casting with
dimensions of 200 mm × 150 mm × 25 mm. The ADC12
alloy (Al-10 % Si-2.5 % Cu-1 % Fe-0.2 % Mg) was used
as a comparative material for the properties evaluation of
the developed alloys.
Thermophysical modeling using the commercial
software JMatPro 5.0 was performed to obtain the ther-
mophysical properties relating to castability and the
phase-equilibria information. For microstructural analy-
sis, including phase characterization, the cross-sections
of the fluidity test channels were examined using a
field-emission scanning electron microscope (FESEM),
model FEI Quanta 200F, equipped with an energy-
dispersive spectroscopy (EDS) probe. To understand the
solidification paths, cooling curve analyses (CCA) were
carried out based on the two-thermocouple method.8,9
Two K-type thermocouples, one at the center (TC) of the
graphite mold and one near the wall (TW), were
positioned to record the solidification history, as shown
in Figure 2. Then the graphite mold was dipped into the
melt for about 30 s to fill it up and allow its temperature
to equilibrate with the melt temperature. According to
Farahany’s method,8 the temperature at which the first
maximum difference between the thermocouples TC and
TW occurred is regarded as the dendrite coherency point
(DCP).
J. SHIN et al.: DEVELOPMENT OF LOW-Si ALUMINUM CASTING ALLOYS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 195–202 197
Figure 2: Cooling-curve analysis, set-up with a graphite mold and two
K-type thermocouples
Slika 2: Sestav za analizo ohlajevalnih krivulj z grafitno kokilo in
dvema termoelementoma vrste K
Figure 1: a) Parting plane of the metal mold for fluidity test and b)
fluidity test casting
Slika 1: a) Na~rt debeline delov pri preizkusu teko~nosti in b) ulitek
po preizkusu teko~nosti
3 RESULTS AND DISCUSSION
The measured thermal conductivity, fluidity, and
ultimate tensile strength of the developed Al-xMg-
1Fe-0.5Si and Al-xSi-1Fe-1Zn alloys were compared to
those of the ADC12 alloy and the results are summarized
in Figure 3. The developed alloys showed a higher
thermal conductivity of 160–180 W/(m K), which is
approximately 70–90 % higher than that of the ADC12
alloy, as shown in Figure 3a. The thermal conductivity
of the alloys decreased with the increment of Mg and Si,
the major alloying elements. The difference in the
thermal conductivity between the two alloy systems was
fairly insignificant. The average flow lengths for the
diameter channels 1 mm and 2 mm are summarized in
Figure 3b. The channels larger than 2 mm in diameter
were completely filled for both alloy melts. The fluidity
of the developed alloys reached 60–85 % that of the
ADC12 alloy, which was obtained by averaging the flow
length for the diameter channels 2 mm. It was found that
the fluidity of the Al-xMg-1Fe-0.5Si alloys decreased
with increasing Mg content, whereas that of the Al-xSi-
1Fe-1Zn alloys increased with increasing Si content.
Interestingly, the opposite tendency was observed
between the two alloy systems with a changing channel
diameter, i.e., for the diameter channels 2 mm, the
Al-xMg-1Fe-0.5Si alloys showed a higher fluidity than
the Al-xSi-1Fe-1Zn alloys, but this tendency was
reversed for the diameter channels 1 mm. The tensile
strength of the two alloy systems increased with
increasing Mg and Si contents, reaching an equal or
higher tensile strength than the ADC12 alloy, as shown
in Figure 3c. The strength of the Al-xMg-1Fe-0.5Si
alloys was higher than that of the Al-xSi-1Fe-1Zn alloys.
To understand the fluidity behavior, which is more
complex than understanding the thermal conductivity or
strength, important thermophysical properties relating to
the castability were calculated by JMatPro software and
are summarized in Figure 4. Figure 4a shows the vari-
ation of the energy release for solidification as a function
of the contents of Mg and Si. From a higher perspective,
the two alloy systems showed similar levels of energy
release for solidification. However, looking within each
alloy system, the energy release of solidification for the
Al-xMg-1Fe-0.5Si alloys did not change until reaching a
mass fraction 1.0 % Mg and then it decreased. Consider-
ing that Mg is an element with almost the same energy
release for solidification as Al, the solidification path and
crystallization phases seemed to change as the Mg con-
tent increased from 1.0 % to 1.5 %. In the Al-xSi-
1Fe-1Zn alloys, the energy release for solidification
increased in proportion to the amount of Si. To summa-
J. SHIN et al.: DEVELOPMENT OF LOW-Si ALUMINUM CASTING ALLOYS ...
198 Materiali in tehnologije / Materials and technology 48 (2014) 2, 195–202
Figure 4: a) Heat release for solidification, b) viscosity and c) surface energy of Al-(0.5–1.5)Mg-1Fe-0.5Si and Al-(1.0–1.5)Si-1Fe-1Zn alloys
calculated by JMatPro
Slika 4: a) Spro{~anje toplote pri strjevanju, b) viskoznost in c) povr{inska energija zlitin Al-(0,5–1,5)Mg-1Fe-0,5Si in Al-(1,0– 1,5)Si-1Fe-1Zn,
izra~unana z JMatPro
Figure 3: Measured: a) thermal conductivity, b) fluidity and c) ultimate tensile strength of Al-(0.5–1.5)Mg-1Fe-0.5Si and Al-(1.0–
1.5)Si-1Fe-1Zn alloys
Slika 3: Izmerjena: a) toplotna prevodnost, b) teko~nost in c) natezna trdnost zlitin Al-(0,5–1,5)Mg-1Fe-0,5Si in Al-(1,0–1,5)Si-1Fe-1Zn
rize, the variation of energy release for the solidification
coincided relatively well with the fluidity behavior, but
the variation in the quantity of energy release for soli-
dification was too small to fully explain the fluidity
variation. Figure 4b shows the viscosity of the deve-
loped alloys calculated under the same superheat con-
ditions. For both alloy systems, the viscosity increased in
proportion to the amount of Mg and Si. It seemed that in
this investigation, the viscosity did not have a significant
effect on the fluidity of these experimental alloys. For
reference, the reason why the viscosity decreased with
increasing Mg and Si content in the work of other
researchers5,7 was that the viscosity was obtained not for
a constant superheat condition but for a constant pour-
ing-temperature condition, i.e., the variation of the
liquidus temperature of the various alloy compositions
was not considered. On the other hand, it is notable that
the difference in the surface energy between the two
alloy systems was significant, as shown in Figure 4c. It
seemed that the superior fluidity for the small diameter
channel of the Al-xSi-1Fe-1Zn alloys was because of the
lower backpressure caused by a lower surface energy. It
is apparent that the Al-xSi-1Fe-1Zn alloys showed a
similar energy release during the solidification and even
a higher viscosity, but their surface energies were nearly
half those of the Al-xMg-1Fe-0.5Si alloys. Actually, in
this fluidity test experiment, which used a low-pressure
casting machine, the backpressure in the diameter chan-
nels 1 mm calculated by the capillary effect amounted to
approximately 80 % of the applied pouring pressure.
Because the solidification range and secondary
phases also have an influence on the fluidity of alloy
melts,10–15 the phase-equilibria calculations were carried
out prior to the microstructural examination. Figure 5
shows the phase equilibria during solidification calcu-
lated by the Scheil equation (the nonequilibrium lever
rule) using JMatPro. The Al-xMg-1Fe-0.5Si alloys were
typified by the fact that the amount of Al3Fe phase,
which has been reported to have a plate shape,11
increased with increasing Mg content, whereas the soli-
dification range decreased significantly as a result of the
change of the quaternary eutectic point. In the Al-xSi-
1Fe-1Zn alloys, the liquidus temperature decreased
slightly with Si content but the quaternary eutectic point
did not change, resulting in a slight decrease in the
solidification range. In addition, the amount of Al3Fe
phase, which was formed by the first monovariant
eutectic reaction, decreased with increasing Si content.
At 1.5 % Si, the amount of the Al3Fe phase became less
than the amount of the -AlFeSi phase, which was
formed by the subsequent monovariant eutectic reaction.
From these calculated solidification characteristics, it is
thought that the Al3Fe phase played a role in lessening
the fluidity of the Al-xMg-1Fe-0.5Si and Al-xSi-
1Fe-1Zn alloys. The decrease in the solidification range
in the Al-xMg-1Fe-0.5Si alloys might also have pro-
voked the negative effect of the Al3Fe phase on the
fluidity, which is different from the general tendency of a
narrower solidification range improving the melt flui-
dity.10 In other words, the decrease in the solidification
range resulted in an increase in the residual liquid
fraction at the stage when the Al3Fe phase started to
form. As the Mg increased from 0.5 % to 1.5 %, the
residual liquid fraction increased from 50 % to 60 %,
thus enlarging the Al3Fe phase with its large aspect ratio
and finally obstructing the melt flow.
To investigate the actual nonequilibrium solidifica-
tion characteristics, cooling-curve analyses and micro-
scopic tests were carried out. In Figure 6, the cooling-
curve analyses results, which consist of cooling curves,
J. SHIN et al.: DEVELOPMENT OF LOW-Si ALUMINUM CASTING ALLOYS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 195–202 199
Figure 5: Phase equilibria calculated by JMatPro: a) Al-0.5Mg-1Fe-0.5Si, b) Al-1.5Mg-1Fe-0.5Si, c) Al-1.0Si-1Fe-1Zn and d) Al-1.5Si-
1Fe-1Zn alloys
Slika 5: Fazna ravnote`ja v zlitinah, izra~unana z JMatPro: a) Al-0,5Mg-1Fe-0,5Si, b) Al-1,5Mg-1Fe-0,5Si, c) Al-1,0Si-1Fe-1Zn in d) Al-1,5Si-
1Fe-1Zn
their first derivatives, and baselines, and the scanning
electron microscopy images with backscattered electrons
[SEM(BSE)] in the as-cast state are shown. In the
Al-xMg-1Fe-0.5Si alloys (Figures 6a and 6b), three
thermal events are observed in the first-derivative curves.
At the onset of solidification of any phase, the first
derivative increases in value and decreases upon the
completion of solidification.8 It appears that the peaks
labeled (a), (b), and (c) correspond to the crystallization
of primary Al dendrites, the formation of secondary
phases at the inter-dendritic regions between the secon-
dary dendrite arms, and the final eutectic solidification
reaction, respectively. From the results of the combined
SEM (BSE) and EDS analyses, which are shown in
Table 3, it was verified that in the Al-0.5Mg-1Fe-0.5Si
alloy (Figure 6a) the plate-shaped particle (1) between
the secondary dendrite arms was a -AlFeSi phase and
the irregular-shaped particle (2) at the final solidification
regions was an -AlFeSi phase. According to Rosefort et
al.,16 - and -AlFeSi phases in as-cast aluminum alloys
can be identified with the help of SEM-EDS, i.e., the
phase with the plate-like structure at high Si(w/%)/
Fe(w/%) ratios above approximately 0.4 is a -AlFeSi
phase and the phase with the Chinese-script-structure at
low Si(w/%)/Fe(w/%) ratios is an -AlFeSi phase. In the
high-Mg-containing Al-1.5Mg-1Fe-0.5Si alloy (Figure
6b), a dotted bright particle (3) and a dark particle (4)
were observed at the inter-dendritic regions between the
secondary dendrite arms. These dotted bright and dark
particles proved to be -AlFeSi and Mg2Si phases,
respectively. The Mg2Si phase was also observed at the
final solidification regions [particle (5)]. It is likely that
the peak (c) on the first derivative curve of Figure 6b,
which is larger than that of Figure 6a, was related to a
larger latent-heat evolution due to Mg2Si phase
crystallization. In the Al-xSi-1Fe-1Zn alloys, four
thermal events were observed on the first-derivative
curves, as shown in Figures 6c and 6d, and it was
determined that the peaks (a), (b), (c), and (d)
corresponded to the crystallization of primary Al
dendrites, the -AlFeSi phase at the inter-dendritic
regions between secondary dendrite arms, and -AlFeSi
and Si phases at the final solidification regions,
respectively. The results indicate that in the Al-1Si-
1Fe-1Zn alloys (Figure 6c) the peak (c) tended to be
relatively larger than the peak (d). The reverse findings
in the Al-1.5Si-1Fe-1Zn alloys (Figure 6d) were attri-
buted to the microstructural differences in the final
solidification regions with Si content. That is, in a
low-Si-containing Al-1Si-1Fe-1Zn alloy, an -AlFeSi
phase was mostly observed in the final solidification
regions, but in a high-Si-containing Al-1.5Si-1Fe-1Zn
alloy, the Si and -AlFeSi phases prevailed.
Table 3: EDS analysis of the compositions of the particles (1)–(10)
marked in Figure 6
Tabela 3: EDS-analiza sestave delcev (1)–(10), ozna~enih na sliki 6
Point
Mg Si Fe Zn
Al
w/% x/% w/% x/% w/% x/% w/% x/%
(1) 0.75 0.88 5.03 5.11 10.46 5.35 – – bal.
(2) 0.43 0.51 1.55 1.58 10.97 5.62 – – bal.
(3) 1.82 2.16 1.12 1.15 12.86 6.65 – – bal.
(4) 6.79 7.54 6.42 6.17 1.01 0.49 – – bal.
(5) 8.05 8.92 4.63 4.44 1.17 0.56 – – bal.
(6) 1.93 2.32 1.38 1.44 14.96 7.82 – – bal.
(7) – – 7.42 7.56 10.54 5.40 0.16 0.19 bal.
(8) – – 4.95 5.16 12.92 6.78 1.78 0.80 bal.
(9) – – 9.45 9.71 9.91 5.12 1.71 0.76 bal.
(10) – – 17.15 17.11 3.02 1.51 2.46 1.05 bal.
Unlike the calculated phase equilibria during solidifi-
cation, shown in Figure 5, an Al3Fe phase (or the meta-
stable Al6Fe phase) was not observed in the metallo-
graphic inspection for any of the Al-xMg-1Fe-0.5Si and
Al-xSi-1Fe-1Zn alloys, as shown in Figure 6. Moreover,
it is interesting to note that -AlFeSi, a low-tempera-
ture-stable phase, crystallized earlier than -AlFeSi, a
high-temperature-stable phase, at the inter-dendritic
J. SHIN et al.: DEVELOPMENT OF LOW-Si ALUMINUM CASTING ALLOYS ...
200 Materiali in tehnologije / Materials and technology 48 (2014) 2, 195–202
Figure 6: Cooling-curve analysis results and SEM (BSE) microstruc-
tural images in as-cast state: a) Al-0.5Mg-1Fe-0.5Si, b) Al-1.5Mg-
1Fe-0.5Si, c) Al-1.0Si-1Fe-1Zn and d) Al-1.5Si-1Fe-1Zn alloys
Slika 6: Rezultati analize ohlajevalnih krivulj in SEM- (BSE)-posnet-
ki mikrostruktur zlitin v litem stanju: a) Al-0,5Mg-1Fe-0,5Si, b)
Al-1,5Mg-1Fe-0,5Si, c) Al-1,0Si-1Fe-1Zn in d) Al-1,5Si-1Fe-1Zn
regions between the secondary dendrite arms. Except for
the high-Mg-containing alloy (the Al-1.5Mg-1Fe-0.5Si
alloy), the -AlFeSi phase was observed only in the final
solidification regions. These solidification characteristics
could be attributed to the segregation behavior of the
solute atoms, particularly Si. That is because the cooling
rate was high and the solidification occurred as a fine,
mushy type, the Si atoms ejected into the inter-dendritic
liquid had little chance of diffusing out to a residual
liquid reservoir because of the limited time and space.
As a result, the -AlFeSi phase with a high Si/Fe ratio
was crystallized at interdendritic regions between the
secondary dendrite arms rather than the -AlFeSi. Thus,
for the final solidification regions, which had relatively
longer solidification times and lower Si concentrations
than the inter-dendritic regions, an -AlFeSi phase could
crystallize. These are consistent with Darvishi’s experi-
mental results13 regarding the presence of silicon-sub-
stituted Al3Fe and Al6Fe phases with - and -AlFeSi
phases and with Dutta’s simulation results,17 which show
that the increase in the cooling rate increased the Si/Fe
ratio in the eutectic liquid. In the high-Mg-containing
alloy (Al-1.5Mg-1Fe-0.5Si alloy) it is likely that the
consumption of Si atoms due to the formation of a Mg2Si
phase enabled the -AlFeSi phase to crystallize, even at
the inter-dendritic regions between the secondary den-
drite arms under the same cooling conditions. Hosseini-
far et al.18 reported a similar phase selection example in a
6xxx Al alloy: the addition of La resulted in the
formation of a La(Al,Si)2 phase and a decrease of the
Si/Fe ratio in the eutectic liquid, favoring the crystalli-
zation of an -AlFeSi phase rather than a -AlFeSi
phase.
The characteristic solidification parameters, includ-
ing the solidification range, were calculated from the
cooling and first-derivative curves and the results are
summarized in Table 4. However, it seems that the
Al-xMg-1Fe-0.5Si and Al-xSi-1Fe-1Zn alloys did not
follow the general relation between the solidification
range and the fluidity length of the alloy very well, in
which the fluidity length is inversely proportional to the
solidification range. DCP and the first intermetallic
crystallization point (FICP) also did not seem to show a
direct relation with the fluidity length. The onset
temperatures for the (b) peaks in Figure 6, which were
determined by the tangential line method, were used as
the FICP.
Table 4: Characteristic solidification parameters for Al-xMg-Fe-Si
(alloy 1 series) and Al-xSi-Fe-Zn (alloy 2 series) alloys
Tabela 4: Zna~ilni parametri strjevanja za Al-xMg-Fe-Si (zlitina 1.
serije) in Al-xSi-Fe-Zn (zlitina 2. serije)
Alloy
Start of
solidifica-
tion (°C)
End of
solidifica-
tion (°C)
Solidifica-
tion range
(°C)
FICP
(°C)
DCP
(°C)
1
1-1 652 571 81 642 640
1-2 649 558 91 637 636
1-3 646 562 84 635 632
2
2-1 651 550 101 637 641
2-2 648 550 98 630 633
2-3 648 546 102 626 634
From the cooling-curve analyses, the latent heat and
liquid fraction at characteristic solidification points can
be obtained. Knowing these values is important in order
to understand the solidification characteristics of the
alloy. The latent heat can be calculated by multiplying
the accumulative area between the first derivative and the
baseline by the specific heat.9 The liquid fraction can be
obtained by calculating the accumulative area between
the first derivative and the baseline from the characte-
ristic solidification points to the solidification end point
as a fraction of the total area between these curves.8
Figure 7 shows the calculated latent heat and the liquid
fractions at the DCP and FICP. The latent heat increased
with increasing Mg and Si in the Al-xMg-1Fe-0.5Si and
Al-xSi-1Fe-1Zn alloys, as shown in Figure 7a. These
results are different from the thermophysical modeling
results of Figure 3a, in which the increase of Mg content
decreased the latent heat of the Al-xMg-1Fe-0.5Si alloys.
This contrasting tendency of the latent heat with
increasing Mg content is responsible for the difference in
the solidification path between the phase equilibria cal-
culation and the actual non-equilibrium solidification. In
any case, it seems that the latent heat did not have a
significant effect on the fluidity of these experimental
alloys. In the liquid fractions at the DCP and FICP of
Figure 7b, two interesting findings were obtained. First,
in the Al-xMg-1Fe-0.5Si alloys, the liquid fraction was
J. SHIN et al.: DEVELOPMENT OF LOW-Si ALUMINUM CASTING ALLOYS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 195–202 201
Figure 7: a) Latent heat and b) liquid fraction at DCP and FICP of
Al-(0.5–1.5)Mg-1Fe-0.5Si and Al-(1.0–1.5)Si-1Fe-1Zn alloys, which
were obtained from a cooling-curve analysis
Slika 7: a) Latentna toplota in b) dele` taline pri DCP- in FICP-zlitin
Al-(0,5–1,5)Mg-1Fe-0,5Si in Al-(1,0–1,5)Si-1Fe-1Zn, dobljeni iz
analize ohlajevalnih krivulj
higher at FICP than at DCP, but this tendency is reversed
in the Al-xSi-1Fe-1Zn alloys. Second, in the Al-xMg-
1Fe-0.5Si alloys, the liquid fractions at DCP and FICP
increased with increasing Mg content, whereas in the
Al-xSi-1Fe-1Zn alloys, the liquid fractions decreased
with increasing Si content. The dendrite coherency
seriously decreases the fluidity.19 Therefore, it is likely
that in the Al-xSi-1Fe-1Zn alloys, although the crystalli-
zation of the plate-like -AlFeSi phase during solidifi-
cation has a negative effect on the fluidity of Al alloys13
and the amount of the -AlFeSi phase increased with
increasing Si content, as shown in Figures 6c and 6d, its
crystallization after DCP had little influence on the
fluidity. It was thus concluded that in the Al-xSi-1Fe-1Zn
alloys, the increase of the Si content effectively
improved the fluidity by increasing the latent heat and
lowering the DCP without disturbing the secondary
phases, whereas in the Al-xMg-1Fe-0.5Si alloys the
secondary phases could have a negative effect on the
fluidity because they crystallized before the dendrite
coherency. Paes14 and Zhang15 reported that the presence
of a Mg2Si phase increased the viscosity of Al alloys.
Considering that the liquid fraction of the Al-1.5Mg-
1Fe-0.5Si alloy was approximately 50 % at FICP, the
increase of the melt viscosity could be regarded as the
most important factor in terms of decreasing the fluidity.
Therefore, it was concluded that in the Al-xMg-
1Fe-0.5Si alloys, the increase of Mg content significantly
deteriorated the fluidity by increasing the viscosity and
increasing the DCP, in spite of the increase in the latent
heat.
4 CONCLUSIONS
For the development of an aluminum alloy that com-
bines a high thermal conductivity with a good castability
and anodizability, the low-Si-containing aluminum
alloys Al-(0.5–1.5)Mg-1Fe-0.5Si and Al-(1.0–1.5)Si-
1Fe-1Zn were investigated. The obtained results are as
follows:
1. The developed aluminum alloys exhibited a thermal
conductivity of 170–190 % level (160–180 W/(m K)), a
fluidity of 60–85 % level, and an equal or higher
ultimate tensile strength compared to those of the
ADC12 alloy.
2. In each developed alloy system, the thermal con-
ductivity decreased and the strength increased with
increasing amounts of Mg and Si, the major alloying
elements. The fluidity exhibited an inverse rela-
tionship with the Mg content and a direct relationship
with the Si content.
3. The contradictory fluidity variation behavior in the
two alloy systems with the compositions of Mg and
Si was caused by the opposing tendencies of DCP
and FICP and the relatively different occurring
sequences of DCP and FICP.
4. In the experimental aluminum alloys with a low Si
content, the prevailing Fe-containing intermetallic
compound and the solidification path were observed
to be mainly dependent on the Si segregation beha-
vior and the Mg alloying level, rather than on the
initial Si/Fe alloying ratio.
5. It was found that the Al-Mg-Fe-Si-based aluminum
alloys that show a higher strength and good fluidity
in channels greater than 2 mm in diameter are
potential materials for general cast heat-dissipating
components, and that the Al-Si-Fe-Zn-based alumi-
num alloys that possess a lower surface energy are
potential materials for thin-wall cast heat-dissipating
components.
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badi, A. Ourdjini, In-situ thermal analysis and macroscopical charac-
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Analysis of A356 Aluminum Alloy, Met. & Mat. Inter., 10 (2004),
89–96
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Trans., 70 (1962), 400–409
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primary Al3Fe phase in Al–Fe alloy under a high magnetic field,
Materials Letters, 61 (2007), 983–986
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of iron-intermetallics on the fluidity of 413 aluminum alloy, Journal
of Alloys and Compounds, 468 (2009), 539–545
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effect of iron and manganese on microstructure and mechanical pro-
perties of aluminium –silicon alloy, MJoM, 16 (2010), 11–24
14 M. Paes, E. J. Zoqui, Semi-solid behavior of new Al–Si–Mg alloys
for thixoforming, Mat. Sci. & Eng., A406 (2005), 63–73
15 J. Zhang, Z. Fan, Y. Wang, B. Zhou, Hypereutectic aluminium alloy
tubes with graded distribution of Mg Si particles prepared by cen-
trifugal casting, Materials and Design, 21 (2000), 149–153
16 M. Rosefort, C. Matthies, H. Buck, H. Koch, Light Metals 2011,
TMS 2011, 711–715
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behaviour of Al–Fe–Si alloys, Mat. Sci. & Eng., A283 (2000),
218–224
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mation during the solidification of an Al-Mg-Si alloy containing La,
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202 Materiali in tehnologije / Materials and technology 48 (2014) 2, 195–202
M. BABI^ et al.: A NEW METHOD FOR ESTIMATING THE HURST EXPONENT H FOR 3D OBJECTS
A NEW METHOD FOR ESTIMATING THE HURST
EXPONENT H FOR 3D OBJECTS
NOVA METODA ZA OCENJEVANJE HURSTOVEGA EKSPONENTA
H ZA 3D-OBJEKTE
Matej Babi~1, Peter Kokol2, Nikola Guid2, Peter Panjan3
1Emo-Orodjarna, d .o. o., Slovenia
2University of Maribor, Faculty of Electrical Engineering and Computer Science, Slovenia
3Jo`ef Stefan Institute, Slovenia
babicster@gmail.com
Prejem rokopisa – received: 2012-10-23; sprejem za objavo – accepted for publication: 2013-06-10
Mathematics and computer science are very useful in many other sciences. We use a mathematical method, fractal geometry, in
engineering, specifically in laser techniques. Characterization of the surface and the interfacial morphology of
robot-laser-hardened material is crucial to understand its properties. The surface microstructure of robot-laser-hardened material
is rough. We aimed to estimate its surface roughness using the Hurst parameter H, which is directly related to the fractal
dimension. We researched how the parameters of the robot-laser cell impact on the surface roughness of the hardened specimen.
The Hurst exponent is understood as the correlation between the random steps X1 and X2, which are followed by time for the
time difference 	t. In our research we understood the Hurst exponent H to be the correlation between the random steps X1 and
X2, which are followed by the space for the space difference 	d. We also have a space component. We made test patterns of a
standard label on the point robot-laser-hardened materials of DIN standard GGG 60, GGG 60 L, GGG 70, GGG 70 L and
1.7225. We wanted to know how the temperature of point robot-laser hardening impacts on the surface roughness. We developed
a new method to estimate the Hurst exponent H of a 3D-object. This method we use to calculate the fractal dimension of a
3D-object with the equation D = 3 – H.
Keywords: fractal structure, Hurst parameter H for 3D-objects, robot, laser, hardening
Matematika in ra~unalni{tvo sta zelo uporabni veji drugih vrst znanosti. Uporabili bomo matemati~no metodo – fraktalno
geometrijo – v in`enirstvu, natan~neje, v laserski tehniki. Karakterizacija povr{ine in morfologija robotsko lasersko kaljenih
materialov ima klju~ni pomen za razumevanje lastnosti materialov. Povr{ina robotsko lasersko kaljenega materiala je hrapava.
To hrapavost pa bi radi ocenili s Hurstovim eksponentom H, ki ga dobimo direktno iz fraktalne geometrije, kot meritev
hrapavosti povr{ine. Raziskali smo, kako parameter robotske celice vpliva na hrapavost kaljenih vzorcev. Hurstov eksponent H
razumemo kot korelacijo med korakoma X1 in X2, ki ga dobimo med ~asovno razliko 	t. V na{i raziskavi razumemo Hurstov
eksponent H kot relacijo med korakoma X1 in X2, ki ga dobimo med prostorsko komponento. Osredinili smo se na to~kovno
robotsko lasersko kaljenje vzorcev standardne oznake po DIN-standardu GGG 60, GGG 60 L, GGG 70, GGG 70 L in 1.7225.
@eleli smo izvedeti, kako temperatura to~kovnega robotskega laserskega kaljenja vpliva na hrapavost povr{ine. Razvili smo
novo metodo za ocenjevanje Hurstovega eksponenta H za 3D-objekte. To metodo smo uporabili za ra~unanje fraktalne
dimenzije 3D-objektov po formuli D = 3 – H.
Klju~ne besede: fraktalna struktura, Hurstov parameter H za 3D-objekte, robot, laser, kaljenje
1 INTRODUCTION
The Hurst parameter1 is understood as the correlation
between the random steps X1 and X2, using space for the
space difference 	d. It occurs in many areas of applied
mathematics, including fractals and chaos theory, and is
used in many fields, ranging from biophysics to network
computers. The parameter was originally developed in
hydrology. However, modern techniques for estimating
the Hurst parameter H are emerging from fractal
mathematics. For example, the fractal dimension was
used to measure the roughness of sea coasts. The
relationship between the fractal dimension D and the
Hurst parameter H is given by the equation D = 2 – H
for 2D-objects and by D = 3 – H for 3D-objects. There is
also a form called statistical self-similarity, where if we
have one data set of a seemingly endless string of data
sets, we can assume that each data set has the same sta-
tistical properties as any other. Statistical self-similarity
occurs in a surprisingly large number of areas in engi-
neering. Fractal geometry2,3 has offered a new perspec-
tive on maths and science and allows observation of the
world around us from a new and completely different
point of view. Nature is full of shapes and images that
from a distance are seen to be similar. A well-known
example of such a self-similar pattern is Sierpinski’s
pyramid. Let us assume that we have a statistically
self-similar finite data series. Any part of this series
would have the same statistical characteristics as the
entire series or any other part of this series. Such series
are often used in engineering. We used one in the
analysis of robot-laser-hardened materials.4 A robot-laser
surface-hardening heat treatment is complementary to
conventional flame or inductive hardening. The energy
source for the laser hardening is a laser beam, which
heats up very quickly. Laser hardening is a process of
controlled energy intake, high performance constancy,
and an accurate positioning process. A hard martensitic
microstructure provides improved surface properties
such as wear resistance and high strength. Point
Materiali in tehnologije / Materials and technology 48 (2014) 2, 203–208 203
UDK 621.785.6 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)203(2014)
robot-laser hardening5–8 is a classic case of robot-laser
hardening. This means that the speed of the laser beam is
no longer a parameter that can be changed. In point
robot-laser hardening cells we are interested in finding
the optimal parameters that give the maximum hardness
of the hardened material. In this work we have used a
scanning electronic microscope (SEM) to search and
analyse the fractal structure of the robot-laser-hardened
material. First, we introduce the Hurst parameter to find
how the temperature of the robot laser cell impacts on
the optimal Hurst parameter H of the hardened material.
We then developed a new method to estimate the Hurst
parameter H of a 3D-object.
2 MATERIALS PREPARATION AND METHOD
2.1 Materials preparation
We made test patterns of a standard label on point
robot-laser-hardened materials of DIN-standard GGG
60, GGG 60 L, GGG 70, GGG 70 L and 1.7225. We
hardened the materials at different temperatures. So we
changed the temperature parameter of the robot laser
cells, T 
 [800, 2000] °C, with a step of 100 °C. In all
the experiments we recorded the microstructure. We
wanted to know how the temperature of point robot-laser
hardening impacts on the surface roughness. Figure 1
shows the transverse and longitudinal section of a
hardened material. Each sample was etched and polished
(IMT, Institute of Metals and Technology Ljubljana,
Slovenia) and then examined under a microscope (IJS,
Jo`ef Stefan Institute). The images were obtained using a
JEOL JMS-7600F field-emission scanning electron
microscopy. Figures 2 to 6 present microstructure of
robot laser hardened specimens. Figure 7 presents boun-
dary between hardened and non-hardened DIN 1.7225
standard material.
M. BABI^ et al.: A NEW METHOD FOR ESTIMATING THE HURST EXPONENT H FOR 3D OBJECTS
204 Materiali in tehnologije / Materials and technology 48 (2014) 2, 203–208
Figure 2: Microstructure of hardened DIN 1.7225 standard material
(SEM)
Slika 2: Mikrostruktura kaljenega materiala po DIN-standardu 1.7225
(SEM)
Figure 1: Transverse and longitudinal section of hardened material
Slika 1: Pre~ni in vzdol`ni prerez kaljenega materiala
Figure 5: Microstructure of hardened DIN GGG 60L standard
material (SEM)
Slika 5: Mikrostruktura kaljenega materiala po DIN-standardu GGG
60 L (SEM)
Figure 4: Microstructure of hardened DIN GGG 70 L standard
material (SEM)
Slika 4: Mikrostruktura kaljenega materiala po DIN-standardu GGG
70 L (SEM)
Figure 3: Microstructure of hardened DIN GGG 70 standard material
(SEM)
Slika 3: Mikrostruktura kaljenega materiala po DIN-standardu GGG
70 (SEM)
The random process is evaluated statistically using
the Hurst parameter H or by determining the distribution
function. The Hurst parameter H as a self-similarity
criterion cannot be accurately calculated; it can only be
estimated. There are several different methods9–13 for
producing estimates of the parameter H, which deviate
from one another to some extent. In doing so, we have
no criteria to determine which method gives the best
result.
Different methods for estimating the Hurst exponent
H have been evaluated.14 The assessment methods in the
space component domain are based on a comparison of
the original process and the average process with the
method of aggregation:
• The variance-time plot analysis is based on the pro-
perty of the slowly decaying variance of self-similar
processes undergoing aggregation.
• The R/S method. The adjusted rescaled range method
or adjusted scale is also a graphical method based on
the properties of the Hurst phenomenon.
• In statistics, residual variance is another name for an
unexplained variation, the sum of squares of diffe-
rences between the y-value of each ordered pair on
the regression line and each corresponding predicted
y-value; it is generally used to calculate the standard
error of an estimate. In other words, residual variance
helps us confirm how well the regression line that we
constructed fits the actual dataset. The smaller the
variance, the more accurate the predictions are.
Methods for evaluation in the frequency or wavelet
("wavelet") space are:
• The periodogram method based on the noise 1/f and
the Fourier transform.
• The Whittle estimator is based on minimizing the
likelihood function used in the periodogram method.
There is no graphical method.
• The Hurst exponent is estimated by using the wavelet
transform of the series. A least-squares fit on the
average of the squares of the wavelet coefficients at
different scales is an estimate of the Hurst exponent.
The method produces both a graphical output and a
confidence interval.
Estimation of the Hurst parameter H15–19 using diffe-
rent methods gives results that show significant differen-
ces.
2.2 Method
Here we present our first new method for estimating
the Hurst exponent H for 3D-objects. First of all, we find
all the coordinates (x, y, z) of an SEM picture. Here, we
use the program ImageJ. Then we use only z coordinates
to estimate the Hurst exponent H. We present all the z
coordinates in a 2D-space component graph (Figure 8),
which is continuous. Also, all the points (xi, y0, zi) pre-
sent the first space component in a 2D-graph for all the
M. BABI^ et al.: A NEW METHOD FOR ESTIMATING THE HURST EXPONENT H FOR 3D OBJECTS
Materiali in tehnologije / Materials and technology 48 (2014) 2, 203–208 205
Figure 7: The boundary between hardened and non-hardened DIN
1.7225 standard material (SEM)
Slika 7: Meja med kaljenim in nekaljenim delom materiala po
DIN-standardu 1.7225 (SEM)
Figure 6: Microstructure of hardened DIN GGG 60 standard material
(SEM)
Slika 6: Mikrostruktura kaljenega materiala po DIN-standardu GGG
60 (SEM)
Figure 9: Long space component
Slika 9: Dolga prostorska komponenta
Figure 8: Space component
Slika 8: Prostorska komponenta
points (xi, zi). All the points (xi, y1, zi) represent the
second space component in a 2D-graph for all the points
(xi, zi). We obtained the space component for all yi, i
(Figure 9). Then we combined all of these space com-
ponents into one space component. For this long space
component we can estimate the Hurst exponent H.
Figure 10 shows the 3D-object of the point robot-laser-
hardened microstructure.
3 RESULTS AND DISCUSSION
The pictures in the .jpeg format were converted into
256-grey-level numerical matrices (level 1 for black and
256 for white) with the program ImageJ. Then we entered
information into the program Selfis 01B, with which we
obtained the following graphs. The graphs were made
only for specimen 1.7225 hardened at 800 °C.
Figures 11a to 11d show the different methods for
estimating the Hurst exponent H.
Figures 12 to 16 present the estimated value of the
Hurst parameter H with five methods for five different
robot-laser-hardened materials.
The collected data were entered into the program and
the graphs were obtained, from which we can deduce the
optimum results.
M. BABI^ et al.: A NEW METHOD FOR ESTIMATING THE HURST EXPONENT H FOR 3D OBJECTS
206 Materiali in tehnologije / Materials and technology 48 (2014) 2, 203–208
Figure 11: a) Variance method for hardened material 1.7225, b) R/S
method for hardened material 1.7225, c) Absolute moments estimator
for hardened material 1.7225, d) Variance of residuals method for
hardened material 1.7225, e) Periodogram method for hardened
material 1.7225
Slika 11: a) Variance metoda za kaljeni material 1.7225, b) R/S-me-
toda za kaljeni material 1.7225, c) Ocena absolutnega momenta za
kaljeni material 1.7225, d) Varian~na metoda ostankov za kaljeni
material 1.7225, e) Metoda periodograma za kaljeni material 1.7225
Figure 10: 3D-object of robot-laser-hardened microstructure
Slika 10: 3D-objekt lasersko kaljene mikrostrukture
Figure 13: Relationship between estimated value of the Hurst para-
meter H and temperature of robot-laser-hardened specimens for DIN
standard GGG 70
Slika 13: Odvisnost med ocenjeno vrednostjo Hurstovega parametra
H in temperaturo robotsko lasersko kaljenih vzorcev GGG 70
Figure 12: Relationship between estimated value of the Hurst para-
meter H and temperature of robot-laser-hardened specimens for DIN
standard 1.7225
Slika 12: Odvisnost med ocenjeno vrednostjo Hurstovega parametra
H in temperaturo robotsko lasersko kaljenih vzorcev 1.7225
Material 1.7225: The smallest Hurst exponent H is
obtained at a temperature of 1400 °C with all methods,
which means that the roughness is higher.
Material GGG 70: The highest Hurst exponent H is
obtained at a temperature of 2000 °C with all methods,
which means that the roughness is lower.
Material GGG 70 L: The smallest Hurst exponent H
is obtained at a temperature of 1400 °C with all methods,
which means that the roughness is higher. It is similar to
the material 1.7225.
Material GGG 60 L: The smallest Hurst exponent H
is obtained at a temperature of 2000 °C with all methods,
without for the periodogram method.
Material GGG 60: The smallest Hurst exponent H is
obtained at a temperature of 800 °C with all methods,
which means that the roughness is higher. The highest
Hurst exponent H is obtained at a temperature of 1400
°C with all methods without for the periodogram me-
thod.
Smaller values are obtained with the periodogram
method for all the hardened specimens at all temperatu-
res. Higher values are obtained with the aggregate
variance estimator method for hardened specimens at all
the temperatures.
4 CONCLUSIONS
We conducted experiments on five different mate-
rials. We changed only the temperature parameter, T 

[800, 2000] °C, in 100 °C steps. In total there were 65
samples. We were interested in which parameter of the
robot laser cell achieved the roughest surface of the
hardened material. We found the Hurst parameter H
exactly. The main findings can be summarized with the
following points:
• A fractal structure exists in robot-laser hardening.
• The Hurst parameter H was calculated using different
methods. The results were compared using the pro-
gram Image J and it was found that the best results
were obtained by the periodogram method, which
gave smaller values.
• With the help of the equation D = 2 – H, the fractal
dimensions can be calculated for all the samples. For
calculating the fractal dimensions in the 2D-plane
obtained by a microscope as a two-dimensional
image from which the fractal dimension can be
determined.
• In the 3D-space we use equation D = 3 – H. In addi-
tion, we found a new process for calculating the
fractal dimension of a 3D object.
• We found the optimal Hurst parameter H for different
laser parameters of robot cells on different materials.
• The smaller the Hurst parameter H, the greater the
surface roughness.
• We developed a new method to estimate the Hurst
parameter H of a 3D-object.
The author of the Selfis program has written a great
deal about the Hurst parameter H. The problem with
these methods is that we estimated only the Hurst para-
meter H and that these methods are based on different
bases (some of aggregation, others on wavelet transfor-
mation, etc.), which in turn lead to different estimates.
Also, each method has certain advantages and limitations
with regard to the captured sample (some methods are
M. BABI^ et al.: A NEW METHOD FOR ESTIMATING THE HURST EXPONENT H FOR 3D OBJECTS
Materiali in tehnologije / Materials and technology 48 (2014) 2, 203–208 207
Figure 14: Relationship between estimated value of the Hurst para-
meter H and temperature of robot-laser-hardened specimens for DIN
standard GGG 70 L
Slika 14: Odvisnost med ocenjeno vrednostjo Hurstovega parametra
H in temperaturo robotsko lasersko kaljenih vzorcev GGG 70 L
Figure 16: Relationship between estimated value of the Hurst
parameter H and temperature of robot-laser-hardened specimens for
DIN-standard GGG 60
Slika 16: Odvisnost med ocenjeno vrednostjo Hurstovega parametra
H in temperaturo robotsko lasersko kaljenih vzorcev GGG 60
Figure 15: Relationship between estimated value of the Hurst
parameter H and temperature of robot-laser-hardened specimens for
DIN standard GGG 60 L
Slika 15: Odvisnost med ocenjeno vrednostjo Hurstovega parametra
H in temperaturo robotsko lasersko kaljenih vzorcev GGG 60 L
better for large samples, some for smaller ones). Many
books state that there is no precise method with which to
accurately calculate the Hurst parameter and that it can
only be estimated. The most common method of calcul-
ating the Hurst parameter H is the R/S-method.
Our findings are important from a practical point of
view. The precise parameters of the robot-laser cell tem-
pering influence the hardness of the hardened material.
Materials with such properties have better wear resistan-
ce and a longer lifespan. The Hurst parameter H can give
us information about the correlation between the rough-
ness and the hardness of material, which is important in
many industries.
Acknowledgements
The present work was supported by the European
Social Fund of the European Union.
5 FUTURE WORK PLAN
In the future we plan to explore the Hurst parameter
H as a function of several parameters in robot cells for
laser hardening. The laser parameters include power,
energy density, focal distance, energy density in the fo-
cus, focal position, temperature, and speed of hardening.
In robot-laser hardening, many different problems are
encountered. We are interested in estimating the Hurst
parameter H in:
• two-beam laser robot hardening (where the laser
beam is divided into two parts);
• areas of overlap (where the laser beam covers an
already hardened area);
• robot-laser hardening at different angles (the angles
change depending on the x and y axes).
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pendence, Birkhauser, 2003, 373–407
14 H. Kettani, J. A. Gubner, A novel approach to the estimation of the
Hurst parameter in self-similar traffic, Proceedings of the 27th
Annual IEEE Conference on Local Computer Networks (LCN 2002),
Tampa, Florida, 2002, 160–165
15 J. M. Bardet, G. Lang, G. Oppenheim, A. Phillipe, S. Stoev, M. S.
Taqqu, Semi-parametric estimation of the long-range dependence
parameter: A survey, In: P. Doukhan, G. Oppenheim, M. S. Taqqu
(eds.), Theory and Applications of Long-Range Dependence, Birk-
hauser, 2003, 557–577
16 J. M. Bardet, G. Lang, G. Oppenheim, A. Phillipe, M. S. Taqqu, Ge-
nerators of long-range dependent processes: A survey, In: P. Douk-
han, G. Oppenheim, M. S. Taqqu (eds.), Theory and Applications of
Long-Range Dependence, Birkhauser, 2003, 579–623
17 C. Park, F. Godtliebsen, M. Taqqu, S. Stoev, J. S. Marron, Visuali-
zation and inference based on wavelet coeficients, SiZer, and SiNos,
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18 G. W. Wornell, A. V. Oppenheim, Estimation of fractal signals from
noisy measurements using wavelets, IEEE Transactions on Signal
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parameter in linear fractional stable motion, Signal Processing, 82
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208 Materiali in tehnologije / Materials and technology 48 (2014) 2, 203–208
M. SHARIFIRAD et al.: DESIGN OF A MICROBIAL SENSOR USING A CONDUCTING POLYMER ...
DESIGN OF A MICROBIAL SENSOR USING A CONDUCTING
POLYMER OF POLYANILINE/POLY 4,4’-DIAMINODIPHENYL
SULPHONE-SILVER NANOCOMPOSITE FILMS ON A CARBON
PASTE ELECTRODE
OBLIKOVANJE MIKROBNEGA SENZORJA Z UPORABO PREVODNE
POLIMERNE POLIANILINSKE/POLI 4,4’-DIAMINODIFENIL SULFONSKE
SREBRNE NANOKOMPOZITNE TANKE PLASTI NA ELEKTRODI Z
OGLJIKOVO PASTO
Meysam Sharifirad, Farhoush Kiani, Fardad Koohyar
Department of Chemistry, Faculty of Science, Ayatollah Amoli Branch, Islamic Azad University, Amol, Iran
f_koohyar@yahoo.com
Prejem rokopisa – received: 2012-11-03; sprejem za objavo – accepted for publication: 2013-06-07
A microbial biosensor based on Gluconobacter oxydans cells immobilized on a conducting polymer of polyaniline/Poly
4,4’-diaminodiphenyl sulphone-silver (PANI/PDDS/Ag) coated onto the surface of a carbon-paste electrode was constructed.
The proposed biosensor was characterized using glucose as the substrate. Conducting polymers are electrochemically
polymerized at the carbon-paste electrode substrates. The polymer films are modified by electrochemically depositing
PANI/PDDS/Ag particles. The effect of changing the size of the Poly 4,4’-diaminodiphenyl sulphone-silver particles and the
polymer film thickness on the voltammetric behaviour of the resulting hybrid material showed noticeable changes in the
electrocatalytic current in an acid medium. The morphology of the polymer films and the distribution of the PDDS/Ag particles
in the film were studied with scanning electron microscopy.
Keywords: polyaniline/Poly 4,4’-diaminodiphenyl sulphone-silver, conducting polymers, gluconobacter oxydans, voltammetric
properties, nanoparticles
Konstruiran je bil mikrobni senzor na osnovi glukonobakterijskih oksidacijskih celic, pritrjenih na prevoden polimer iz
polianilina/poli 4,4’-diaminodifenil sulfon–srebra (PANI/PDDS/Ag), nanesenega na povr{ino elektrode z ogljikovo pasto.
Predlagani senzor je bil karakteriziran z uporabo podlage iz glukoze. Prevodni polimeri so bili elektrokemijsko polimerizirani
na elektrodo s podlago iz ogljikove paste. Polimerne tanke plasti so bile modificirane z elektrokemijskim nanosom
PANI/PDDS/Ag-delcev. Spreminjanje velikosti poli 4,4’-diaminodifenil sulfonskih srebrnih delcev in debeline polimerne plasti
je pokazalo ob~utne spremembe pri voltametri~nem vedenju, rezultirajo~i hibridni material pa je pokazal ob~utne spremembe
elektrokataliti~nega toka v kislem mediju. Morfologija polimernih plasti in razporeditev PDDS/Ag-delcev v plasti sta bila
pregledana na vrsti~nem elektronskem mikroskopu.
Klju~ne besede: polianilin/poli 4,4’-diaminodifenil sulfon–srebro, prevodni polimeri, glukonobakterijski oksidant, voltame-
tri~ne lastnosti, nanodelci
1 INTRODUCTION
Many types of microbial sensors have been deve-
loped as analytical tools. Such a microbial sensor
consists of a transducer and a microbe as the sensing
element. The characteristics of microbial sensors are in
complete contrast to those of enzyme sensors or immu-
nosensors, which are highly specific for the substrates of
interest, although the specificity of the microbial sensor
has been improved by genetic modifications of the
microbe used as the sensing element. Microbial sensors
have the advantages of a tolerance to the measuring
conditions, a long lifetime, and cost performance, but
they also have the disadvantage of a long response time.
Since their discovery in the mid-1970s, research on
conducting polymers (CPs) has become an ever-growing
research area in polymer chemistry.1 The redox
behaviour and an unusual combination of the properties
of metals and plastics make conducting polymers a new
class of materials.2 The interest in conducting polymers
is largely due to the wide range of possible applications
due to their facile synthesis, good environmental stability
and the long-term stability of the electrical conductivity.
The advantage of using conducting polymers compared
to more traditional sputtered metal coatings is that the
polymer is soluble, enabling a non-destructive analysis
of the specimens.3 CPs were extensively studied in the
past decade and used for technological applications in
electrochromic devices,4,5 gas-separation membranes,6
enzyme immobilization7 and have been featured in bio-
technical applications since the very early days following
their discovery. The biosensing approach using CPs has
also been widely investigated in previous studies.8
There have been several attempts to produce nano-
particle polymer composites. Overall, we note four diffe-
rent approaches used to date. The first technique consists
of the in-situ preparation of the nanoparticles in the
polymer matrix. This is affected by a reduction of the
metal salts dissolved in the polymer matrix.9 The second
technique involves polymerizing the matrix around the
Materiali in tehnologije / Materials and technology 48 (2014) 2, 209–214 209
UDK 678.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)209(2014)
nanoparticles.10 The blending of pre-formed nanoparti-
cles into pre-synthesized polymer can be considered as
the third technique for the preparation of a nanoparticle
polymer composite material.11 The fourth process
involves the in-situ synthesis of the composite material,
with the metal nanoparticles being formed from an ionic
precursor and the polymer being produced from the
monomer. A major advantage of the latter method lies in
the provision of a better particle-polymer interaction.12
Microbial cells are very promising for biosensor
constructions due to them having several advantages: the
enzyme does not need to be isolated, the enzymes are
usually more stable in their natural environment in the
cell, the co-enzymes and activators are already present in
the system.13 Cell-based biosensors are frequently used
for a determination of the biological oxygen demand
(BOD), the toxic agents and the assimilatable sugars as
well as the selective detection of a single analyte.14
In this study different amounts of polymer nanocom-
posites aniline were examined for the detection of
glucose and ethanol. The measurement was based on the
respiratory activity of the cells. As well as the optimiza-
tion and characterization, an application of the proposed
system was carried out on real samples.
2 EXPERIMENTAL
2.1 Reagents
LiClO4, NaClO4, dialysis membrane, AlCl3, succinyl
chloride, benzene-1,4-diamine propionic acid, nitrome-
thane, iron(III) chloride, propylene carbonate, poly(me-
thylmethacrylate), dichloromethane, toluene, d-glucose,
ethanol and gold colloidal were purchased from Sigma.
Methanol and acetonitrile were purchased from Merck.
All other chemicals were of analytical grade and pur-
chased either from Merck or Sigma. The DDS were
purchased from Merck, and silver nitrate from Alderich.
2.2 Apparatus
Scanning electron microscopy (SEM) images were
taken using a VEGA HV (high potential) 1500 V at
various magnifications. All the experiments were carried
out in air atmosphere at room temperature (25 °C). A
conventional three-electrode set up was employed for the
CV studies. The counter electrode was a platinum sheet
with a surface area of 2 cm2. A SCE electrode was em-
ployed as the reference electrode. Electropolymerization
and all other electrochemical studies were carried out
using a Potentiostat/Galvanostat EG&G Model 263 A;
USA well equipped with M 270 software.
2.3 Synthesis of the composite material
In a typical reaction, 0.8 g DDS was dissolved in a
magnetically stirred 25 mL of methanol in a 50 mL
conical flask. After complete dissolution, 100 mL dilute
silver nitrate (10–1 mol dm–3) was added drop wise. After
the addition of the entire silver nitrate, the precipitated
material collected at the bottom of the flask. This work
accomplished for the solution of silver nitrate that the
colloidal solutions were observed as milky coloured.
2.4 Cultivation of G. oxydans
The strain of G. oxydans was obtained from DSMZ
and maintained on agar containing (g L–1): d-glucose, 90;
yeast extract, 15; calcium carbonate, 18; agar, 25.15 The
cell biomass was prepared by aerobic cultivation at 28
°C. Then, the cells were collected by centrifugation after
reaching the late exponential phase and washed twice
with 0.9 % sodium chloride solution containing CaCl2 2
mM. The biomass concentration was expressed as the
weight matter determined by drying to a constant weight
at 105 °C.
2.5 Microbial electrode method of making
Prior to the electropolymerization, graphite elec-
trodes were polished on wet emery paper and washed
thoroughly with distilled water, sonicated for 2–3 min,
rinsed with bi-distilled water and dried at 105 °C. The
electrochemical polymerization of DDS/Ag (1 mg/mL)
was carried out by scanning the potential between -0.5 V
and 0.8 V via cyclic voltammetry with the scan rate of
300 mV/s in NaClO4 (0.2 M) and LiClO4 (0.2 M)/ace-
tonitrile medium. The concentration of the monomer was
as for the polymerization of the DDS/Ag.
For the cells of bacteria to prove G. oxydans on
electrode coated with polymer, first we spray it and leave
the electrode surface to dry for 1 h. After the removal of
unbound cells by washing with distilled water, the layer
was covered with a dialysis membrane, pre-soaked in
water. The membrane was fixed tightly with a silicone
rubber O-ring. Daily-prepared electrodes with fresh cells
were used in all the experimental steps, unless otherwise
stated. Control experiments using carbon paste electrode
were covered with bacterial cells and a polymer coating
that was not.
2.6 Measurements
All the measurements were carried out at 30 °C under
constant stirring. After each run, the electrode was
washed with distilled water and kept in a phosphate
buffer (pH 7) solution 50 mM at 30 °C for 7 min. The
working buffer solution was renewed after each measure-
ment. The microbial sensor was initially equilibrated in
the buffer and then the substrate was added to the
reaction cell. After 30 min the substrate was added to the
reaction cell. The biosensor responses were registered as
current densities by following the oxygen consumption
at –0.7 V due to the metabolic activity of the bacterial
cells in the presence of glucose.16
M. SHARIFIRAD et al.: DESIGN OF A MICROBIAL SENSOR USING A CONDUCTING POLYMER ...
210 Materiali in tehnologije / Materials and technology 48 (2014) 2, 209–214
3 RESULTS AND DISCUSSION
Enzymes and cells have been used in biosensor
construction for many decades. Both concepts have some
advantages and challenges.17 There have been various
strategies to modify the microbial cells for applications
in microbial biosensors. The principle of the bacterial
biosensor is rather simple, and sensor productions only
require the growth of the microorganisms. There are
multiple strategies for how to use catalytic activities
present in microbial cells ranging from using viable
cells, non-viable cells, permeable cells, or membrane
fractions. These cell-derived biocatalysts serve as an
economical substitute for enzymes; an additional benefit
for the biosensor performance is that the enzymes are
still linked to the respiratory chain.
Since conducting polymers can act as transducers in
biosensors and coating electrodes with CPs under mild
conditions this opens up various possibilities for the
immobilization of biomolecules and biosensing mate-
rials, the enhancement of their electrocatalytic proper-
ties, rapid electron transfer and direct communication.
The co-immobilization of redox mediators or cofactors
by entrapment within electropolymerized films or by
covalent binding on the surface allows the simple
fabrication of reagent-less biosensors.18 The CPs have an
organized molecular structure on different transducers,
which allows them to function as a three-dimensional
matrix for the biomolecule immobilization and preserve
the activity for a long period. This property of the matrix
with their functionality as a membrane has provided
opportunities for sensor development. This process is
reproducible with a high operational stability.19 In this
work the use of an electrochemically polymerized anili-
ne/4,4’-diaminodiphenyl sulphone-silver as a microbial
biosensing platform was examined for G. oxydans cells.
The morphologies of the bacterial sensing surfaces
provide the most precise information about the cells and
matrices used in the system. The scanning electron
microscopy (SEM) technique is utilized to monitor the
surface characteristics and shows the interaction between
the biological materials and the immobilization matrices.
The morphologies of bare graphite, and a polymeric
matrix with and without cells were shown in Figure 1.
Unbound cells were removed by washing the electrode
surfaces several times before analysis. As seen from the
micrographs, PANI/DDS/Ag provided an efficient immo-
bilization platform with a compact structure for the cell
immobilization. Hence, cells could be kept on the
surface where higher sensor responses with high
operational stabilities are obtained. The presence of
amino groups in the structure may also contribute to
attaching the micro-organisms on the matrix due to the
ionic interactions between the cell surface and this
functional group.
3.1 Effect of electropolymerization time
The most convenient electrochemical method
employed for characterization is cyclic voltammetry.
Cyclic voltammograms of bare graphite electrode carbon
paste electrode (A), PDDS/Ag polymer (B), bacteria –
PDDS/Ag (C), bacteria – PANI/PDDS/Ag (D) onto the
graphite electrode are shown in Figure 2. The amount of
conductive polymer on the electrode surface can be
controlled by adjusting the polymerization time, which
has a direct effect on the resulting current values. It has
been previously reported that the microstructure of a
conducting polythiophene film changes with an increase
in the thickness. As the thickness depends on the
deposition time, more and more defects such as voids
and large molecule agglomerates could appear, causing
degradation and an incompact microstructure of the
films.20 In order to observe the effect of electropolyme-
rization time, aniline/4,4’-diaminodiphenyl sulphone-
silver was polymerized onto the carbon paste surface for
different periods ((5, 10 and 20) min, Figure 3), and then
modified electrodes were used to form microbial biosen-
M. SHARIFIRAD et al.: DESIGN OF A MICROBIAL SENSOR USING A CONDUCTING POLYMER ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 209–214 211
Figure 1: SEM images of bare graphite surface: a) PANI/PDDS/Ag
polymer, b) bacteria-PANI/PDDS/Ag
Slika 1: SEM-posnetka gole povr{ine grafita: a) polimer PANI/
PDDS/Ag, b) bakterija PANI/PDDS/Ag
a)
b)
sors, as described in the experimental section. The best
current values were obtained for 10 min of polymeri-
zation time. However, a decrease was observed when the
polymerization time was higher. This could be due to the
improper film structure related to the thickness after 10
min of the electropolymerization time for the cell immo-
bilization.
3.2 Effect of cell amount
In order to determine the appropriate amount of cells,
different biosensors containing (5, 10, 15, 20, 25, 30, 35
and 40) μL of bacterial cells were prepared. The highest
current responses were obtained with the 25 μL of cells.
When the 5 μL of cells was used the lowest current
response was obtained. On the other hand, when the
amount of cells was increased to 30 μL, a lower signal
than that for 25 μL was obtained. This is an expected
result and caused by the diffusion problem due to the
high cell density. Since both amounts caused lower
current values, further experiments were conducted using
the 25 μL cell (Figure 4).
3.3 Effect of pH
The effect of pH on the microbial sensor based on
PANI/PDDS/Ag was achieved by adjusting the pH
between 6.0 and 8 when using phosphate buffer (50
mM). The response of the microbial sensor towards
glucose (10 mM) at different pH values was shown in
Figure 5. Since pH 7 has the maximum current response,
it was chosen as the optimum pH and all the other
experiments were conducted with this pH value.
M. SHARIFIRAD et al.: DESIGN OF A MICROBIAL SENSOR USING A CONDUCTING POLYMER ...
212 Materiali in tehnologije / Materials and technology 48 (2014) 2, 209–214
Figure 3: Effect of electropolymerization time on the biosensor
response (in phosphate buffer, 50 mM, pH 7, 30 °C, –0.7 V)
Slika 3: Vpliv ~asa elektropolimerizacije na odzivnost biosenzorja (v
fosfatnem pufru, 50 mM, pH 7, 30 °C, –0,7 V)
Figure 5: Effect of pH on the biosensor response (pH 6.0–8.0 phos-
phate buffer, 30 °C, –0.7 V)
Slika 5: Vpliv pH na odzivnost biosenzorja (pH 6,0–8,0 fosfatni pufer,
30 °C, –0,7 V)
Figure 4: Effect of cell loading on the biosensor response (in po-
tassium phosphate buffer (50 mM, pH 7), –0.7 V, 10 mM glucose)
Slika 4: Vpliv obremenitve celice na odziv biosenzorja (v kalijevem
fosfatnem pufru (50 mM, pH 7), –0,7 V, glukoza 10 mM)
Figure 2: Cyclic voltammograms of bare carbon paste electrode (A),
PDDS/Ag polymer (B), bacteria – PDDS/Ag (C), bacteria – PANI/
PDDS/Ag (D) onto the graphite electrode (number of scans: 100, in
potassium phosphate buffer (50 mM, pH 7))
Slika 2: Kro`ni voltamogrami elektrode z ogljikovo pasto (A),
polimera PDDS/Ag (B), bakterije – PDDS/Ag (C), bakterije – PANI/
PDDS/Ag (D) na grafitni elektrodi ({tevilo pregledovanj: 100, v kalij
fosfatnem pufru (50 mM, pH 7))
3.4 Effect of temperature
The amperometric response of the microbial sensors
to glucose (50 mM) was followed at different tem-
peratures, varying from 20 °C to 40 °C. From 20 °C to
25 °C an increase was observed up to 30 °C and then the
signal started to decrease at 35 °C (Figure 6). As a
result, the optimum temperature was found to be appro-
ximately 30 °C. And further experiments were conducted
at this temperature.
3.5 Analytical characteristics
A microbial sensor was prepared, as described in the
experimental part, to examine the analytical characteri-
stics. The linear dynamic ranges and the equations were
obtained based on optimized conditions. For the pro-
posed system, a linear calibration graph Figure 7 (A)
was obtained for the current density versus the substrate
concentration between 0.1 mM and 2.5 mM glucose. A
linear relation was defined with the equation y = 0.3116x
(R2 = 0.9376), where y is the sensor response in the
current density (μA/cm2) and x is the substrate concen-
tration (mM).
Another type of electrode was covered using a
polymer PDDS/silver and bacteria. It was previously
reported that metal nanoparticles can display unique
advantages, such as an enhancement of the mass trans-
port, catalysis, a high effective surface area and control
over the electrode microenvironment over macro-electro-
des when used for electro-analysis. For instance, Pt and
Au nanoparticles are very effective as matrices for
enzyme sensors by taking advantage of the biocompati-
bility and large surface area. In our case, the calibration
curve for the modified system based on silver nano-
particles was studied using the same method as described
in the experimental section. However, these nanoparti-
cles in polymer nanocomposites were a DDS. A linear
relation for the glucose substrate was found between 0. 5
mM and 3 mM, as represented by the equation y =
0.8774x (R2 = 0.9786) and a response time of 120 s
(Figure 7 (B)). It is also possible that the high surface
area due to the Ag on the polymer matrix provides the
loading of the largest amount of cells, causing a larger
biosensor response. The effects of different nanoparticles
in terms of sensitivity and stability in microbial
biosensing are now under investigation.
At the end of the calibration curve the carbon-paste
electrode modified with bacteria and polymer PANI/
DDS/Ag was observed (Figure 7 (C)). The currents
recorded were low in this case and irreproducible current
responses were observed. A linear relation was observed
in the range 0.1–3.0 mM glucose and defined by the
equation y = 1.056x (R² = 0.9808). This reveals that the
polymer provides a good contact for the cells on the
electrode surface where they can attach and survive
during the operational conditions, as described in our
previous study.
4 CONCLUSIONS
Conducting polymers that have an organized
molecular structure can serve as proper and functional
immobilizing platforms for biomolecules. These matri-
ces provide a suitable environment for the immobili-
zation and preserve the activity for long durations. In this
paper a measurement method and environment for the
bacteria and the PANI/DDS/Ag-modified electrode are
described. The proposed system does not require any
complicated immobilization procedure for the construc-
tion of a biosensor. The preparation is simple, cheap and
is not time consuming. The biosensor showed a good
linear range, a good repeatability and a high operational
M. SHARIFIRAD et al.: DESIGN OF A MICROBIAL SENSOR USING A CONDUCTING POLYMER ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 209–214 213
Figure 6: Effect of temperature on the electrode response for micro-
bial biosensors (in potassium phosphate buffer, 50 mM, pH 7, –0.7 V)
Slika 6: Vpliv temperature na odzivnost elektrode pri osnovnem
mikrobnem biosenzorju (v kalijevem fosfatnem pufru, 50 mM, pH 7,
–0,7 V)
Figure 7: Calibration curves for three types of microbial biosensor
including bacteria immobilized on the bare carbon paste electrode (A),
PDDS/Ag polymer (B), PANI/PDDS/Ag (C) (in potassium phosphate
buffer (50 mM, pH 6.5), –0.7 V)
Slika 7: Umeritvene krivulje za tri vrste mikrobnih biosenzorjev,
vklju~no z bakterijami, pritrjenimi na elektrodi z ogljikovo pasto (A),
polimer PDDS/Ag (B), PANI/PDDS/Ag (C) (v kalijevem fosfatnem
pufru (50 mM, pH 6,5), –0,7 V)
stability. It can be concluded that the proposed system
could also be a good alternative for BOD and toxicity
estimation.
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M. SHARIFIRAD et al.: DESIGN OF A MICROBIAL SENSOR USING A CONDUCTING POLYMER ...
214 Materiali in tehnologije / Materials and technology 48 (2014) 2, 209–214
I. B. RISTESKI: A NEW GENERALIZED ALGEBRA FOR THE BALANCING OF  CHEMICAL REACTIONS
A NEW GENERALIZED ALGEBRA FOR THE
BALANCING OF  CHEMICAL REACTIONS
NOVA POSPLO[ENA ALGEBRA ZA URAVNOTE@ENJE
KEMIJSKIH REAKCIJ 
Ice B. Risteski
2 Milepost Place # 606, Toronto, Ontario, Canada M4H 1C7
ice@scientist.com
Prejem rokopisa – received: 2012-12-01; sprejem za objavo – accepted for publication: 2013-07-15
In this article we develop a new generalized algebra for the balancing of  chemical reactions. This is a completely new
approach to the balancing of these kinds of chemical reactions that is based on an understanding of reaction analysis and the
elementary theory of inequalities. The generators of the reaction determined all the interactions among the stoichiometric
coefficients.
Keywords:  chemical reactions, generalized algebra, balancing reactions
V tem ~lanku smo razvili novo posplo{eno algebro za uravnote`enje kemijskih reakcij. To je popolnoma nov na~in uravno-
te`enja kemijskih reakcij, ki temelji na navidezni analizi reakcij in elementarni teoriji neenakosti. Generatorji reakcij dolo~ajo
vse interakcije med stehiometri~nimi koeficienti.
Klju~ne besede: kemijske reakcije , posplo{ena algebra, uravnote`enje reakcij
1 INTRODUCTION
Since the balancing of chemical reactions in che-
mistry is a basic and fundamental issue it deserves to be
considered on a satisfactory level. This topic always
draws the attention of students and teachers, but it is
never a finished product. Because of its importance in
chemistry and mathematics, there are several articles
devoted to the subject. However, here we will not
provide a historical perspective about this topic, because
it has been done in so many previous publications. We
can, however, still provide a full balancing of chemical
reactions with the use of a generalized algebra.
In mathematics and chemistry there are several
mathematical methods for balancing chemical reac-
tions.1–7 All of them are based on generalized matrix
inverses and they have formal scientific properties that
need a higher level of mathematical knowledge for their
application. The so-called chemical methods are parado-
xical and out of order.8
The newest approach for balancing  reactions is
developed in9. The present article is a prolongation of the
previous research.9,10
Generally speaking, balancing a chemical reaction
that possesses atoms with fractional oxidation numbers
is a tough problem in chemistry. It is really hard for
reactions that have only one set of coefficients, but for 
reactions that have an infinite number of sets of
coefficients, this problem is extremely hard.11,12
In the next section we shall consider three general
reactions of oxidation. They are examples of elementary
 reactions, which possess atoms with fractional and
integer oxidation numbers. Actually, we balanced three
general  reactions with one, two and three arbitrary
elements. The first reaction plays a very important role
in metallurgy. For instance, this reaction is a basic reac-
tion in the theory of metal corrosion, ferrous metallurgy
as well as the theory of metallurgical processes, but
unfortunately it was not taken into account until today.
The main reason why this reaction was neglected lies in
its balancing. This article will provide its full balancing,
which is neither easy nor simple.
2 MAIN RESULTS
Now we shall consider the announced  reactions.
Reaction 1. Let us consider this general  reaction
with one arbitrary element:
x1 X + x2 O2  x3 X0.987O + x4 X2O3 + x5 X3O4 (1)
The above chemical reaction (1) reduces to the
following system of linear equations:
x1 = 0.987x3 + 2x4 + 3x5,
2x2 = x3 + 3x4 + 4x5 (2)
Since the system (2) has two linear equations and five
unknowns, we can solve it in 5!/[2!(5 – 2)!] = 10 ways.
Actually, we shall determine all the possible general
solutions of the system (2). They are the following pairs:
(x1, x2), (x1, x3), (x1, x4), (x1, x5), (x2, x3), (x2, x4), (x2, x5),
(x3, x4), (x3, x5) and (x4, x5).
1° Let x3, x4 and x5 be arbitrary real numbers, then the
general solution of the system (2) is:
Materiali in tehnologije / Materials and technology 48 (2014) 2, 215–219 215
UDK 54 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)215(2014)
x1 = 0.987x3 + 2x4 + 3x5,
x2 = x3/2 + 3x4/2 + 2x5 (3)
After the substitution of (3) into (1), the balanced
reaction takes on this general form:
(0.987x3 + 2x4 + 3x5) X + (x3/2 + 3x4/2 + 2x5) O2
 x3 X0.987O + x4 X2O3 + x5 X3O4, (4)
 x3, x4, x5 
 .
This means that by finding the coefficients of the
products we find the coefficients of the reactants.
2° Assume x2, x4 and x5 are arbitrary real numbers,
then the general solution of the system (2) is:
x1 = 1.977x2 – 0.961x4 – 0.948x5,
x3 = 2x2 – 3x4 – 4x5 (5)
The balanced reaction (1) obtains this general form:
(1.977x2 – 0.961x4 – 0.948x5) X + x2 O2
 (2x2 – 3x4 – 4x5) X0.987O + x4 X2O3 + x5 X3O4 (6)
Since the generators (5) are positive, it should imme-
diately follow these inequalities:
1.977x2 – 0.961x4 – 0.948x5 > 0,
2x2 – 3x4 – 4x5 > 0 (7)
From (7) we obtain the inequality:
x2 > 1.5x4 + 2x5 (8)
The expression (8) is a necessary and sufficient con-
dition for a general reaction (6) to hold. In other words,
the reaction is possible if and only if the condition (8) is
satisfied.
3° Suppose x2, x3 and x5 are arbitrary real numbers.
The general solution of the system (2) is:
x1 = 4x2/3 + 0.961x3/3 + x5/3,
x4 = 2x2/3 – x3/3 – 4x5/3 (9)
If we substitute (9) into (1), the general form of the
balanced reaction is:
(4x2/3 + 0.961x3/3 + x5/3) X + x2 O2
 x3 X0.987O + (2x2/3 – x3/3 – 4x5/3) X2O3+ x5 X3O4 (10)
Since the generator x4 > 0, it should immediately
follow that:
x2 > 0.5x3 + 2x5 (11)
The reaction (10) is possible if and only if the con-
dition (11) is satisfied. The inequality (11) is a necessary
and sufficient condition to hold (10).
4° Let x2, x3 and x4 be arbitrary real numbers. The ge-
neral solution of the system (2) is:
x1 = 3x2/2 + 0.948x3/4 – x4/4,
x5 = x2/2 – x3/4 – 3x4/4 (12)
After the substitution of (12) into (1), the general
chemical reaction takes on this form:
(3x2/2 + 0.948x3/4 – x4/4) X + x2 O2
 x3 X0.987O + x4 X2O3
+ (x2/2 – x3/4 – 3x4/4) X3O4 (13)
Since the generators x1, x5 > 0, then it must be:
3x2/2 + 0.948x3/4 – x4/4 >0,
x2/2 – x3/4 – 3x4/4 > 0 (14)
After (14) it immediately follows that:
x2 > x3/2 + 3x4/2 (15)
The inequality (15) is a necessary and sufficient con-
dition to hold the general reaction (13), i.e., the reaction
(13) holds if and only if (15) is satisfied.
5° Suppose x1, x4 and x5 are arbitrary real numbers.
The general solution of the system (2) is:
x2 = x1/1.974 + 0.961x4/1.974 + 0.948x5/1.974,
x3 = x1/0.987 – 2x4/0.987 – 3x5/0.987 (16)
The balanced chemical reaction (1) obtains this
general form:
x1 X + (x1/1.974 + 0.961x4/1.974 + 0.948x5/1.974) O2
 (x1/0.987 – 2x4/0.987 – 3x5/0.987) X0.987O
+ x4 X2O3 + x5 X3O4 (17)
Since the generator x3 > 0, then it must be:
x1 > 2x4 + 3x5 (18)
The above inequality (18) is a necessary and suffi-
cient condition to hold (17), i.e., the reaction (17) holds
if and only if (18) is satisfied.
6° Assume x1, x3 and x5 are arbitrary real numbers.
The general solution of the system (2) is:
x2 = 3x1/4 – 0.961x3/4 – x5/4,
x4 = x1/2 – 0.987x3/2 – 3x5/2 (19)
The balanced chemical reaction (1) has this general
form:
x1 X + (3x1/4 – 0.961x3/4 – x5/4) O2
 x3 X0.987O + (x1/2 – 0.987x3/2
– 3x5/2) X2O3 + x5 X3O4 (20)
Since the generators x2, x4 > 0, then it must be that:
3x1 – 0.961x3 – x5 > 0,
x1 – 0.987x3 – 3x5 > 0 (21)
From (21) we obtain:
x1 > 0.987x3 + 3x5 (22)
The above expression (22) is a necessary and suffi-
cient condition to hold the general reaction (20), i.e., the
reaction (20) holds if and only if (22) is satisfied.
7° Assume x1, x3 and x4 are arbitrary real numbers.
The general solution of the system (2) is:
x2 = 2x1/3 – 0.474x3/3 + x4/6,
x5 = x1/3 – 0.987x3/3 – 2x4/3 (23)
The balanced chemical reaction (1) has this general
form:
x1 X + (2x1/3 – 0.474x3/3 + x4/6) O2
 x3 X0.987O + x4 X2O3
+ (x1/3 – 0.987x3/3 – 2x4/3) X3O4 (24)
Since the generators x2, x5 > 0, then these inequalities
should follow:
I. B. RISTESKI: A NEW GENERALIZED ALGEBRA FOR THE BALANCING OF  CHEMICAL REACTIONS
216 Materiali in tehnologije / Materials and technology 48 (2014) 2, 215–219
4x1 – 0.948x3 + x4 > 0,
x1 – 0.987x3 – 2x4 > 0 (25)
From (25) we obtain:
x1 > 0.987x3 + 2x4 (26)
The inequality (26) is a necessary and sufficient con-
dition to hold the general reaction (24), i.e., the reaction
(24) holds if and only if (26) is satisfied.
8° Let us assume x1, x2 and x5 are arbitrary real num-
bers. The general solution of the system (2) is:
x3 = 3x1/0.961 – 4x2/0.961 – x5/0.961,
x4 = – x1/0.961 + 1.974x2/0.961
– 0.948x5/0.961 (27)
After the substitution of (27) into (1), the general
chemical reaction obtains this form:
x1 X + x2 O2  (3x1/0.961 – 4x2/0.961
– x5/0.961) X0.987O + (– x1/0.961 + 1.974x2/0.961
– 0.948x5/0.961) X2O3 + x5 X3O4 (28)
Since the generators x3, x4 > 0, then these inequalities
should follow:
3x1 – 4x2 – x5 > 0,
– x1 + 1.974x2 – 0.948x5 > 0 (29)
From (29) we obtain:
4x2/3 + x5/3 < x1 < 1.974x2 – 0.948x5,
x2 > 2x5 (30)
The inequalities (30) are necessary and sufficient
conditions to hold the general reaction (28). In other
words, the reaction (28) holds if and only if (30) are
satisfied.
9° Suppose x1, x2 and x4 are arbitrary real numbers.
The general solution of the system (2) is:
x3 = 4x1/0.948 – 6x2/0.948 + x4/0.948,
x5 = – x1/0.948 + 1.974x2/0.948 – 0.961x4/0.948 (31)
The balanced chemical reaction (1) has this general
form:
x1 X + x2 O2  (4x1/0.948 – 6x2/0.948
– x4/0.948) X0.987O + x4 X2O3 + (– x1/0.948
+ 1.974x2/0.948 – 0.961x4/0.948) X3O4 (32)
Since the generators x3, x5 > 0, then these inequalities
should follow:
4x1 – 6x2 – x4 > 0,
– x1 + 1.974x2 – 0.961x4 > 0 (33)
From (33) we obtain:
3x2/2 – x4/4 < x1 < 1.974x2 – 0.961x4,
x2 > 3x4/2 (34)
The inequalities (34) are necessary and sufficient
conditions to hold the general reaction (32), i.e., the reac-
tion (32) holds if and only if (34) are satisfied.
10° Let us assume x1, x2 and x3 are arbitrary real num-
bers. The general solution of the system (2) is:
x4 = – 4x1 + 6x2 + 0.948x3,
x5 = 3x1 – 4x2 – 0.961x3 (35)
The balanced chemical reaction (1) has this general
form:
x1 X + x2 O2  x3 X0.987O
+ (– 4x1 + 6x2 + 0.948x3) X2O3
+ (3x1 – 4x2 – 0.961x3) X3O4 (36)
Since the generators x4, x5 > 0, then follow these
inequalities:
– 4x1 + 6x2 + 0.948x3 > 0,
3x1 – 4x2 – 0.961x3 > 0 (37)
From (37) we obtain:
4x2/3 + 0.961x3/3 < x1 < 3x2/2 + 0.474x3/2,
x2 > x3/2 (38)
The inequalities (38) are necessary and sufficient
conditions to hold the general reaction (36), i.e., the reac-
tion (36) holds if and only if (38) are satisfied.
Example 1. For instance, if we substitute X = Fe, Mn,
Pb in (1), we immediately obtain three sub-particular 
balanced reactions.
Next, we shall consider the following two  reac-
tions.
Reaction 2. Let us balance this general  reaction
with two arbitrary elements:
x1 X2 + x2 Y2 + x3 O2
 x4 XYO + x5 XYO2 + x6 XYO3 (39)
The above chemical reaction (39) reduces to the fol-
lowing system of linear equations:
2x1 = x4 + x5 + x6,
2x2 = x4 + x5 + x6,
2x3 = x4 + 2x5 + 3x6 (40)
Since the system (40) has three linear equations and
six unknowns, we can solve it in 6!/[3!(6 – 3)!] = 20
ways. Actually, we must determine all the possible gene-
ral solutions of the system (40). They are the following
triads: (x1, x2, x3), (x1, x2, x4), (x1, x2, x5), (x1, x2, x6), (x1, x3,
x4), (x1, x3, x5), (x1, x3, x6), (x1, x4, x5), (x1, x4, x6), (x1, x5,
x6), (x2, x3, x4), (x2, x3, x5), (x2, x3, x6), (x2, x4, x5), (x2, x4,
x6), (x2, x5, x6), (x3, x4, x5), (x3, x4, x6), (x3, x5, x6) and (x4,
x5, x6).
Since the size of our article is limited, we shall deter-
mine only one general solution of the system (40). It is
the solution (x1, x2, x3).
1° Let x4, x5 and x6 be arbitrary real numbers, then the
general solution of the system (40) is:
x1 = x2 = (x4 + x5 + x6)/2,
x3 = (x4 + 2x5 + 3x6)/2 (41)
After the substitution of (41) into (39), the balanced
reaction obtains this general form:
[(x4 + x5 + x6)/2] X2 + [(x4 + x5 + x6)/2] Y2
+ [(x4 + 2x5 + 3x6)/2] O2
 x4 XYO + x5 XYO2 + x6 XYO3 (42)
I. B. RISTESKI: A NEW GENERALIZED ALGEBRA FOR THE BALANCING OF  CHEMICAL REACTIONS
Materiali in tehnologije / Materials and technology 48 (2014) 2, 215–219 217
 x4, x5, x6 
  .
For reaction (39) to be fully balanced, the remaining
19 triads must be determined.
Example 2. For X = H and Y = Cl, we obtain a sub-
particular reaction.
Reaction 3. Now we shall balance this general 
reaction with three arbitrary elements:
x1 XY2 + x2 Z  x3 XY + x4 X2Y3
+ x5 X3Y4 + x6 X3Z + x7 ZY + x8 ZY2 (43)
The above chemical reaction (43) reduces to the fol-
lowing system of linear equations:
x1 = x3 + 2x4 + 3x5 + 3x6,
2x1 = x3 + 3x4 + 4x5 + x7 + 2x8,
x2 = x6 + x7 + x8 (44)
Since the system (44) has three linear equations and
eight unknowns, we can solve it in 8!/[3!(8 – 3)!] = 56
ways. Actually, we must determine all the possible gene-
ral solutions of the system (44). They are the following
triads: (x1, x2, x3), (x1, x2, x4), (x1, x2, x5), (x1, x2, x6), (x1, x2,
x7), (x1, x2, x8), (x1, x3, x4), (x1, x3, x5), (x1, x3, x6), (x1, x3,
x7), (x1, x3, x8), (x1, x4, x5), (x1, x4, x6), (x1, x4, x7), (x1, x4,
x8), (x1, x5, x6), (x1, x5, x7), (x1, x5, x8), (x1, x6, x7), (x1, x6,
x8), (x1, x7, x8), (x2, x3, x4), (x2, x3, x5), (x2, x3, x6), (x2, x3,
x7), (x2, x3, x8), (x2, x4, x5), (x2, x4, x6), (x2, x4, x7), (x2, x4,
x8), (x2, x5, x6), (x2, x5, x7), (x2, x5, x8), (x2, x6, x7), (x2, x6,
x8), (x2, x7, x8), (x3, x4, x5), (x3, x4, x6), (x3, x4, x7), (x3, x4,
x8), (x3, x5, x6), (x3, x5, x7), (x3, x5, x8), (x3, x6, x7), (x3, x6,
x8), (x3, x7, x8), (x4, x5, x6), (x4, x5, x7), (x4, x5, x8), (x4, x6,
x7), (x4, x6, x8), (x4, x7, x8), (x5, x6, x7), (x5, x6, x8), (x5, x7,
x8) and (x6, x7, x8).
As we mentioned previously, the size of the article is
limited, and so we shall determine only one general solu-
tion for the system (44). It is the solution (x1, x2, x3).
1° Let us assume x4, x5, x6, x7 and x8 are arbitrary real
numbers, then the general solution of the system (44) is:
x1 = x4 + x5 – 3x6 + x7 + 2x8,
x2 = x6 + x7 + x8,
x3 = – x4 – 2x5 – 6x6 + x7 + 2x8 (45)
After the substitution of (45) into (43), the balanced
reaction obtains this general form:
(x4 + x5 – 3x6 + x7 + 2x8) XY2 + (x6 + x7 + x8) Z
 (– x4 – 2x5 – 6x6 + x7 + 2x8) XY
+ x4 X2Y3 + x5 X3Y4 + x6 X3Z + x7 ZY + x8 ZY2 (46)
Since the generators x1, x3 > 0, these inequalities
should immediately follow:
x4 + x5 – 3x6 + x7 + 2x8 > 0,
– x4 – 2x5 – 6x6 + x7 + 2x8 > 0 (47)
From (47) we obtain:
x7 + 2x8 > x4 + 2x5 + 6x6 (48)
The above inequality (48) is a necessary and suffi-
cient condition to hold the general reaction (46), i.e., the
reaction (46) holds if and only if (48) is satisfied.
For reaction (43) to be fully balanced the remaining
55 triads must be determined.
Example 3. For X = Mn 
 Fe, Y = O, and Z = C we
obtain a sub-particular reaction.
3 DISCUSSION
The  chemical reactions are a special kind of reac-
tions that have non-unique coefficients. In chemistry,
until now, they were balanced like a reaction with an
infinite number of coefficients, which is incorrect. Every
 reaction has n!/[k!(n – k)!] general reactions, where n
is the number of reaction molecules and k is the number
of reaction elements. Each of these general reactions has
an infinite number of sets of coefficients. In other words,
every  reaction reduces to n!/[k!(n – k)!] general reac-
tions with an infinite number of particular sub-reactions
for each of them.
In this article we determined all the general reactions
of the reaction (1), which are given by the expressions
(4), (6), (10), (13), (17), (20), (24), (28), (32) and (36).
Also, for all of them we determined the necessary and
sufficient conditions for which they hold. In three exam-
ples we showed that this approach to the balancing of 
reactions works successfully. We would also like to men-
tion that the examples 1, 2 and 3 are derived sub-parti-
cular reactions, which are not fully balanced. The readers
can derive the other general solutions very easily,
because they are similar to those of reaction (1), which
we derived using the technique of generalized algebra.
4 CONCLUSION
In this article three  general chemical reactions are
balanced. All the chemical reactions looked similar to
elementary molecular reactions, but they were very hard
to balance. Using this method of generalized algebra, the
author proved again that balancing chemical reactions
does not have anything to do with chemistry – it is a
purely mathematical issue.
The strengths of the method of generalized algebra
are:
1. This method provides an alternative approach for
balancing  chemical reactions. This method showed
that matrix methods can be substituted by the method
of generalized algebra.
2. Since this method of generalized algebra is well for-
malized, it belongs to the class of consistent methods
for balancing chemical reactions.
3. This method of generalized algebra showed that for
any  chemical reaction a topology of its solutions
can be introduced.
4. In fact, the offered method of generalized algebra
simplifies the mathematical operations provided by
the previous well-known matrix methods and is very
suitable for daily practice. The method of generalized
algebra has this advantage, because it fits for all 
I. B. RISTESKI: A NEW GENERALIZED ALGEBRA FOR THE BALANCING OF  CHEMICAL REACTIONS
218 Materiali in tehnologije / Materials and technology 48 (2014) 2, 215–219
chemical reactions, which previously were only
balanced by the methods of generalized matrix in-
verses.
5. For a determination of general reactions any method
for the solution of a system of linear equations can be
used.
6. Using this method the general forms of the balanced
chemical reactions are determined much faster than
by other matrix methods.
7. From the general balanced reactions the other parti-
cular and sub-particular reactions can be determined.
8. Using the method of generalized algebra the dimen-
sion of the solution space can be determined.
9. Using this method the basis of the solution space can
be determined.
10. Necessary and sufficient conditions for which some
reaction holds can be determined by this method as
well. These conditions determine the possibility of
the reaction interval.
11. This method gives an opportunity to be extended with
other numerical calculations necessary for  reac-
tions.
12. The method of generalized algebra represents a good
basis for building a software package.
The weak sides of the method are:
1. Using this method the minimal reaction coefficients
cannot be determined.
2. This method cannot recognize when a chemical reac-
tion reduces to one generator reaction.
3. It cannot predict quantitative relations among the
reaction coefficients.
4. This method cannot arrange the molecules’ disposi-
tion.
5. The method of generalized algebra cannot predict
reaction stability.
This method opens doors in chemistry and mathematics
for new research on  chemical reactions, which unfor-
tunately today cannot be balanced using a computer,
because there is not such a method. The method of
generalized algebra creates a large challenge for
researchers to extend and adapt its usage for computer
application. This is not an easy and simple job, but it
deserves to be realized as soon as possible.
5 REFERENCES
1 I. B. Risteski, The new algebraic criterions to even out the chemical
reactions, In: 22nd October Meeting of Miners & Metallurgists:
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2 I. B. Risteski, A new approach to balancing chemical equations,
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4 I. B. Risteski, A new pseudoinverse matrix method for balancing
chemical equations and thier stability, J. Korean Chem. Soc., 52
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I. B. RISTESKI: A NEW GENERALIZED ALGEBRA FOR THE BALANCING OF  CHEMICAL REACTIONS
Materiali in tehnologije / Materials and technology 48 (2014) 2, 215–219 219

A. REDJECHTA et al.: ELECTRODEPOSITION AND CHARACTERIZATION OF Cu-Zn ALLOY FILMS ...
ELECTRODEPOSITION AND CHARACTERIZATION OF
Cu-Zn ALLOY FILMS OBTAINED FROM A SULFATE
BATH
ELEKTRONANOS IN KARAKTERIZACIJA PLASTI ZLITIN Cu-Zn,
NASTALIH IZ SULFATNE KOPELI
Abdelouahab Redjechta1, Kzmel Loucif1, Loubna Mentar2, Mohamed Redha Khelladi2,
Abdelkrim Beniaiche3
1Laboratoire des Matériaux Non Métalliques, Institut d’Optique et Mécanique de Précision, Université Ferhat Abbas-Sétif 1, 19000 Sétif, Algeria
2Laboratoire de Chimie, Ingénierie Moléculaire et Nanostructures, Ferhat Abbas-Sétif 1, 19000 Sétif, Algeria
3Laboratoire des Systèmes Photoniques et Optiques Non Linéaires, Institut d’Optique et Mécanique de Précision, Université Ferhat
Abbas-Sétif 1, 19000 Sétif, Algeria
redjechtaabdelouahab@yahoo.fr
Prejem rokopisa – received: 2013-02-26; sprejem za objavo – accepted for publication: 2013-04-16
In this work, we report the influence of the deposition potential on the electrodeposition process, current efficiency, surface
morphology and microstructure of Cu-Zn alloys deposited on a Ru substrate from a sulfate solution with an addition of EDTA.
The study was carried out by means of cyclic voltammetry (CV), chronoamperometry, atomic force microscopy (AFM) and
X-ray diffraction (XRD) techniques analyzing the electrochemical behavior, surface morphology and structural characterization,
respectively. The experimental results show that the electrochemical behavior of Cu-Zn electrodeposits varied with the
deposition potential. The AFM measurement showed that the Cu-Zn thin films obtained at all the potentials are homogenous in
appearance being of a small crystallite size, and a variation in the film roughness with deposition potentials is established. An
analysis of X-ray diffraction patterns indicates that the electrodeposited Cu-Zn alloys exhibit - and -phases.
Keywords: copper-zinc, electrodeposition, cyclic voltammetry, morphology, X-ray diffraction
V tem delu poro~amo o vplivu potenciala nanosa pri postopku elektronana{anja na u~inkovitost toka, morfologijo povr{ine in
mikrostrukturo zlitine Cu-Zn, nanesene na podlago iz Ru iz sulfatne raztopine z dodatkom EDTA. Za analizo elektrokemijskega
vedenja, morfologije povr{ine in zna~ilnosti strukture so bile uporabljene cikli~na voltametrija (CV), kronoamperometrija,
mikroskopija na atomsko silo (AFM) in rentgenska difrakcija (XRD). Rezultati raziskav ka`ejo, da se elektrokemijsko vedenje
elektronanosov Cu-Zn spreminja s spreminjanjem potenciala pri nana{anju. AFM-meritve so pokazale, da so tanke plasti Cu-Zn,
dobljene pri vseh potencialih, na videz homogene z majhnimi kristalnimi zrni, spreminjanje potenciala pa vpliva na hrapavost
povr{ine. XRD-analize poka`ejo, da zlitina Cu-Zn po elektronanosu vsebuje - in -faze.
Klju~ne besede: baker-cink, elektronanos, cikli~na voltametrija, morfologija, rentgenska difrakcija
1 INTRODUCTION
The production of the coatings made of zinc and its
alloys has recently been of interest since alloy coatings
provide a better corrosion protection than pure-zinc
coatings. In addition, alloy coatings are very interesting
due to their high strength, good plasticity and excellent
mechanical properties. There are several methods for
obtaining these alloys: physical vapor deposition (PVD),
chemical vapor deposition (CVD), sputtering and mole-
cular beam epitaxy (MBE) techniques are just a few of
them. These methods have several advantages and are
used for specific applications. However, due to certain
limitations, such as high capital and high-energy costs,
an alternative method is required. Recently, the electro-
chemical deposition (electrodeposition) has been used as
an alternative technique for producing these structures on
different surfaces. Electrochemical processes offer many
advantages, including a room-temperature operation,
low-energy requirements, fast deposition rates, a fairly
uniform deposition over complex three-dimensional
objects, low costs and a simple scale-up with an easily
maintainable equipment.1 The control of the solution
composition and deposition parameters determines the
properties of a deposit. In electrodeposition, the mecha-
nism growth, the morphology and the micro-structural
properties of a film depend on electrodeposition condi-
tions such as the electrolyte composition, the electrolyte
pH and the deposition potential.2
Numerous studies of the electrodeposition of Cu-Zn
alloys from aqueous baths have been carried out.3,4 It has
been reported that different electrochemical deposition para-
meters such as deposition potential or current density,
temperature, pH, substrate-surface preparation and bath
composition affect considerably the properties of depo-
sits.5–9 It is known that high-quality micrometer-thick
films (smooth and bright deposits) can be prepared at a
reasonably high deposition rate using baths with a high
metallic-ion concentration and small amounts of additi-
ves.10,11 In the same way, due to a large difference bet-
ween the standard electrode potentials of Cu and Zn (≅1.1
V), these ions should be complex in electrodeposition
solutions to facilitate their codeposition. Therefore, in
Materiali in tehnologije / Materials and technology 48 (2014) 2, 221–226 221
UDK 621.793 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)221(2014)
this work, an organic additive of C10H14Na2O8, 2H2O,
called EDTA, was added to the sulfate bath.
The objective of the present work was to study the
electrodeposition process and properties of the Cu-Zn
alloys from a sulfate electrolyte with EDTA. The mor-
phology and structure of the deposits were examined.
2 EXPERIMENTAL WORK
A deposition of Cu-Zn alloys was carried out in a
bath of 0.14 M CuSO4 for Cu, 0.06 M ZnSO4 for Zn
(Aldrich) with the Na2SO4 support electrolyte, 0.5 M
H3BO3 (in order to control the pH of the solution and
improve the quality of the deposit) and 0.35 M
C10H14Na2O8, 2H2O (EDTA) at pH  4.2 (Table 1). All
the measurements were made at room temperature.
Ethylenediaminetetra acetic acid, widely abbreviated as
EDTA, was chosen as the complexing agent for a
deposition of theses alloys. Plating baths were prepared
from the chemicals of analytical grade and bidistilled
water, and the pH was adjusted with dilute sulfuric acid
when needed. Before the electrodeposition, each solution
was stirred with a nitrogen gas flow. The conventional
electrochemical measurements were taken using a glass
cell consisting of a three-electrode assembly that was
connected to a VoltaLab 40 (PGZ301 and Volta Master
4) controlled by a personal computer. A platinum sheet
was used as the counter electrode (anode) and the
cathode (the Ru substrate) potentials were referred with
respect to the saturated calomel electrode (SCE). The
working electrode was an thick Ru layer approximately
200 nm deposited by sputtering onto each silicon wafer
at a low temperature (150 °C) to get a better adherence.
Before the electrodeposition, the substrates were first
cleaned ultrasonically in acetone and ethanol and then
also with distilled and deionized water. The Cu-Zn
thin-film deposition onto the Ru surface (an area of
0.5 cm2) was studied by means of cyclic voltammetry
(CV) and chronoamperometry (CA) techniques. The CV
for all the solutions was initially carried out between
–1.2 V and 0.2 V (SCE) at a scan rate of 20 mV s–1. The
surface morphologies of the deposits were examined
with atomic force microscopy (AFM). The roughness
(the root-mean-square height deviation) of the samples
was obtained directly from the AFM software (PicoScan
5.3 from Molecular Imaging). The crystalline structures
of the deposits were identified with X-ray diffraction
using a Philips diffractometer with a 2 range 10–100°
and Cu K radiation ( = 0.15406 nm).
3 RESULTS AND DISCUSSIONS
3.1 Electrochemical study
Cyclic voltammetry was performed to understand the
electrochemical behavior of the Cu(II) and Zn(II) species
on the Ru electrode. Figure 1 shows the cyclic voltam-
mograms of the Ru electrode recorded in different ion
solutions of Cu, Zn and Cu-Zn. In effect, Figure 1a
shows a cyclic voltammogram of a solution containing
0.14 M CuSO4 with a cathodic scan limit of –0.8 V vs.
SCE. Two sharp peaks are observed at –0.141 V and
0.072 V, corresponding to a reduction and a dissolution
of Cu, respectively. In the Cu electrodeposition, the
charge transfer step is fast and the growth rate is con-
trolled with the rate of the Cu-ion mass transfer to the
A. REDJECHTA et al.: ELECTRODEPOSITION AND CHARACTERIZATION OF Cu-Zn ALLOY FILMS ...
222 Materiali in tehnologije / Materials and technology 48 (2014) 2, 221–226
Figure 1: Cyclic voltammograms obtained for: a) 0.14 M CuSO4, b)
0.06 M ZnSO4 + 0.35 M EDTA and c) 0.14 M CuSO4 + 0.06 M
ZnSO4 + 0.35 M EDTA with the cathodic scan limit of –1.20 V vs.
SCE, at the scan rate of 20 mV s–1; supporting electrolyte is 1 M
Na2SO4 + 0.5 M H3BO3 (pH 4.2)
Slika 1: Cikli~ni voltamogrami, dobljeni z: a) 0,14 M CuSO4, b) 0,06
M ZnSO4 + 0,35 M EDTA in c) 0,14 M CuSO4 + 0,06 M ZnSO4 +
0,35 M EDTA s katodno omejitvijo –1,20 V proti SCE, pri hitrosti
skeniranja 20 mV s–1; osnovni elektrolit je 1 M Na2SO4 + 0,5 M
H3BO3 (pH 4,2)
Table 1: Bath composition and conditions for the Cu-Zn electrodepo-
sition
Tabela 1: Sestava kopeli in razmere pri elektronana{anju Cu-Zn
Bath ZnSO4,7H2O/M
CuSO4,
5H2O/M
Na2SO4
/M
H3BO3
/M
EDTA
/M
Zn 0.06 1 0.5 0.35
Cu 0.14 1 0.5
Cu-Zn 0.06 0.14 1 0.5 0.35
growing centers. The consistency of the cyclic-voltam-
metry behavior upon the potential cycling indicates that
the anodic stripping process completely removes Cu
from the electrode surface. The data in this figure
indicates the absence of an underpotential deposition
peak, with the Cu reduction occurring at the significant
overpotential to the Nernstian value. This is due to a
weak deposit/substrate interaction, as the early stages of
an electrodeposition of Cu on Ru surfaces correspond to
the Volmer-Weber growth mechanism.12
Figure 1b shows a cyclic voltammogram obtained in
0.06 M ZnSO4. During a direct scan, it is possible to note
that the increase in the current begins at –0.7 V; this is
due to the electrodeposition of Zn and hydrogen evolu-
tion. In the reverse potential scan, the absence of the
peak corresponding to the dissolution of the previously
deposited Zn is observed. For the Cu and Zn solution
(Figure 1c), the voltammogram obtained shows the
presence of cathodic and anodic peaks related to the
deposition and dissolution of the metals. In the cathodic
scan, it can be observed that the increases in the current
were detected at –0.152 V and –0.75 V, being charac-
teristic of the potential deposition processes of Cu and
Zn onto Ru surfaces, respectively. After this limit, the
hydrogen evolution is predominant. During the inverse of
the potential scan, it is possible to observe, in all the
curves, the presence of crossovers which are typical of
the formation of a new phase involving a nucleation pro-
cess.13
To elucidate the role of an applied potential, the
current efficiency (CE) during the codeposition process
was determined. The deposition CE was calculated from
the ratio of the cathodic electric charge, which passed
during the electrodeposition of the Cu-Zn alloy, to that of
the anodic one required for the total alloy dissolution.
The Cu-Zn alloy thin films were obtained in the poten-
tiostatic mode at different deposition potentials. The
dependence of CE on different deposition potentials is
shown in Figure 2. The efficiencies of the deposits
decrease considerably with the potential and then reach
their minima at more negative potentials. This decrease
in CE is due to the process of hydrogen evolution. It is
clear from these results that there is an appreciable
decrease in the value of CE as the deposition potentials
become more negative than  –1.0 V. This is due to the
hydrogen evolution reaction (HER) that becomes more
significant than the Zn and Cu electrodeposition. This, in
turn, increases the pH level at the cathode, causing the
metal hydroxide to be included in the deposit.11 These
observations indicate that the control of the deposition
potential is very important for realizing a high CE in the
Cu-Zn deposition process. Similar results for Cu-Zn
deposits are observed by de Almeida et al.14
The current was recorded as a function of time to
study the deposition mechanisms of Cu-Zn alloys during
their growth. The electrochemical deposition was per-
formed using the standard chronoamperometry technique
to study the nucleation and growth mechanism of Cu-Zn
on ruthenium. Deposition potentials were chosen
according to the reduction peaks appearing on the cyclic
voltammograms. Deposition is conducted at the constant
potentials in the potentiostatic mode, during which
current transients are recorded. Figure 3 shows the
current transients obtained at four different potentials:
–1.0 V, –1.1 V, –1.2 V and –1.3 V vs. SCE. In this figure,
at the beginning of the applied potentials, a high cathodic
current is seen for a short time. After that, the current
rapidly decreases due to a depletion of the metal-ion
concentrations close to the electrode surface, subse-
quently reaching a stable value. The current-time tran-
sients have a normal dependence on the overpotentials,
whereas the current density increases with an increase in
the overpotential. This is specific to the nucleation and
growth process and for longer times, merging into a
common curve caused by the diffusion-controlled pro-
A. REDJECHTA et al.: ELECTRODEPOSITION AND CHARACTERIZATION OF Cu-Zn ALLOY FILMS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 221–226 223
Figure 3: Evolution of current densities versus deposition time during
the deposition of Cu-Zn on the Ru surface at different deposition
potentials
Slika 3: Razvoj gostote toka v primerjavi s ~asom nanosa Cu-Zn na
povr{ino Ru pri razli~nih potencialih nana{anja
Figure 2: Effect of the deposition potential on the current efficiency
of the Cu-Zn electrodeposition process
Slika 2: Vpliv potenciala nanosa na u~inkovitost elektronana{anja
Cu-Zn
cess and described with the Cottrell equation.15 Accord-
ing to the i–t curves, each transient has one well-defined,
recognizable current maximum seen as a clear first peak
followed by a sharp fall and subsequent growth. The i–t
transients have a normal dependence on the overpoten-
tials, whereas the current density increases with an
increase in the overpotential. The peak corresponds to
the nucleation of the metallic sites on the surface and it
is followed by a reduction in the current exhibiting a
three-dimensional (3D) growth. An increase in the peak
current at higher overpotentials means that the number of
sites (nucleation rate) increases due to a higher nucle-
ation rate.15
3.2 Morphological and Structural Analysis
The morphology of the electrodeposited surface was
imaged ex situ after the Cu-Zn electrodeposition using
AFM measurements. Figure 4 shows 2 μm × 2 μm AFM
images of the deposited Cu-Zn alloy films obtained at
different deposition potentials. The figures reveal that,
with all the applied potentials, the images have a granu-
lar surface. However, the dimensions of the visible
features in the images are different. It is known that a
film electrodeposited on a polycrystalline substrate is
also polycrystalline. During electrolysis, Cu-Zn crystal-
lites randomly grow on the polycrystalline substrate and
may form conglomerates. The composition and crystal-
lite sizes strongly depend on the applied potential. If the
current density is very small, there will be insufficient
crystalline growth centers and the deposited layer will be
rough-grained. If the current density is high, the depo-
sited layer will be porous and soft.
The surface topography is traditionally analyzed with
surface-roughness measurements such as the root-mean-
square (RMS) roughness, the average roughness and the
peak-to-valley roughness.16 In brief, the surface rough-
ness Rq (denoted also as RMS) and the mean roughness
Ra were calculated using the standard software (Table 2).
The surface roughness increases with the film thickness
for the films deposited using both potentials, and the
variation in the surface morphology with the applied
potential is consistent with the general theory.17
Table 2: Dependence of the surface roughness and crystallite size of
the electrodeposited Cu-Zn thin films on deposition potentials
Tabela 2: Odvisnost hrapavosti povr{ine in velikosti kristalnih zrn
Cu-Zn tanke plasti po elektronana{anju od potenciala pri nanosu
E/(V vs. SCE) Rq/nm Ra/nm D/nm
–1.10 42.72 33.12 48.20
–1.20 53.04 65.60 46.50
–1.30 108.62 82.08 46.10
The effect of the current density on this surface mor-
phology can be explained since a high current density
results in higher rates of the crystal nucleation (a higher
mobility of atoms), giving rise to finer crystal structures
and, hence, a smoother surface.17 However, a new
theory18,19 has emerged, proposing that the concentration
of metallic ions does change the bath homogeneously,
but rather preferentially increases near the substrate
(cathode). This relative concentration in the discharging
zone is of a little significance at a low current density, at
which the surface of a deposited film is rough. At a high
current density, the convexity of the film increases, being
associated with a relative concentration of ions in the
discharging zone. The surface of a film deposited under
these conditions is very smooth. The convex part of the
film attracts more ions by acting as a focus of discharge,
further increasing the convexity of the film. This may
explain the mechanism of deposition.20
Figure 5 shows the XRD patterns of the Cu-Zn films
deposited on the Ru substrates in the sulfate/EDTA bath
under different deposition potentials of –1.1 V, –1.2 V
and –1.3 V vs. SCE in the 2 scan range of 25–60°. All
the XRD patterns show many peaks corresponding to
two distinct phases,  and , respectively. It is clear that
A. REDJECHTA et al.: ELECTRODEPOSITION AND CHARACTERIZATION OF Cu-Zn ALLOY FILMS ...
224 Materiali in tehnologije / Materials and technology 48 (2014) 2, 221–226
Figure 4: AFM images of the Cu-Zn thin films prepared at different
cathodic potentials: a) –1.1 V, b) –1.2 V and c) –1.3 V vs. SCE
Slika 4: AFM-posnetki tankih plasti Cu-Zn, pripravljenih pri razli~nih
katodnih potencialih: a) –1,1 V, b) –1,2 V in c) –1,3 V v odvisnosti od
SCE
the XRD patterns from the Cu-Zn electrodeposits are
different from those of pure Zn and Cu, indicating that
crystalline alloys are indeed formed in these Cu-Zn
electrodeposits. Furthermore, the  phase diffraction
lines increase in intensity as the deposition potential
becomes more negative; in other words, the  phase
increases as the Cu content in the electrodeposited
Cu-Zn alloy decreases. At a higher deposition potential,
a decrease in the Cu concentration is explained with the
fact that, at these potentials, a reduction in Cu is mass-
transport limited. A further increase in the deposition
overpotential would only increase the amount of Zn
being deposited.21–23 From these results, the  phase was
more dominant than the  phase in the Cu-Zn alloy thin
films.
The average crystallite size of the particles is calcul-
ated from the full width at half maximum (FWHM) of
the respective peaks using the Scherrer relation:24
D =
0 9.
cos

 
(1)
where D is the crystallite size,  is the wavelength of
X-ray radiation ( = 0.15406 nm),  is the FWHM of the
peak and  is the diffraction angle.
Also, Table 2 shows the average crystallite size
obtained from XRD for [001] planes, for the  phase of
the alloys electrodeposited at three different applied
potentials. The average crystallite size decreases with the
increasing Zn concentration in the films and with the
increasing applied potentials. This observation shows
that, at a more negative potential, the deposition rate is
high and, hence, the atoms are incorporated in the film
with little surface migration, thus limiting the grain size.
4 CONCLUSIONS
Smooth, compact and bright binary Cu-Zn alloys
were deposited on a Ru substrate from a sulfate
electrolyte with an EDTA additive. The electrodeposition
behavior of the sulfate electrolyte was studied using
cyclic voltammetry. A possibility of depositing pure
copper and zinc with a trace of copper was revealed
during a cathodic scan of the substrate potential. The
effects of deposition potentials on the microstructures of
Cu-Zn were investigated by means of AFM and XRD
techniques. The AFM images showed Cu-Zn clusters of
an equivalent size randomly distributed in the surface
defects acting as active sites. An X-ray diffraction
measurement reveals that the Cu-Zn alloy exhibits two
phases, the  and  phases. According to these results,
the  phase was more dominant than the  phase in the
Cu-Zn alloy thin films.
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Materiali in tehnologije / Materials and technology 48 (2014) 2, 221–226 225
Figure 5: X-ray diffraction patterns of the electrodeposited Cu-Zn
alloy films obtained at different deposition potentials: a) –1.1 V, b)
–1.2 V and c) –1.3 V vs. SCE
Slika 5: XRD-posnetki tankih plasti Cu-Zn, dobljeni pri razli~nih
potencialih nana{anja: a) –1,1 V, b) –1,2 V in c) –1,3 V v primerjavi s
SCE
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226 Materiali in tehnologije / Materials and technology 48 (2014) 2, 221–226
M. T. OZKAN et al.: EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL NETWORK MODEL ...
EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL
NETWORK MODEL FOR TURNING THE 50CrV4
(SAE 6150) ALLOY USING COATED CARBIDE/CERMET
CUTTING TOOLS
EKSPERIMENTALNA ZASNOVA IN MODEL UMETNE
NEVRONSKE MRE@E ZA STRU@ENJE JEKLA 50CrV4 (SAE 6150)
Z UPORABO ORODIJ S KARBIDNO ALI KERMETNO PREVLEKO
Murat Tolga Ozkan1, Hasan Basri Ulas2, Musa Bilgin3
1Gazi University, Faculty of Technology, Department of Industrial Design Engineering, 06500 Ankara, Turkey
2Gazi University, Faculty of Technical Education, Mechanical Education Department, 06500 Ankara, Turkey
3Erzincan University, Vocational Technical School, Mechanical Program, 24000 Erzincan, Turkey
mtozkan06@yahoo.com
Prejem rokopisa – received: 2013-03-08; sprejem za objavo – accepted for publication: 2013-06-18
In this experimental study, the 50CrV4 (SAE 6150) steel was subjected to the machining tests with coated carbide and cermet
cutting tools in a turning operation. The tests were carried out at various cutting speeds, feed rates and cutting depths. In the
light of these parameters, cutting forces and surface-roughness values were determined. Three components (Fa, Fr and Fc) of the
cutting forces were measured during the tests using a dynamometer, while the machined surface-roughness values were
determined using a surface roughness measuring unit. A multiple regression analysis and experimental design were performed
statistically. The measured surface-roughness values were used for the modeling with an artificial neural network system
(ANNS). The relations between the cutting forces and the surface-roughness values were also defined.
Keywords: turning operations, coated carbide/cermet cutting tools, cutting force, surface roughness, artificial neural network
V tej {tudiji je bilo jeklo 50CrV4 (SAE 6150) strojno obdelano s stru`enjem z orodji za rezanje s karbidno ali kermetno pre-
vleko. Preizkusi so bili opravljeni pri razli~nih hitrostih rezanja, podajanja in globine rezanja. Iz teh parametrov so bile
ugotovljene sile pri rezanju in vrednosti hrapavosti povr{ine. Tri komponente sil rezanja (Fa, Fr in Fc) so bile merjene med
poskusom z dinamometrom, medtem ko so bile vrednosti hrapavosti povr{ine izmerjene z merilnikom za hrapavost povr{ine.
Izvr{ena je bila ve~kratna regresijska analiza in statisti~na obdelava izvedbe poskusov. Izmerjene vrednosti hrapavosti so bile
uporabljene za modeliranje s sistemom umetne nevronske mre`e (ANNS). Dolo~ena je bila tudi odvisnost med silami rezanja in
hrapavostjo povr{ine.
Klju~ne besede: postopek stru`enja, rezilna orodja s karbidno ali kermetno prevleko, sila rezanja, hrapavost povr{ine, umetna
nevronska mre`a
1 INTRODUCTION
Machining experiments are usually costly and time
consuming. In order to eliminate the cost and to reduce
the time, some advanced techniques were developed.
These are generally called modeling. There are several
modeling methods. Empirical modeling, analytical
modeling, mechanistic modeling, finite-element model-
ing and artificial neural network modeling are some of
these modeling methods. With these techniques, predic-
tions are made using the experimentally obtained results.
Boubekri et al.1 investigated different cutting condi-
tions on three grades of steel, namely, 1018 (low-carbon,
cold-rolled steel), 304 (austenitic stainless steel) and
4140 (low-alloy steel) using uncoated carbide inserts as
the tool material. Mathematical models were developed
for predicting the forces acting on the tool for different
cutting conditions and materials.
Kumar et al.2 studied the performance of an alumi-
na-based ceramic cutting tool as an attractive alternative
for carbide tools in the machining of steels in its
hardened condition. Two types of ceramic cutting tools,
namely, the Ti[C, N] mixed alumina-ceramic cutting tool
and zirconia-toughened alumina-ceramic cutting tool
were used for the investigation. The performance of
these ceramic cutting tools related to the surface finish
was also discussed.
Tekiner and Yeþilyurt3 carried out machining tests to
determine the best cutting conditions and cutting
parameters in the turning of AISI 304 stainless steels by
taking into consideration the process sound. The ideal
cutting parameters and cutting-process sounds were
determined.
The tests involving the AISI 4340 steel were per-
formed using two hardness values, 42 HRC and 48 HRC;
for the former, a coated carbide insert was used as the
cutting tool, whereas for the latter a polycrystalline cubic
boron nitride insert was employed. The machining tests
on the AISI D2 steel hardened to 58 HRC were conduc-
ted using a mixed-alumina cutting tool. The cutting
forces, surface roughness, tool life and wear mechanisms
were assessed. The results indicated that when turning
Materiali in tehnologije / Materials and technology 48 (2014) 2, 227–236 227
UDK 621.941:004.032.26 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)227(2014)
the AISI 4340 steel, using low feed rates and depths of
cut, the forces were higher when machining the softer
steel and that the surface roughness of the machined part
was improved as the cutting speed was elevated and it
deteriorated with the increasing feed rate.4
The wear mechanisms of the cutting tools made of
tungsten-carbide (WC), PCBN and PCD were investiga-
ted using the tool life and the temperature results avail-
able in the literature. For the tool/work combinations of
WC/steel and PCBN/hardened-steel, under practical con-
ditions, the tool wear was found to be greatly influenced
by the temperature.5
The AISI 4340 low-alloy steel was assumed for the
workpiece. Cubic boron-nitride (CBN) or polycrystalline
(PCBN) inserts are widely used as the cutting-tool mate-
rial in high-speed machining of hardened tool steels due
to high hardness, high abrasive-wear resistance and che-
mical stability at high temperature. Cutting forces and
feed forces were determined with numerical simulations.
The cutting force and feed force increased with the
increasing feed, tool-edge radius, negative rake angle,
and workpiece hardness. The results from the simula-
tions were compared to the experimental results reported
in the literature.6
The effects of the cutting speed, cutting force and
feed rate on the flank wear land and tool life of the TiN
coated carbide inserts in boring operations were studied
when machining the 38MnVS5 alloy. The results showed
that the machinability of microalloyed steel is better than
that of heat-treated alloy steels under identical condi-
tions.7
The tribological influences of PVD-applied TiAlN
coatings on the wear of cemented carbide inserts and the
microstructure influence on the wear behaviors of the
coated tools were investigated under dry and wet
machining. The turning test was conducted with variable
high cutting speeds ranging from 210 m/min to 410
m/min.8
The tested work materials were plain carbon steel JIS
S45C and BN free-machining steels. The JIS S45C was
used as the standard. The tool wear in turning the BN
free-machining steel was smaller than that in turning the
standard steel. In the case of turning BN1 with P30 at
200 m/min, 300 m/min, the wear-progress rate of the
flank wear and crater depth were about half as much as
that in turning the standard steel. The BN free-machining
steel showed a slightly lower cutting temperature and a
smaller cutting force in comparison with the standard
steel at the tested cutting speeds.9
The cutting-force data used in the analyses was
gathered with a tool-breakage detection system detecting
the variations of the cutting forces measured with a
three-dimensional force dynamometer. The workpiece
materials used in the experiments were: cold-work tool
steel, AISI O2 (90 MnCrV8); hot-work tool steel, AISI
H10 (X32CrMo33); and mould steel, AISI improved 420
(X42Cr13). The cutting tools used were HSS tools,
uncoated WC and coated TiAlN and TiC + TiCN + TiN
inserts (ISO P25). No cutting fluid was used during the
turning operations. During the experiments, the cutting
forces, flank wear and surface-roughness values were
measured throughout the tool life and the machining
performances of the tool steels were compared.10
Orthogonal cutting experiments were performed on a
high-speed machining center with the surface speeds of
up to 500 m/min and the uncut chip thicknesses ranging
from 0.1 mm to 0.3 mm. The results indicate that in cer-
tain critical regions of the thermal field, the improved
machinability correlates with significant reductions in
the temperature exceeding the measurement uncertain-
ties. Such micro-scale temperature measurements will
help to design the materials with further improved
machinability.11
Three different carbide cutting tools were used,
namely, TiCN + TiC + TiCN + Al2O3 + TiN-coated car-
bide tools with multilayer coatings of 7.5 μm and 10.5
μm and an uncoated WC/Co tool. The turning experi-
ments were carried out at four different cutting speeds,
which were (125, 150, 175 and 200) m/min. The feed
rate (f) and depth of cut (ap) were kept fixed at 0.25 mm/r
and 1.5 mm/r, respectively, throughout the experiments.
The tool performance was evaluated with respect to the
tool wear, the surface finish produced and the cutting
forces generated during turning.12
Cakir and Isik13 performed a study about the tool-life
testing with single-point turning tools. The cutting-force
data used in the analysis was gathered with a tool-break-
age detection system detecting the variations of the
cutting forces measured with a three-dimensional force
dynamometer. Six pairs of ductile iron specimens,
austempered at (300, 350 and 400) °C for 1 h and 2 h
were tested. The cutting tools used in the tests were
coated carbide inserts, ISO SNMG 120408 (K10),
clamped on the tool holders, CSBNR 2525 M12. No
cutting fluid was used during the turning operations.
During the experiments, the cutting forces, flank wear
and surface-roughness values were measured throughout
the tool life and the machining performances of ADI
having different structures were compared.
Lalwani et al.14 carried out a machining study focus-
ing on the effect of the cutting parameters (cutting speed,
feed rate and depth of cut) on the cutting forces (feed
force, thrust force and cutting force) and the surface
roughness in the finish hard turning of the MDN250 steel
(equivalent to the 18Ni(250) maraging steel) using a
coated ceramic tool. A non-linear quadratic model best
describes the variation in the surface roughness with the
major contribution of the feed rate and the secondary
contributions of the interaction effect between the feed
rate and the depth of cut, the second-order (quadratic)
effect of the feed rate and the interaction effect between
the speed and depth of cut. The suggested models of the
cutting forces and surface roughness are adequately map-
M. T. OZKAN et al.: EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL NETWORK MODEL ...
228 Materiali in tehnologije / Materials and technology 48 (2014) 2, 227–236
ped within the limits of the cutting parameters consi-
dered.
In another study, coated tungsten-carbide ISO P-30
turning-tool inserts were subjected to a deep cryogenic
treatment (–176 °C). The machining studies were con-
ducted on a C45 workpiece using both untreated and
deep cryogenic treated tungsten-carbide cutting-tool
inserts. The cutting force during the machining of the
C45 steel is lower in the case of the deep cryogenic
treated carbide tools when compared with the untreated
carbide tools. The surface finish produced when machin-
ing the C45 steel workpiece is better with the deep cryo-
genic treated carbide tools than with the untreated
carbide tools.15
The objective of the work is to determine the
influence of cutting fluids on the tool wear and surface
roughness during turning AISI 304 with a carbide tool. A
further attempt was made to identify the influence of
coconut oil on reducing the tool wear and surface rough-
ness during the turning process. The performance of
coconut oil was also compared with two other cutting
fluids, namely, an emulsion and a neat cutting oil (immi-
scible with water).16
Ebrahimi and M. M. Moshksar17 machined micro-
alloyed steel (30MnVS6) and quenched-tempered (QT)
steels (AISI 1045 and AISI 5140) under different cutting
conditions. An experimental investigation was conducted
to determine the effects of the cutting speed, feed rate,
hardness, and workpiece material on the flank wear and
tool life of the coated cemented carbide inserts in the
hard-turning process. A statistical analysis was used for
evaluating different factors of the cutting forces. Chip
characteristics and the chip/tool contact length were also
investigated.
In another study, the machining of the uncoated AISI
1030 steel (i.e., the orthogonal cutting), PVD- and
CVD-coated cemented carbide inserts with different feed
rates of (0.25, 0.30, 0.35, 0.40 and 0.45) mm/r, with the
cutting speeds of (100, 200 and 300) m/min and a con-
stant depth of cut (i.e., 2 mm), without using a cooling
liquid was accomplished. The effects of the surface
roughness, the coating method, the coating material, the
cutting speed and the feed rate on the workpiece were
investigated. Afterwards, these experimental studies
were carried out on artificial neural networks (ANNs).
The training and test data of the ANNs were prepared
using experimental patterns for the surface roughness.
Therefore, the surface-roughness value was determined
with an ANN with an acceptable accuracy.18
Gaitonde et al.19 made a study analyzing the effects
of the depth of cut and the machining time on the machi-
nability aspects such as machining force, power, specific
cutting force, surface roughness and tool wear using
second-order mathematical models during turning high-
chromium AISI D2 cold-work tool steel with the CC650,
CC650WG and GC6050WH ceramic inserts.
The effects of the machining parameters on the cutt-
ing force, specific cutting pressure, cutting temperature,
tool wear and surface-finish criteria were investigated
during the experimentation. The present approach and
the results will help manufacturing engineers to under-
stand the machinability of Inconel 718 during high-speed
turning.20
Two AISI 4140 steels with different machinability
ratings and three types of tools were compared. The con-
trol-volume approach was used to estimate the energy
partition from the thermal images and the energy out-
flows were compared for the measurement of the cutting
power. This provides a new physical tool for examining
machinability, tool wear and subsurface damage.21
The work is an experimental study of hard turning
the AISI 52100 bearing steel with a CBN tool. The
relationships between the cutting parameters (cutting
speed, feed rate and depth of cut) and machining output
variables (surface roughness, cutting forces) are analyzed
and modeled with the response-surface methodology
(RSM). Finally, the depth of cut exhibits the maximum
influence on the cutting forces as compared to the feed
rate and cutting speed.22
There are many studies in the literature about the arti-
ficial neural network modeling of the surface roughness.
The main purpose of this study is to carry out the turning
tests on the 50CrV4 (SAE 6150) steel using coated car-
bide/cermet cutting tools and to carry out the modeling
with an artificial neural network. There are many studies
in the literature about the SAE 6150 material. These
studies are especially related to physical and mechanical
properties. However, there is no sufficient study on the
machinability of the SAE 6150 material.
2 MATERIALS AND METHOD
The turning tests were carried out on the commer-
cially available 50CrV4 (equivalent to SAE 6150) steel
workpiece specimens. These steel specimens were pro-
duces by CEMTAS, TR and their chemical composition
is shown in Table 1.
The 50CrV4 (SAE 6150) workpiece specimens were
prepared in accordance with the requirements of the ISO
M. T. OZKAN et al.: EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL NETWORK MODEL ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 227–236 229
Table 1: Workpiece material 50CrV4 (SAE 6150) and its chemical composition (w/%)
Tabela 1: Material obdelovancev 50CrV4 (SAE 6150) in njegova kemijska sestava (w/%)
Material Hardness (HB)
Material content
C Si Mn P S Cr V
50CrV4 - SAE(AISI)6150 311 0.50 0.31 0.78 0.009 0.008 1.06 0.15
3685 standard. The size of the specimens was
determined in accordance with the diameter/length ratio
that should not exceed 1/10. Prior to the machining tests
through a surface turning operation, both faces of the
specimens were machined and center drilled on a
universal lather. The dimensions of the workpiece
specimens are shown in Figure 1.
Tool holder was also chosen according to the require-
ments specified in ISO 3685 for the machinability tests.
The specifications of the tool holder, produced by
SANDVIK, were PCLNR 2525 12. The tool-holder
features are shown in Figure 2.
The cutting inserts with six different features were
used for the machining tests. These cutting inserts were
coated carbide and cermet cutting tools. In addition, the
cutting-insert tip radii were (0.4, 0.8 and 1.2) mm. The
cutting-insert grades and technical specifications are
shown in Table 2.
The machining tests were carried out at the cutting
speeds of (150, 200, 250, 250 and 300) m/min, the feed
rates of (0.12, 0.16, 0.20) mm/r and the depths of cut of
(0.5, 1, 1.5) mm without a coolant, using the cutting
tools whose specifications are given in Table 2. The
cutting speed, depth of cut and feed rate were chosen
according to the manufacturer’s recommendations, the
M. T. OZKAN et al.: EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL NETWORK MODEL ...
230 Materiali in tehnologije / Materials and technology 48 (2014) 2, 227–236
Figure 2: Tool-holder specifications
Slika 2: Specifikacija nosilca orodja
Figure 1: Experimental-workpiece-specimen geometry
Slika 1: Geometrija poskusnega vzorca
Table 2: Cutting-tool technical specifications
Tabela 2: Tehni~ne zna~ilnosti poskusnega orodja za rezanje
Cutting-tool
number
Cutting-tool
material Cutting tool
Coating
method
Coating
material Code Tip radius ISO grade
1 Carbide CVD Al2O3+TiCN GC4215 0.4 CNMG 12 04 04-PF-P15
2 Carbide CVD Al2O3+TiCN GC4215 0.8 CNMG 12 08 04-PF-P15
3 Carbide CVD Al2O3+TiCN GC4215 0.12 CNMG 12 12 04-PF-P15
4 Cermet PVD TiN+TiCN GC1525 0.4 CNMG 12 04 04-PF-P15
5 Cermet PVD TiN+TiCN GC1525 0.8 CNMG 12 08 04-PF-P15
6 Cermet PVD TiN+TiCN GC1525 0.12 CNMG 12 12 04-PF-P15
Table 3: Cutting parameters used in the experimental study
Tabela 3: Parametri rezanja, uporabljeni pri eksperimentih
Cutting tool Tip radius Rå/mm Depth of cut d/mm Feed rate f/(mm/r) Cutting speed v/(mm/min)
Coated carbide 0.4–0.8–1.2 0.5–1–1.5 0.12–0.16–0.20 150–200–250–250–300
Coated cermet 0.4–0.8–1.2 0.5–1–1.5 0.12–0.16–0.20 150–200–250–250–300
related literature and the ISO 3685 standard. Table 3
shows the types of cutting tool, tip radius, depth of cut,
feed rate and cutting speed.
An industrial-type Johnford TC-35 CNC lathe was
used as the machine tool. To measure the axial (Fx),
radial (Fy) and main (Fz) cutting-force components, a
KISTLER 9257B dynamometer was used. This dynamo-
meter was connected with a KISTLER TYPE 5019
multichannel charge amplifier and the cutting-force
signals were sent to a computer with an RS-232C cable
connection. The graphs were obtained with the Dyno-
Ware Type 2825A1-2 software. Fx (Fa), Fy (Fr) and Fz
(Fc) force values were determined during the machining
processes. These values were calculated as the average
scalar values in newtons (N) by the DynoWare software.
A schematic presentation of the experimental setup is
shown in Figure 3, while the experimental setup is
shown in Figure 4.
For the measurement of the surface-roughness values,
a MAHR-Perthometer M1 model device was used
(Figure 5). The cut-off length and the sampling length
were chosen to be 0.8 mm and 5.6 mm, respectively.
Three measurements were made parallel to the longitu-
dinal axis of the workpiece at the intervals of 120°. The
average surface roughness, Ra, was taken into account in
all the measurements.
3 ARTIFICIAL NEURAL NETWORK MODEL
The concept of an artificial neural network has
emerged with the idea that it simulates the operating
principles of a human brain. The first studies were made
with mathematical modeling of biological neurons that
make up the brain cells. An artificial neural network
consists of a large number of interconnected processing
elements. Artificial neural network processing elements
are called simple nerves.
An artificial neural network contains a large number
of interconnected nodes. A nerve is the basic unit of an
artificial neural network. An artificial nerve is, thus,
simpler than a biological nerve. Figure 6 shows an
artificial neural network algorithm. All the artificial
neural networks are derived from this basic structure.
The differences in the structure of artificial neural
networks result in different classification types.
The learning procedure can be affected by establish-
ing correct correlations between the input and the output
in artificial neural networks. This process falls below a
certain value of the error between the foreseen output
and the desired output. Artificial neural networks learn
like a human. The more samples are used for learning,
the more accurate is the obtained result. When a certain
input is entered, the network can make changes to the
data in order to give similarly accurate answers. The
Levenberg-Marquardt (Levenberg, 1944; Marquardt,
1963) method uses a search direction that is a solution of
the linear set of equations:
( ( ) ( ) ) ( ) ( )J x J x x I d J x F xk
T
k k k k
T
k+ = − (1)
or, optionally, of the equations:
( ( ) ( ) ( ( ) ( )))
( ) ( )
J x J x l J x J x d
J x F x
k
T
k k k
T
k k
k
T
k
diag+ =
= −
(2)
M. T. OZKAN et al.: EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL NETWORK MODEL ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 227–236 231
Figure 6: Artificial neural network
Slika 6: Umetna nevronska mre`a
Figure 5: Surface-roughness measurement device
Slika 5: Naprava za merjenje hrapavosti povr{ine
Figure 4: Experimental setup
Slika 4: Eksperimentalni sestav
Figure 3: Schematic presentation of the experimental assembly
Slika 3: Shematski prikaz eksperimentalnega sestava
where scalar k controls both the magnitude and direc-
tion of dk. It is necessary to set the Scale Problem option
to šnone’ to choose Equation 1 and to šJacobian’ to
choose Equation 2.
Multilayer networks often use the log-sigmoid trans-
fer function (logsig). The function generates the outputs
between 0 and 1 as a neuron’s net input goes from the
negative to the positive infinity. Alternatively, multilayer
networks can use the tan-sigmoid transfer function
(tansig). Sigmoid output neurons are often used for
pattern-recognition problems, while linear output neu-
rons are used for function-fitting problems. The linear
transfer function is shown below (Figure 7).
4 RESULTS AND DISCUSSIONS
4.1 Experimental study results
In this study, the influence of the machining para-
meters (cutting speed, feed rate and depth of cut) on the
cutting forces (Fa, Fr and Fc) and surface-roughness
values were investigated when machining the 50CrV4
(SAE 6150) steel. The surface-roughness values were
analyzed on the basis of the process variables such as
cutting speed, feed rate, depth of cut, and cutting-tool
type. In addition, the influences of these variables on the
cutting force were also examined. The relations between
the cutting forces and the surface roughness were
defined. A total of 270 experiments were performed to
examine the effects of these variables. The minimum sur-
face-roughness value was obtained with the coated cer-
met cutting tool, while the maximum surface-roughness
value was obtained with the coated carbide cutting tool.
The cutting-force components were measured (Fa, Fr,
Fc). Maximum Fa, Fr and Fc were obtained with the
coated carbide cutting tool. However, minimum Fa and
Fr were obtained with the coated cermet cutting tool.
Maximum Ra was obtained with the coated cermet cutt-
ing tool, while minimum Ra was obtained with the
coated carbide cutting tool. A direct relation can be seen
between the cutting forces and surface-roughness values.
The lower are Fa and Fr, the lower are the surface-rough-
ness values (Table 4).
M. T. OZKAN et al.: EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL NETWORK MODEL ...
232 Materiali in tehnologije / Materials and technology 48 (2014) 2, 227–236
Figure 7: Used ANN functions
Slika 7: Uporabljene funkcije ANN
Figure 8: Comparison of the cutting forces: machining with coated
carbide/cermet cutting tools
Slika 8: Primerjava sil pri rezanju: rezanje z orodjem s karbidno ali
kermetno prevleko
Table 4: Maximum/minimum values of the cutting force and surface roughness
Tabela 4: Maksimalne/minimalne vrednosti sil pri rezanju in hrapavost povr{ine
Cutting tool (coated
carbide/cermet)
Cutting
tool (z)
Depth of
cut d/mm
Cutting speed
v/(mm/min)
Feed rate
f/(mm/r) Fx (Fa)/N Fy (Fr)/N Fz (Fc)/N
Surface
roughness
Ra/μm
Min/max
Cermet 6 0.5 250 0.12 74.37 118.83 218.68 0.345 Fa min
Carbide 2 1.5 150 0.2 454.17 203.88 836.12 1.619 Fa max
Cermet 4 1.5 200 0.12 321.63 51.72 516.02 0.702 Fr min
Carbide 3 1.5 250 0.2 399.06 319.58 804.36 1.336 Fr max
Carbide 1 0.5 175 0.12 100.11 94.04 202.54 0.634 Fc min
Carbide 3 1.5 200 0.2 409.48 305.07 859.52 1.315 Fc max
Cermet 6 0.5 250 0.12 74.37 118.83 218.68 0.345 Ra min
Carbide 1 1.5 150 0.2 437.23 110.79 814.09 1.893 Ra max
The cutting-force values were found to influence the
resulting surface-roughness values. When Fa, Fr and Fc
increase, the surface roughness also increases. There is a
direct correlation between the Fa, Fr, Fc values and Ra.
When the resulting surface-roughness values are exami-
ned, it is seen that the surface-roughness values obtained
with the coated cermet cutting tools are lower than those
obtained with coated carbide tools. The resultant force
components of Fa, Fr and Fc have a direct correlation
with Ra. Maximum Fa, Fr and Fc appeared with the coated
carbide cutting tools (Figure 8). However, the minimum
surface-roughness values were obtained with the coated
cermet cutting tools (Figure 9). The lower surface-
roughness values obtained with the coated cermet cutting
tools can be explained with a lower built-up-edge (BUE)
formation tendency of the cermet cutting tools. The pre-
sence of BUE on the cutting tool significantly increases
the surface-roughness values.
4.2 Performance of the ANN model
An experimental design was accomplished. A 27 fully
factorial experimental design was planned. Totally,
8 column × 270 line data was obtained from the experi-
ments. This data was divided into groups. 70 % of the
data was used for the training, 15 % of the data was used
for the validation of the results and 15 % of the data was
used to test the results. A multiple regression analysis
was accomplished (R2 = 0.81810949, adjusted
M. T. OZKAN et al.: EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL NETWORK MODEL ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 227–236 233
Figure 9: Comparison of the surface-roughness values: machining
with coated carbide/cermet cutting tools
Slika 9: Primerjava vrednosti hrapavosti povr{ine: rezanje z orodjem s
karbidno ali kermetno prevleko
Figure 10: ANN model
Slika 10: ANN-model
Table 5: Iterations to find the best ANN model
Tabela 5: Pribli`ki pri iskanju najbolj{ega ANN-modela
Net name Trainingperformance Test performance Training error Test error Training algorithm
MLP 7-26-1 0.942287 0.910443 0.002352 0.003235 BFGS 52
MLP 7-14-1 0.957514 0.913514 0.001748 0.003132 BFGS 54
MLP 7-29-1 0.957661 0.911581 0.001738 0.003189 BFGS 90
MLP 7-26-1 0.940432 0.904966 0.002425 0.003485 BFGS 59
MLP 7-20-1 0.947744 0.907168 0.002135 0.003365 BFGS 73
MLP 7-54-1 0.960559 0.916491 0.001624 0.003013 BFGS 109
MLP 7-26-1 0.943531 0.910626 0.002302 0.003236 BFGS 50
MLP 7-37-1 0.987255 0.952368 0.000532 0.001757 BFGS 132
RBF 7-53-1 0.929539 0.905647 0.002851 0.003411 RBFT
MLP 7-50-1 0.955443 0.904887 0.001829 0.003456 BFGS 67
MLP 7-43-1 0.962073 0.909532 0.001561 0.003283 BFGS 52
MLP 7-74-1 0.944994 0.907058 0.002245 0.003363 BFGS 66
MLP 7-45-1 0.943063 0.908330 0.002321 0.003342 BFGS 52
R2 = 0.81324982). The results obtained with the multiple
regression analyses were not enough to interpret the
experiment results. For this reason, an Annova analysis
was also performed. But Annova did not make any
contribution towards interpreting the experimental
results either. The Statistica software was used for the
statistical analysis. The artificial neural network (ANN)
model could not be improved. A code was developed for
modeling the ANN using Matlab. Different ANN models
were tried (Table 5) and the best ANN model with the
highest learning level was chosen. The code used in the
model resulted in very sound results (RMS =
0.0055829150, R2 = 0.9997334826 and MEP/% =
–0.0000134181). These are the average values of all the
ANN-model results.
There are many commercial ANN software products.
In this study, the Matlab neural network toolbox was
used to obtain a neural network model. There are seven
inputs (type of cutting tool, coated carbide/cermet
cutting tools, depth of cut, cutting speed, feed rate, Fa, Fr,
Fc forces and one output (surface roughness). A multi-
layer feed-forward perceptron (MLP 12-13-1) was used
in the ANN model. The tansig, logsig and purelin
functions were used on the Matlab code and the
Levenberg-Marquardt training method was used in the
ANN model. Figure 10 shows the ANN model.
Figure 11 shows the training performance of the
model. According to the graphs, the training percentage
is a maximum. This shows that the training value can be
acceptable. The resulting value is very close to 1. Figure
12 shows the ANN results: the training, validation and
the test regression-analysis results. The performance of
the ANN model relates to the deviation between the
actual output values and predicted output values. Table 6
shows a comparison of the experimental values and the
Matlab neural-network predictions. The error amounts
were used for the analysis of three statistical values.
M. T. OZKAN et al.: EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL NETWORK MODEL ...
234 Materiali in tehnologije / Materials and technology 48 (2014) 2, 227–236
Figure 12: ANN results: training, validation and the test
Slika 12: Rezultati ANN: usposabljanje, ocena in preizkus
Figure 11: ANN training result
Slika 11: Rezultati ANN-usposabljanja
Table 6: ANN test data
Tabela 6: Podatki ANN-preizkusa
Cutting
tool
(z)
Depth
of cut
d/mm
Cutting
speed
v/(mm/min)
Feed rate
f/(mm/r) Fx (Fa)/N Fy(Fr)/N Fz (Fc)/N
Experimental
surface rough-
ness (Ra/μm)
ANN model
test
(MATLAB)
RMS R2 MEP
1 0.5 150 0.12 118.5 95.61 224.01 0.708 0.703709 0.004291 0.999963 0.006097
1 0.5 150 0.16 130.96 118.18 280.97 1.046 1.041602 0.004398 0.999982 0.004223
2 1 175 0.12 235.44 163.89 403.76 0.603 0.604894 0.001894 0.99999 –0.00313
2 1 175 0.16 221.03 164.91 474.29 0.866 0.868187 0.002187 0.999994 –0.00252
2 1 175 0.2 237.59 177.01 567.95 1.106 1.103077 0.002923 0.999993 0.00265
3 0.5 200 0.16 83.21 155.97 249.81 0.711 0.713359 0.002359 0.999989 –0.00331
3 0.5 200 0.2 97.96 173.58 310.1 1.018 1.014781 0.003219 0.99999 0.003172
3 0.5 225 0.12 86.49 148.89 226.71 0.465 0.470102 0.005102 0.999882 –0.01085
4 0.5 225 0.2 123.95 102.78 316.82 1.515 1.502162 0.012838 0.999927 0.008546
4 0.5 250 0.12 113.08 82.18 226.48 0.596 0.582287 0.013713 0.999445 0.02355
4 0.5 250 0.2 121.36 121.4 301.84 1.568 1.565077 0.002923 0.999997 0.001868
5 0.5 250 0.2 92.58 126.44 292.2 1.387 1.385913 0.001087 0.999999 0.000784
5 1 150 0.12 204.73 134.97 394.89 0.637 0.636777 0.000223 0.999999 0.000351
Figure 13 shows ANN function fit for output. Figure 14
shows ANN error histogram (30 Bins and 20 Bins).
These are the statistical errors of the RMSE (the root
mean square error), R2 (the absolute fraction of variance)
and the MEP (the mean error percentage). If these error
amounts are calculated according to the value of the
output surface roughness, the RMSE is found to be
smaller than 0.004291, R2 is 0.999963 and the MEP is
around 0.006097 for the training and test data.
The RMSE, R2 and MEP values are obtainable with
the following equations:
RMSE = ( )1 2
p
j t oj j∑ − (3)
RMSE = 0 708 0 703709
2
. .− = 0.004291
The statistical error amount:
R
j t o
j o
R
j j
j
2
2
2
2
2
1
1
0 708 0 703709
0 703
= −
−
= −
−
∑
∑
( )
( )
( . . )
( . 709
0 9999632)
.=
(4)
and the average percent error:
MEP/% =
j
t o
t
p
j j
j
−
⋅
⎛
⎝
⎜
⎞
⎠
⎟∑ 100
(5)
MEP/% =
0 708 0 703709
0 703709
100 0 006097
. .
.
.
−
⋅ =
values are obtained, where t is the target value, o is the
output and p is the number of the samples.
5 CONCLUSIONS
An ANN was developed for predicting the surface-
roughness values in turning the 50CrV4 (SAE 6150)
steel using the cutting forces (Fa, Fr and Fc) and
machining parameters. The influences of the machining
parameters (cutting speed, feed rate, depth of cut and
cutting-tool material) on the cutting forces and surface
roughness were investigated. The optimum configuration
of the ANN consisted of three layers with the LM neural
network approach. The statistical-analysis results are:
RMSE = 0.004291, R2 = 0.999963, MEP/% =0.006097.
The predicted surface-roughness values using the ANN
model are in good agreement with the experimentally
obtained surface-roughness values. The ANN based on
the calculation can be used to predict the surface rough-
ness depending on the machining parameters. These
results can be used to predict the cutting forces and
surface roughness in machining the 50CrV4 (SAE 6150)
steel using the coated carbide and cermet cutting tools.
In conclusion, a surface-roughness prediction not
requiring an experimental study with ANN models can
provide both simplicity and fast calculation. It is shown
that an ANN model can be used as an effective and
alternative method of experimental studies improving
both the time and economical optimization of the
machining.
M. T. OZKAN et al.: EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL NETWORK MODEL ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 227–236 235
Figure 14: ANN error histograms, 30 Bins and 20 Bins
Slika 14: Histogram ANN-napak, 30 Bins in 20 Bins
Figure 13: ANN function fit for the output
Slika 13: ANN-funkcja, primerna za izhod
6 REFERENCES
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M. T. OZKAN et al.: EXPERIMENTAL DESIGN AND ARTIFICIAL NEURAL NETWORK MODEL ...
236 Materiali in tehnologije / Materials and technology 48 (2014) 2, 227–236
Z. NAIT ABDELLAH, M. KEDDAM: ESTIMATION OF THE BORON DIFFUSION COEFFICIENTS IN FeB AND Fe2B LAYERS ...
ESTIMATION OF THE BORON DIFFUSION
COEFFICIENTS IN FeB AND Fe2B LAYERS DURING
THE PACK-BORIDING OF A HIGH-ALLOY STEEL
DOLO^ANJE KOEFICIENTA DIFUZIJE BORA V PLASTEH FeB IN
Fe2B MED BORIRANJEM VISOKO LEGIRANEGA JEKLA V
SKRINJI
Zahra Nait Abdellah1,2, Mourad Keddam1
1Laboratoire de Technologie des Matériaux, Département de Sciences des Matériaux, Faculté de Génie Mécanique et Génie des Procédés,
USTHB, B.P N°32, 16111 El-Alia, Bab-Ezzouar, Alger, Algérie
2Département de Chimie, Faculté des sciences, Université Mouloud Mammeri, 15000 Tizi-Ouzou, Algérie
keddam@yahoo.fr
Prejem rokopisa – received: 2013-03-31; sprejem za objavo – accepted for publication: 2013-06-07
In this work we propose a diffusion model to estimate the boron diffusion coefficients in FeB and Fe2B layers during the
pack-boriding of AISI M2 steel in the temperature range 1173–1323 K for a treatment time of 4–8 h.
The proposed model is based on the mass-balance equations at the two growth fronts – FeB/Fe2B and Fe2B/substrate – under
certain assumptions. The estimated values of the boron activation energies in the FeB and Fe2B layers were compared with the
literature data. The present model was extended to predict the thickness of each boride layer for the borided samples at different
temperatures for 10 h. Iso-thickness diagrams were established to be used as a tool for predicting the thickness of each boride
layer as a function of the two parameters: temperature and time. Finally, a simple equation was proposed to estimate the
required time to obtain a single Fe2B layer by diffusion annealing.
Keywords: boriding, incubation times, Fick’s laws, simulation, growth kinetics, annealing
Predstavljeno delo predlaga model difuzije za dolo~anje koeficienta difuzije bora v plasteh FeB in Fe2B med boriranjem v
skrinji jekla AISI M2 v temperaturnem obmo~ju 1173–1323 K pri spreminjanju trajanja postopka od 4 h do 8 h.
Predlagani model temelji na ena~bi masne balance na dveh rasto~ih mejnih ploskvah (FeB/Fe2B) in (Fe2B/osnova) pri dolo~enih
predpostavkah. Dolo~ena vrednost aktivacijske energije bora v FeB- in Fe2B-plasti je bila primerjana s podatki iz literature.
Predstavljeni model je bil raz{irjen, da bi lahko napovedal debelino vsake od obeh boridnih plasti za borirane vzorce pri
razli~nih temperaturah in trajanju do 10 h. Postavljeni so bili diagrami enake debeline, ki so uporabni kot orodje za
napovedovanje debeline vsakega od boriranih slojev v odvisnosti od dveh parametrov (temperature in ~asa). Predlagana je
preprosta ena~ba za dolo~anje potrebnega ~asa za nastanek plasti Fe2B z difuzijskim `arjenjem.
Klju~ne besede: boriranje, inkubacijski ~as, Fickovi zakoni, simulacija, kinetika rasti, `arjenje
1 INTRODUCTION
One of the surface-modification methods for improv-
ing the surface properties of ferrous and non-ferrous
alloys is boriding. According to the Fe-B binary system,
two kinds of iron borides, i.e., FeB and Fe2B, with a
narrow range of composition can be identified.1 The
boriding process applies in the temperature range
1073–1323 K between 1 h to 10 h and it can be carried
out in solid, liquid or gaseous media. The possible for-
mation of the FeB and Fe2B iron borides depends upon
various factors, such as the boron activity of the boriding
medium, the chemical composition of the substrate, the
process temperature and the treatment time. The mor-
phology of the boride layers is influenced by the pre-
sence of alloying elements in the matrix. Saw-tooth-
shaped layers are obtained in low-alloy steels, whereas in
high-alloy steels, the interfaces tend to be flat. The
modelling of the boriding kinetics is considered as a
suitable tool to match the case depth with the intended
industrial applications for this borided steel. So, the
modelling of the growth kinetics for boride layers has
gained much attention to simulate the boriding kinetics
during recent decades.2–26
In the present work an original diffusion model is
proposed to estimate the boron diffusion coefficients in
the FeB and Fe2B layers grown on AISI M2 steel by
considering the boride incubation times. A non-linear
boron-concentration profile is assumed through the
boride layers. The mass-balance equations were applied
to the two diffusion fronts: the FeB/Fe2B and Fe2B/sub-
strate interfaces in the temperature range 1173–1323 K.
In addition, a simple equation was proposed to estimate
the required time to obtain a single Fe2B layer by
diffusion annealing.
2 THE DIFFUSION MODEL
The model takes into account the FeB/Fe2B bilayer
growth on the saturated substrate with boron atoms, as
shown in Figure 1.
C up
FeB and C low
FeB (= 16.23 % B) are the upper and lower
boron mass concentrations in the FeB, while C up
FeB (= 9 %
B) and C low
FeB (= 8.83 % B) are, respectively, the upper
Materiali in tehnologije / Materials and technology 48 (2014) 2, 237–242 237
UDK 621.793:669.14 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)237(2014)
and lower boron concentrations in the Fe2B. Cads denotes
the adsorbed concentration of boron,15 while u is the
position of the FeB/Fe2B interface, and v is the position
of the Fe2B/substrate interface. C0 is the boron solubility
in the matrix and is equal to 35 · 10–4 % B.2 The upper
boron content in the FeB phase (C up
FeB ), imposed by the
boriding medium, gives rise to the two iron borides: FeB
and Fe2B. From a thermodynamic point of view, the FeB
phase exhibits a narrow composition range (of about the
mole fraction x = 1 % B or the mass fraction w = 0.2 %
B), as identified by Massalski.27 The upper boron content
in the FeB phase was taken in the composition range of
mass fractions 16.25–16.43 % B to obtain a bilayer
configuration consisting of the two iron borides, FeB and
Fe2B.
The following assumptions are considered during the
formulation of the diffusion model:
• The kinetics is dominated by the diffusion-controlled
mechanism
• The growth of the boride layers is a consequence of
the boron diffusion perpendicular to the sample sur-
face
• The range of homogeneity of the iron borides is
about x = 1 % B
• The iron borides nucleate after a certain incubation
time
• The boride layer is thin in comparison to the sample
thickness
• A local equilibrium occurs at the phase interfaces
• A planar morphology is assumed for the phase inter-
faces
• The volume change during the phase transformation
is ignored
• The diffusion coefficient of boron in each iron boride
does not vary with the boron concentration and
follows an Arrhenius relationship
• A uniform temperature is assumed throughout the
sample
• The alloying elements have no effect on the boron
diffusion
• The presence of porosity is neglected during the
boron diffusion.
The initial conditions of the diffusion problem are set
up as follows:
{ }
{ }
{ }
C x t
C x t
C x t
FeB
Fe B
Fe
2
( )
( )
( )



0 0 0
0 0 0
0 0 0
= =
= =
= =
(1)
The boundary conditions are given by the following
equations:
{ }C x t t T CFeB FeB upFeB[ ]= = =0 0( ) for C wtads B16 23. .% (2)
{ }C x t t T CFeB FeB lowFeB[ ]= = =0 0( ) for C wtads B 16 23. .%
and with the FeB phase: (3)
{ }C x t t T CFe B Fe B upFe B2 2 2[ ]= = =0 0( ) for
883 16 23. .% . .%wt C wtB Bads  and without the FeB
phase: (4)
{ }C x t t T CFe B Fe B lowFe B2 2 2[ ]= = =0 0( ) for C wtads B 883. .%
and without the FeB phase: (5)
C x t t u CFeB low
FeB( ( ) )= = = (6)
C x t t u CFe B up
Fe B
2
2( ( ) )= = = (7)
C x t t v CFe B low
Fe B
2
2( ( ) )= = = (8)
C x t t v CFe 0( ( ) )= = = (9)
The mass-balance equations28 are given by the equa-
tions (10) and (11):
[ ]w u
t
J J
x uFeB B
FeB
B
Fe Bd
d
2⎛
⎝
⎜ ⎞
⎠
⎟ = −
=
(10)
[ ]w v
t
w
u
t
J
x vFe B B
Fe B
2
2
d
d
d
d
⎛
⎝
⎜ ⎞
⎠
⎟ + ⎛⎝
⎜ ⎞
⎠
⎟ =
=
' (11)
with
[ ]w C C C C
w
FeB up
FeB
low
FeB
low
FeB
up
Fe B
Fe B
2
2
= × − + −
=
05
0
. ( ) ( )
[ ]. ( ) ( )
' . (
5
05
× − + −
= ×
C C C C
w C
up
Fe B
low
Fe B
low
Fe B
0
up
Fe B
2 2 2
2 −C low
Fe B2 )
The boron flux through a given boride layer is
obtained from Fick’s first law as follows:
J D
C x t
x
i i i
B B= −
∂
∂
( , )
with i = (FeB or Fe2B) (12)
DB
FeB and DB
Fe B2 are, respectively, the diffusion coeffi-
cients of boron in the FeB and Fe2B phases. The boron
concentration profile in the FeB layer is given by:
C x t C
C C
erf
u
D t
FeB up
FeB low
FeB
up
FeB
B
FeB
( , )
( )
= +
−
⎛
⎝
⎜
⎜
⎞
⎠2
⎟
⎟
⋅
⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟erf
x
D t2 B
FeB
For 0 ≤ ≤x u (13)
In the same way, the boron concentration profile in
the Fe2B layer can be obtained as follows:
238 Materiali in tehnologije / Materials and technology 48 (2014) 2, 237–242
Z. NAIT ABDELLAH, M. KEDDAM: ESTIMATION OF THE BORON DIFFUSION COEFFICIENTS IN FeB AND Fe2B LAYERS ...
Figure 1: Boron concentration profile through the FeB/Fe2B bilayer
Slika 1: Profil koncentracije bora skozi plasti (FeB/Fe2B)
C x t C
C C
erf
u
D t
Fe B up
Fe B low
Fe B
up
Fe B
B
Fe B
2
2
2 2
2
( , )
( )
= +
−
⎛
2⎝
⎜
⎜
⎞
⎠
⎟
⎟ −
⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟
⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥
⋅
⋅
erf
v
D t
erf
u
D t
2
2
B
Fe B
B
Fe B
2
2
⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟ −
⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟
⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥
erf
x
D t2 B
Fe B2
For u x v≤ ≤ (14)
The FeB layer thickness u grows parabolically
according to equation (15), where kFeB represents the
parabolic growth constant at the FeB/Fe2B interface:
[ ]u k t t T= −FeB 0FeB ( )
/1 2
(15)
The distance v is the location of the Fe2B/substrate
interface and k its parabolic growth constant (equation
(16)) and the difference (l = v – u) denotes the layer
thickness of the Fe2B (equation 17):
[ ]v k t t T= − 0 ( )
/1 2
(16)
[ ] [ ]l v u k t t T k t t T= − = − − −0 FeB 0FeB( ) ( )
/ /1 2 1 2
(17)
with t T t T0
FeB ( ) ( ) 0 and k k FeB
where t0(T) is the boride incubation time of the total
boride layer and t T0
FeB ( ) is the boride incubation time of
the FeB layer. To take into account the effect of the
boride incubation times when solving the mass-balance
equations, it is necessary to define the two parameters
FeB(T) and (T):
FeB
0
FeB
( )
( )
.
T
t T
t
= −
⎡
⎣⎢
⎤
⎦⎥
1
0 5
(18)
and ( )
( )
.
T
t T
t
= −
⎡
⎣
⎢
⎤
⎦
⎥1
0 5
0
(19)
The layer thickness of the FeB (u) is related to the
FeB(T) parameter by equation (20):
u k T t= FeB FeB ( ) (20)
In the same way, the layer thickness of the Fe2B (l) is
expressed using equation (21):
[ ]l k T k T t= − ( ) ( )FeB FeB (21)
3 ESTIMATION OF THE BORON DIFFUSION
COEFFICIENTS IN THE FeB AND Fe2B LAYERS
To estimate the boron diffusion coefficients in the
FeB and Fe2B layers, the experimental results published
by Campos-Silva et al.29 on borided AISI M2 steel were
used. In this reference work, the powder-pack boriding
was carried out at four temperatures, (1173, 1223, 1273
and 1323) K, for three exposure times, (4, 6 and 8) h,
using the B4C Durborid as a boriding medium. Eighty
measurements were performed on different cross-sec-
tions of the borided samples from the AISI M2 steel to
determine the thickness of each boride layer.
Tables 1 and 2 list the experimental parabolic growth
constants for each phase interface with the correspond-
ing incubation times. The experimental values of the
parabolic growth constants at each phase interface were
obtained from the slopes of the curves relating the
squared boride layer thickness to the boriding time. The
boride incubation times were deduced for a null boride
layer thickness.
Table 1: Experimental values of the parabolic growth constants at the
FeB/Fe2B interface in the temperature range 1173–1323 K with the
corresponding boride incubation times
Tabela 1: Eksperimentalne vrednosti konstant paraboli~ne rasti na
stiku (FeB/Fe2B) v temperaturnem obmo~ju 1173–1323 K, z ustrez-
nim inkubacijskim ~asom borida
T/K Experimental growthConstants: kFeB/(ìm s–1/2)
t T0
FeB( )/s
1173
1223
1273
1323
0.065
0.121
0.179
0.238
10131
6085.7
4347.8
3815.5
Table 2: Experimental values of the parabolic growth constants at the
Fe2B/substrate interface in the temperature range 1173–1323 K with
the corresponding boride incubation times
Tabela 2: Eksperimentalne vrednosti konstant paraboli~ne rasti na
stiku (Fe2B/podlaga) v temperaturnem obmo~ju 1173–1323 K, z
ustreznim inkubacijskim ~asom borida
T/K Experimental growthConstants: k/(μm s–1/2) t0(T)/s
1173
1223
1273
1323
0.168
0.305
0.448
0.589
8806.2
4729
4323
3742.7
It was demonstrated that the higher boriding tem-
peratures involve the shorter incubation times,24 as
shown in Tables 1 and 2. The two respective parameters
FeB(T) and (T) are linearly dependent on the boriding
temperature and can be approximated by equations (22)
and (23) from a linear fitting of the experimental data
displayed in Figure 2:
FeB ( ) ( . )T T= ⋅ × −
−1 39 10 085793 (22)
and ( ) ( . )T T= ⋅ × −−1 40 10 084703 (23)
For this purpose, a computer code written in Matlab
(version 6.5) was used to estimate the boron diffusivity
in each boride layer. This program requires the following
input data: the time, the temperature, the lower and upper
boron concentrations at each phase interface as well as
the two parameters FeB(T) and (T). By solving the
mass-balance equations (equations (10) and (11)) via the
Newton-Raphson method,30 it is possible to determine
the boron diffusion coefficients in the FeB and Fe2B
layers. Table 3 summarizes the estimated values of the
boron diffusion coefficients in the FeB and Fe2B layers
for an upper boron content equal to w = 16.40 % in the
FeB phase.
Figure 3 depicts the temperature dependence of the
boron diffusion coefficients in the FeB and Fe2B layers
according to the Arrhenius equation. The value of the
Z. NAIT ABDELLAH, M. KEDDAM: ESTIMATION OF THE BORON DIFFUSION COEFFICIENTS IN FeB AND Fe2B LAYERS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 237–242 239
boron activation energy in each boride layer can be
easily obtained from the slopes of the corresponding
curves. So, the boron diffusion coefficients in the FeB
and Fe2B layers are, respectively, given by equations (24)
and (25):
D RTB
FeB
kJ /mol
m s= × −
−
−28 10 3
220 2
2 1. exp ( )
.
(24)
D RTB
Fe B
kJ /mol
2 m s= × −
−
−16 10 3
213
2 1. exp ( ) (25)
where R is the universal gas constant (= 8.314 J/(mol K)),
and T represents the absolute temperature in Kelvin.
The reported values of the activation energies24,29,31–33
of the borided steels are listed in Table 4 together with
the values from this work. The obtained values of the
activation energies are found to be dependent on the
boriding method and on the chemical composition of the
substrates.
In Table 5, a comparison was achieved between the
experimental boride layer thicknesses and the simulated
ones at different temperatures for 10 h of treatment. The
simulated results were obtained from equations (20) and
(21). The present model was able to predict the boride
layer thickness (FeB or Fe2B) for the given boriding
conditions.
Z. NAIT ABDELLAH, M. KEDDAM: ESTIMATION OF THE BORON DIFFUSION COEFFICIENTS IN FeB AND Fe2B LAYERS ...
240 Materiali in tehnologije / Materials and technology 48 (2014) 2, 237–242
Figure 3: An Arrhenius relationship between the boron diffusion
coefficient and the temperature: a) FeB layer, b) Fe2B layer
Slika 3: Arrheniusova odvisnost med koeficientom difuzije bora in
temperaturo: a) FeB-plast, b) Fe2B-plast
Figure 2: Evolution of the two parameters as a function of the
boriding temperature: a) FeB(T) and b) (T)
Slika 2: Razvoj dveh parametrov v odvisnosti od temperature bori-
ranja: a) FeB(T) in b) (T)
Table 3: Determination of the boron diffusion coefficient in each
boride layer for an upper boron mass fraction content of w =16.40 %
in the FeB layer
Tabela 3: Dolo~anje koeficienta difuzije bora v vsaki boridni plasti za
zgornji masni dele` vsebnosti bora 16,40 % v plasti FeB
T/K DB
FeB m s( )2 1 1210− −× DB
Fe B2 m s( )2 1 1210− −×
1173
1223
1273
1323
0.376
1.282
2.797
4.915
0.462
1.502
3.227
5.539
Table 4: Values of the boron activation energies obtained for different borided steels
Tabela 4: Vrednosti aktivacijske energije bora, dobljene iz razli~nih boriranih jekel
Material Boriding method Activation energy of FeBE/(kJ mol–1)
Activation energy of Fe2B
E/(kJ mol–1) Reference
AISI M2
AISI4140
AISI H13
AISI 316L
AISI M2
AISI M2
Paste
Paste
Powder-pack
Powder-pack
Powder-pack
Powder-pack
283
–
–
204
223
220.2
239.4
168.5
186.2
198
207
213
32
33
31
24
29
Present study
Table 5: Experimental (exp.) and simulated (sim.) values of the boride
layer thickness in the temperature range 1173–1323 K for 10 h of treat-
ment, with an upper boron mass fraction of content of w = 16.40 % in
the FeB phase
Tabela 5: Eksperimentalne (exp.) in simulirane (sim.) vrednosti za
debelino plasti borida v temperaturnem obmo~ju 1173–1323 K, za
10 h obdelave pri gornjem masnem dele`u vsebnosti bora 16,40 %, v
FeB plasti
T/K FeB (μm)exp.
FeB (μm)
sim.
Fe2B (μm)
exp.
Fe2B (μm)
sim.
1173
1223
1273
1323
10.17
20.98
28.30
40.24
10.30
17.89
29.75
47.60
19.66
32.81
51.83
72.28
16.70
28.23
45.81
71.67
In Table 6, the predicted values of the boride layer
thicknesses are compared with the experimentally
determined values in the temperature range 1173–1323
K for a treatment time varying from 4 h to 8 h. Good
agreement was observed between the experimental data
and the simulation results for an upper boron content
equal to w = 16.40 % in the FeB phase.
Table 6: Experimental (exp.) and simulated values (sim.) of the boride
layer thickness in the temperature range 1173–1323 K for different
treatment times with an upper boron content w = 16.40 % in the FeB
phase
Tabela 6: Eksperimentalne (exp.) in simulirane (sim.) vrednosti debe-
line plasti borida pri temperaturah 1173–1323 K za razli~ne ~ase
obdelave in zgornjo vsebnostjo bora w = 16,40 % v FeB-fazi
T/K Time(h)
FeB (μm)
exp.
FeB (μm)
sim.
Fe2B (μm)
exp.
Fe2B (μm)
sim.
1173
4 4.24 6.47 8.31 10.53
6 6.96 7.92 12.04 12.89
8 8.88 9.15 14.87 14.89
1223
4 11.03 11.32 18.96 17.85
6 15.07 13.86 24.54 21.86
8 18.23 16.00 29.08 25.24
1273
4 17.94 18.81 27.02 28.97
6 23.51 23.04 35.37 35.48
8 28.00 26.60 42.09 40.97
1323
4 24.48 24.89 36.31 35.90
6 31.73 32.27 46.97 46.43
8 37.61 38.25 55.61 54.98
Figure 4 displays the iso-thickness diagrams des-
cribing the evolution of the boride layer thickness as a
function of the time and the boriding temperature. The
results derived from Figure 4 can be used as a tool to
predict the boride layer thickness in relation with its
practical use in an industrial area.
4 OBTAINING OF A SINGLE LAYER OF Fe2B
BY DIFFUSION ANNEALING
In industrial practice it is possible to reduce the
brittleness of boride layers by controlling their micro-
structure. It is known that a single Fe2B boride layer is
more desirable than a dual FeB-Fe2B layer.34 This makes
it possible to reduce the FeB layer thickness by applying
a diffusion annealing in a hydrogen atmosphere. During
this stage, the supply of boron is stopped since the
concentration gradient of boron in the FeB is null (i.e.,
C up
FeB = C low
FeB = 16.23 %), the FeB layer will be converted
into an Fe2B layer. The time required to eliminate the
FeB layer during the diffusion annealing can be obtained
from equation (26):
t
u l C C
D C C
uFeB =
=
× × −
−0
( )
(
low
FeB
up
Fe B
B
Fe B
up
Fe B
low
Fe
2
2 2 2
B
)
(26)
where u is the FeB layer thickness (μm), l the Fe2B layer
thickness (μm) and DB
Fe B2 represents the boron diffusion
coefficient in Fe2B. It is clear that the annealing time
depends on the boron diffusion coefficient in Fe2B, and
also on the thickness of each boride layer. During the
diffusion annealing, an infinitesimal reduction of the
FeB layer is related to the infinitesimal growth of the
Fe2B layer by equation (27):
Δ Δ Δu
w
w w w
l l= −
+ +
⎛
⎝
⎜⎜
⎞
⎠
⎟⎟ = −
Fe B
FeB Fe B
2
2
'
.05493 (27)
The value of the Fe2B layer thickness l' (μm) after
diffusion annealing becomes:
l l
u
'
.
= +
Δ
05493
(28)
Table 7 presents the simulation results obtained from
equations (26) and (28) to estimate the Fe2B layer
thickness after diffusion annealing and the time required
Z. NAIT ABDELLAH, M. KEDDAM: ESTIMATION OF THE BORON DIFFUSION COEFFICIENTS IN FeB AND Fe2B LAYERS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 237–242 241
Figure 4: Iso-thickness diagrams describing the evolution of the
boride layers: a) FeB, b) Fe2B
Slika 4: Diagram enakih debelin opisuje razvoj boridnih plasti: a)
FeB, b) Fe2B
to eliminate the FeB layer in the case of the borided
samples treated at different temperatures for 10 h. The
obtained annealing times are increased with an increase
of the boriding temperature since the boride layer
becomes thicker. In this context, Kulka et al.26 have
experimentally determined the annealing time using a
hydrogen atmosphere to obtain a single Fe2B layer on
gas borided Armco Fe at 1173 K for 2 h in a gas mixture
(H2–BCl3). They found that the total elimination of the
FeB layer took about 1 h.
Furthermore, it was shown by Dybkov et al.35 that
annealing of a borided Fe–Cr sample for 6 h resulted in
the disappearance of the FeB layer.
Table 7: Estimation of the Fe2B layer thickness and the time required
to eliminate the FeB layer for the borided samples at different tem-
peratures for 10 h
Tabela 7: Dolo~anje debeline plasti Fe2B in ~as, potreben za odpravo
FeB plasti, za vzorce, borirane 10 h pri razli~nih temperaturah
T/K
FeB
(μm)
sim.
Fe2B
(μm)
sim.
Fe2B (μm)
After diffusion
annealing
Equation (28)
Annealing time
tu FeB = 0/h
Equation (26)
1173 10.30 16.70 35.45 4.15
1223 17.89 28.23 60.79 4.99
1273 29.75 45.81 99.96 5.92
1323 47.60 71.67 158.32 6.93
5 CONCLUSIONS
In this work an original diffusion model was pro-
posed to estimate the boron diffusion coefficients in the
FeB and Fe2B layers grown on AISI M2 steel. To deter-
mine the boron activation energy in each boride layer,
the mass-balance equations were formulated, including
the effect of the boride incubation times. The estimated
boron activation energies were compared with the lite-
rature data. The present model was extended to predict
the thickness of each boride layer for the borided sam-
ples at different temperatures for 10 h. Iso-thickness
diagrams were established to be used as a tool to predict
the thickness of each boride layer as a function of the
two parameters (temperature and time). The required
time to obtain a single Fe2B layer by diffusion annealing
was estimated on the basis of a simple equation. The for-
mation of a single Fe2B layer on AISI M2 steel depended
on the boriding parameters.
Acknowledgements
This work was carried out in the framework of the
CNEPRU project under code number J0300220100093
of the Algerian Ministry of Higher Education and Scien-
tific Research.
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Z. NAIT ABDELLAH, M. KEDDAM: ESTIMATION OF THE BORON DIFFUSION COEFFICIENTS IN FeB AND Fe2B LAYERS ...
242 Materiali in tehnologije / Materials and technology 48 (2014) 2, 237–242
O. CULHA: FINITE ELEMENT MODELLING OF SUBMERGED ARC WELDING PROCESS FOR A SYMMETRIC T-BEAM
FINITE ELEMENT MODELLING OF SUBMERGED ARC
WELDING PROCESS FOR A SYMMETRIC T-BEAM
MODELIRANJE POSTOPKA OBLO^NEGA VARJENJA POD
PRA[KOM SIMETRI^NEGA T-NOSILCA Z METODO KON^NIH
ELEMENTOV
Osman Culha
Celal Bayar University, Engineering Faculty, Department of Materials Engineering, Muradiye Campus, Manisa, Turkey
osman.culha@cbu.edu.tr
Prejem rokopisa – received: 2013-04-11; sprejem za objavo – accepted for publication: 2013-06-10
Metallurgical welding joints are extensively used in the fabrication industry, including ships, offshore structures, steel bridges
and pressure vessels. The merits of such welded structures include a high joint efficiency, water and air tightness, and low
fabrication costs. However, residual stresses and distortions can occur near the weld bead due to localized heating by the
welding process and subsequent rapid cooling. This paper is focused on deriving a simulation solution to predict the design
parameters, such as the temperature-stress distribution, the approximate gradient and the nodal displacement on the plates
during the process of submerged arc welding (SAW). During the construction of an AH 36 quality T-beam profile using the
SAW process, thermal residual stress and distortion occurs due to heat fusion from the source to the joint part of the symmetric
T-beam. The value of the design parameter is achieved by performing a thermal elasto-plastic analysis using finite-element
techniques. Furthermore, this investigation provides an available process analysis to enhance the fabrication process of welded
structures.
Keywords: submerged arc welding (SAW), finite element modelling (FEM), stress-temperature distribution
Metalur{ko varjeni spoji se splo{no uporabljajo v industriji, vklju~no z ladjedelni{tvom, za naftne plo{~adi, jeklene mostove in
tla~ne posode. Prednosti takih varjenih konstrukcij so velika zmogljivost spoja, tesnost za vodo in zrak, nizki proizvodni stro{ki.
Vendar pa se v okolici kopeli zvara zaradi lokalnega segrevanja in hitrega ohlajanja lahko pojavijo zaostale napetosti in izkriv-
ljanje. ^lanek je osredinjen na izpeljavo re{itve simulacije za predvidenje parametrov kot so: razporeditev temperature – nape-
tosti, pribli`ek gradienta in izkrivljanja vozlov na plo{~ah med postopkom oblo~nega varjenja pod pra{kom (SAW). Med
sestavljanjem AH 36 kvalitete T-nosilca s postopkom SAW se pojavijo termi~ne zaostale napetosti in izkrivljanja zaradi toplote
zlivanja od vira do spojnega mesta simetri~nega T-nosilca. Vrednosti parametrov so dobljene z izvajanjem termi~ne elasto-
plasti~ne analize s tehniko kon~nih elementov. Poleg tega ta preiskava omogo~a analizo procesa za pospe{itev izdelave varjenih
konstrukcij.
Klju~ne besede: oblo~no varjenje pod pra{kom (SAW), modeliranje z metodo kon~nih elementov (FEM), razporeditev napetosti
– temperature
1 INTRODUCTION
Metallurgical welding joints are extensively used in
the fabrication industry, which includes ships, offshore
structures, steel bridges and pressure vessels. Among the
merits of such welded structures are a high joint effi-
ciency, water and air tightness, and low fabrication costs.
However, residual stresses and distortions can occur near
the weld bead due to localized heating by the welding
process and subsequent rapid cooling. Additionally,
phase transformations that occur in the weld metal and
adjacent heat-affected zone (HAZ), e.g., in structural
steels, contribute to the evolution of residual stress.
Stress regions near the weld may promote brittle fractu-
res, fatigue, or stress-corrosion cracking. Meanwhile,
residual stresses in the base plate may reduce the
buckling strength of the structure members.1 The inhe-
rent characteristics of weldments, such as metallurgical
or geometrical defects, the presence of stress raisers and
heterogeneous material properties, make them particu-
larly vulnerable to failure. In material selection it is
general practice to specify weld metals with strengths
higher than those of the parent plate. This overmatching
policy is thought to be able to offset potential problems,
such as a reduced toughness and the presence of defects
in the weld metals.2 Therefore, the residual stresses of
welding must be minimized to control them according to
the respective requirements. Previous investigators deve-
loped several methods, including heat treatment,
hammering, preheating, vibration stress relieving, and
weld sequencing, to reduce the residual stresses attri-
buted to welding.3–5 As known, many welded structures
that cannot be subjected to post-weld manufacturing
measures after the welding contain residual stresses of
varying degree that can result in unintended deforma-
tions of the welded component, increase the suscepti-
bility to hydrogen-induced cold cracking, and also
combine with tensile stresses experienced during service
to promote brittle fracture, fatigue failure, and stress-
corrosion cracking. Thus, developing an available
welding sequence and accurately predicting the welding
residual stresses for a welds system are necessary in
order to achieve the safest design.1–9
Materiali in tehnologije / Materials and technology 48 (2014) 2, 243–248 243
UDK 621.791.75 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)243(2014)
On the other hand, traditional methods for welding-
induced residual stress and strain characterization are
mainly experimental, and include hole drilling, X-ray,
neutron diffraction, ultrasonic and demountable mecha-
nical gauge measurement. However, the application of
these methods in practice is usually limited by either
their cost or accuracy. Numerical simulations based on
finite-element modelling are used to study the influence
of welding sequences on the distribution of the residual
stress and distortion generated when welding a flat-bar
stiffener to a steel plate. The simulation consists of
sequentially coupled thermal and structural analyses
using an element birth-and-death technique to model the
addition of weld metal to the workpiece. The tempe-
rature field during welding and the welding-induced
residual stress and distortion fields are predicted and the
results are compared with experimental measurements
and analytical predictions.10,11
The intense heat input from the welding process and
the molten steel, which is deposited into the joint, con-
tribute to the thermal expansion and contraction of the
heated parts. Also, residual stresses released in the rolled
and raw-cut material will have some influence on the
final shape of the beam. If the welding is carried out on
only one side of the web the plate will bend upwards and
sideways (sweep and camber). If two welds are made
simultaneously, one on each side of the web, the beam
tends only to exhibit beam camber due to the shrinkage
produced after welding. Distortion can be minimized in
symmetrical designs by applying the heat into the mate-
rial in a symmetrical manner. The remaining distortion
on the final part can rarely be tolerated in subsequent
assemblies and must be eliminated with either thermal or
mechanical methods. Both mechanical and thermal
straightening techniques are used to straighten the beams
subjected to plate shrinkage. Mechanical straightening
requires large, expensive machines, which produce pro-
cess "wrinkles". As with thermal straightening, this tech-
nique adds another step to the production line for beams.
Thermal straightening requires that the opposite (upper)
side of the beam is heated – which later induces a com-
pensating shrinkage. This is supposed to compensate for
the initial shrinkage induced during the welding process
mentioned above and can be carried out with a propane
burner or induction line heating. In accordance with
these explanations, the aims of the study were to obtain
the temperature, stress and displacement region of a sub-
merged arc welding (SAW) process for welded T-beam
dimensions of 180 mm × 28 mm and 630 mm × 14 mm,
AH 36 quality steel bars using the finite-element method.
2 SUBMERGED ARC WELDING (SAW)
PROCEDURES FOR FLAT BARS
Submerged arc welding is a method in which the heat
required to fuse the metal is generated by an electric
current passing between the welding wire and the work-
piece. The tip of the welding wire, the arc and the weld
area are covered by a layer of granular flux. A hopper
and a feeding mechanism are used to provide a flow of
flux over the joint being welded. A conveyor tube is
244 Materiali in tehnologije / Materials and technology 48 (2014) 2, 243–248
O. CULHA: FINITE ELEMENT MODELLING OF SUBMERGED ARC WELDING PROCESS FOR A SYMMETRIC T-BEAM
Figure 1: Schematics of submerged arc welding process
Slika 1: Shematski prikaz oblo~nega varjenja pod pra{kom
Figure 2: Dimensional representation of steel bars for finite-element
model analysis
Slika 2: Predstavitev dimenzij jeklenih palic, uporabljenih za analizo
z metodo kon~nih elementov
Table 1: Process parameters of submerged arc welding (SAW)
Tabela 1: Procesni parametri pri oblo~nem varjenju pod pra{kom
(SAW)
Wire diameter
(mm)
Amperage
(A)
Voltage
(V)
Welding speed
(mm/min)
2.4 780 28 740
provided to control the flow of the flux and is always
kept ahead of the weld zone to ensure an adequate
supply of flux ahead of the arc. The intense heat evolved
by the passage of the electric current through the
welding zone melts the end of the wire and the adjacent
edges of the workpieces, creating a puddle of molten
metal. The puddle is in a liquid state and is turbulent. For
this reason any slag or gas bubble is quickly swept to the
surface. The flux completely shields the welding zone
from any contact with the atmosphere. Additionally,
SAW can use a much higher heat input and has slower
solidification and cooling characteristics. Also, the
silicon content will be much higher in submerged arc
welding if care is not exercised in selecting the proper
flux material. SAW can be used for the welding of mate-
rials in higher gauges and has the advantage of a high
weld-metal quality and a smooth and uniform weld
finish. The deposit rate, deposition efficiency and the
weld speed are high. This means that smoke and arc
flash are absent in this procedure (Figure 1).
According to these advantages of the SAW process, it
was applied to two different steel plates (as shown in
Figure 2; 180 mm × 28 mm and 630 mm × 14 mm) at
Özkan Iron Steel Industry Co., Ltd. for the production of
a symmetrical T-beam process. The SAW production-
line parameters such as wire diameter, amperage, voltage
and welding speed are listed Table 1. In particular, since
the heat input amount per millimetre was affected by
these parameters, they were controlled and modified
depending on the dimensions of the steel plate.
3 FINITE-ELEMENT MODELLING OF
WELDING FLAT BARS
The simulation technologies were developed in
accordance with the development of computers. In many
applications, the numerical simulation is applied to save
costs. In the field of the welding, the numerical simu-
lations were proposed and applied to the practical field in
order to determine the welding conditions. The accuracy
of the dimensions after the welding of the steel structure
becomes an important factor for the product costs. The
dimensions become inaccurate due to the welding distor-
tion. Therefore, the control of the welding distortion is
demanded in the steel structure welding so as to improve
the productivity. For this purpose, an estimation of the
amount of deformation is needed and its behaviour is
investigated in this study.
The finite-element modelling (FEM) procedure for
SAW was used to determine: a) the thermal stress forma-
tion and distribution, b) the temperature gradient, c) the
deformation characteristics and d) the thermal straighten-
ing regions of the produced T-beam by welding. The
SAW process under consideration was applied for the
manufacture of a T-beam 180 mm × 28 mm and 630 mm
× 14 mm steel plates, as shown in Figure 3. A two-
dimensional (2D) FE model was used to investigate the
detailed residual stress and strain distributions around
the welded region of the T-beam. Therefore, the FE
model 2D coupled temperature-displacement transient
model was structured for the dimensions of the real steel
plates. Additionally, the same thermal properties of the
base and weld metals, such as the thermal conductivity,
convection, expansion and the specific heat capacity,
were taken for this analysis.
On the other hand, the yield strengths of the base and
weld metals were taken at room temperature. The
O. CULHA: FINITE ELEMENT MODELLING OF SUBMERGED ARC WELDING PROCESS FOR A SYMMETRIC T-BEAM
Materiali in tehnologije / Materials and technology 48 (2014) 2, 243–248 245
Figure 3: 2D model assembly of FE analysis
Slika 3: 2D modelni sestav za FE-analizo
Table 2: Composition of the steel used in the experiments
Tabela 2: Sestava jekla, uporabljenega pri poskusu
Material C Si Mn Cu Al V P S
AH 36 0.14 0.20 0.90 0.25 0.027 0.07 0.028 0.030
Table 3: Thermal and mechanical properties of AH 36 quality steel12
Tabela 3: Toplotne in mehanske lastnosti jekla AH 3612
Te
m
pe
ra
tu
re
,
o C E
la
st
ic
M
od
ul
us
,
G
Pa
P
oi
ss
on
ra
ti
o,

T
he
rm
al
C
on
du
c-
tiv
it
y,
K
/W
/(
m
°C
)
S
pe
ci
fi
c
he
at
,
J/
(k
g
°C
)
T
he
rm
al
ex
pa
ns
io
n
co
ef
fi
ci
en
t,

/(
10
–6
/°
C
)
T
he
rm
al
co
nv
ec
ti
on
co
ef
fi
ci
en
t,
h/
W
/(
m
2
°C
)
20 207 0.30 52 485 11.8 3.2
100 202 0.31 50 486 11.8 5.5
200 200 0.33 48 495 12.1 6.3
300 198 0.34 45 513 12.7 6.8
400 181 0.36 42 532 13.2 7.4
500 112 0.38 38 555 13.7 7.7
600 65 0.40 34 586 14.2 7.8
700 42 0.42 30 636 14.7 8.1
800 33 0.44 27 683 14.8 8.3
900 24 0.46 26 698 14.7 8.5
1000 13 0.48 28 698 14.7 8.6
1100 7 0.48 30 698 14.7 8.7
1200 7 0.48 31 698 14.7 8.8
1300 7 0.47 32 698 14.7 8.9
1400 7 0.47 34 698 14.7 9.0
1450 7 0.47 35 698 14.7 9.1
1500 7 0.47 95 698 14.7 9.2
Young’s modulus and Poisson’s ratio variations were
assumed to be the same in the base and weld metals
between 40 °C and 1450 °C.10,12 The Von Mises yield
criterion and the associated flow rule were used together
with kinematic hardening and a bilinear representation of
the stress – strain curve. Kinematic hardening is believed
to be the best model that can simulate the reverse
plasticity and Bauschinger effect that is expected to
occur during multipass welding. The density of both
materials was assumed to be constant at a value of 7800
kg/m3. The composition of the AH 36 quality steel plates
and weld metal used in the T-beam production is shown
in Table 2 and the properties used for the modelling of
the temperature distributions and distortions are listed in
Table 3.
Furthermore, creating the mesh design of the entire
model was very important, especially when the problem
had different thermal regions. So, the elements were
finest in the welded area and became coarser away from
the steel plates. In Figure 4, the welded regions of the
steel plates were partitioned and meshed as the struc-
tured elements type and numbered as CPE4T: A4-node
plane strain thermally coupled quadrilateral, bilinear
displacement and temperature with 2000. In the region
away from the location of the weld, the temperature gra-
dient is significantly lower and the mesh was then made
less dense in this region to reduce the number of degrees
of freedom and thus the time required for the solution.
4 RESULTS AND DISCUSSION
4.1 Thermal history
A temperature-displacement analysis of the symme-
trical T-beam welding with the given welding condition
was performed using the 2D finite-element method.
During this step, the temperature histories for each node
were computed during the multipass welding process.
Ogawa et al.13 states that the weld pass can be divided
into a number of small parts (mesh blocks) with the same
length to simulate the deposition of the weld metal. In
each weld pass, the volume of the moving heat source is
equal to that of these elements composing the corres-
ponding weld bead in one mesh block. The elements
involved in each block are successively activated and
then heated to model the moving heat source. However,
since the main aim of study was to obtain the tempera-
ture gradient, distributions and thermal stress formation
for whole body of the welded T-beam, a cooling pro-
cedure was applied to the welding zone from 1450 °C to
40 °C as soon as the welding process for the steel plates
was finished. In this way the temperature gradient-dis-
tribution and the thermal stress-deformation amount,
owing to cooling to room temperature, and the thermal
expansion can be achieved using the finite-element mo-
del with the transient temperature-displacement analysis.
According to this cooling procedure, the temperature
contours of the weldments are shown in Figures 5a and
5b. It indicated that the grey region was the welding pla-
ce and its temperature was 1450 °C. When the cooling
step of the symmetrical T-beam started and the tempe-
rature reduced from 1450 °C to 40 °C in 377 increments,
the temperature distributions and thermal gradient took
place from the cross-section of the welding region to flat
bars, as shown in Figures 6a and 6b. According to the
nodal temperature variation versus time results, as
graphically represented in Figure 6a, for example, at the
node 348 (near the bottom plate), the reduction of the
temperature from 1450 °C to 40 °C is faster than for the
other nodes because of the cold bottom plate. So as the
location of the temperature measurement changed from
near the bottom plate to the centre of the welding region,
the reduction of the temperature rate decreased and the
cooling effect on the thermal gradient was small. On the
other hand, the temperature distribution of the bottom
steel plate is shown in Figure 6b. In accordance with the
results, the temperature is increased from 40 °C to 88 °C
at the node 145 (near the welding region) due to the heat
diffusion from the welding region to the bottom steel
plate over time. In addition, the temperature increase of
O. CULHA: FINITE ELEMENT MODELLING OF SUBMERGED ARC WELDING PROCESS FOR A SYMMETRIC T-BEAM
246 Materiali in tehnologije / Materials and technology 48 (2014) 2, 243–248
Figure 5: Temperature contour of welded region for: a) 2D model and
b) representation of extruded 2D model
Slika 5: Kontura temperaturnega podro~ja varjenja za: a) 2D model in
b) ekstrudiran 2D model
Figure 4: Mesh designs of 2D symmetrical welded T-beam: a) normal
view and b) magnified view of welded region
Slika 4: Postavitev mre`e 2D simetri~no varjenega T-nosilca: a) nor-
malen pogled, b) pove~ano podro~je zvara
the surface of the bottom steel plate is slowed down as it
moves away from the source region (from node 145 to
node 137).
4.2 Mechanical analysis
The same finite-element mesh as used in the thermal
analysis was employed in the mechanical analysis. The
analysis was conducted using temperature histories com-
puted with the thermal analysis as the input information.
Similar to the thermal analysis, in the mechanical
analysis, the temperature-dependent mechanical proper-
ties,10,12,13 such as Young’s modulus, yield strength and
thermal expansion coefficient, are employed. The elastic
strain-stress relationship is modelled using the isotropic
Hooke’s law, and the plastic behaviour is considered
through the Von Misses criterion. According to the
counter analysis of the welding region as represented in
Figure 7a, the nodal stress distribution shows that the
thermal residual stress exceeded the yield stress of the
welding material and the region is plastically deformed
(stress exceeded 350 MPa). In detail, it can be seen that
the nodal stress distribution is decreased from node 145
to node 142, as approximately 280 MPa to 70 MPa with
time. When the temperature gradient effect started to
disappear at the surface of the bottom steel plate (move
away from the welding region), the thermal stress
formation began to reduce, as shown in Figure 7b.
In this study an elastic-perfectly-plastic material
model is used with Von Mises failure criteria and the
O. CULHA: FINITE ELEMENT MODELLING OF SUBMERGED ARC WELDING PROCESS FOR A SYMMETRIC T-BEAM
Materiali in tehnologije / Materials and technology 48 (2014) 2, 243–248 247
Figure 7: Von Mises stress formation and distribution stress distribu-
tion of: a) welding surface and b) bottom steel plate
Slika 7: Nastanek in razporeditev Misesovih napetosti: a) varjena po-
vr{ina, b) spodnja jeklena plo{~a
Figure 6: Nodal temperature distributions of: a) welding zone and b)
surface of bottom steel plate for coupled temperature-displacement
analysis
Slika 6: Razporeditev temperaturnih vozlov: a) podro~je zvara in b)
povr{ina spodnje plo{~e za skupno analizo temperatura – izkrivljanje
Figure 8: Nodal x-axis displacement variation of bottom steel plate
between: a) 0–40 s, b) 40–3600 s after welding, c) nodal y-axis dis-
placement variation of bottom steel plate between 0–40 s, d) 40–3600 s
after welding
Slika 8: Variiranje izkrivljanj vozlov po x-osi na spodnji plo{~i: a)
med 0–40 s, b) 40–3600 s po varjenju, c) variiranje izkrivljanj vozlov
po y-osi na spodnji plo{~i med 0–40 s, d) 40–3600 s po varjenju
associated flow rule, which states that the plastic flow is
orthogonal to the yield surface. Nonlinearities due to the
large strain and displacement are considered. Strain
hardening is not included and was also neglected in
several previous studies.14,15 Experiments by Karlsson
and Josefson16 showed a nearly ideal plastic behaviour
for the material at temperatures above 800 °C. Since
most plastic strains during welding occur at high tempe-
rature, this indicates that a perfectly plastic behaviour is
suitable.10 So, the bottom plate of the welded material
model has a displacement due to the thermal gradient,
the stress and the strain distribution with time. Accord-
ing to the results of the finite-element simulation of the
model, the x-axis nodal displacement (node number
system: 10, 131, 132, 133, 134, 135 and 136) of the
bottom steel plate is illustrate in Figure 8a: 40 s after
welding and Figure 8b: 40–3600 s after welding. Fig-
ures 8a and 8b explain that when the cooling procedure
began at the welding region, the nodal displacement of
the bottom-edge steel plate was firstly increased
(positive direction) from starting point to 0.034 mm until
13 s, then decreased to 0.015 mm at the end of the cool-
ing step. On the other hand, the y-axis net nodal displa-
cement of the bottom region after 2 s, between 2–10 s
and 10–3600 s was decreased to the –0.11 mm (negative
direction), increased to the 0.0 point and increased form
0.0 point to the +0.081 mm, respectively. So a detailed
analysis showed that the thermal expansion and the stress
caused a deformation at the welding region and joint
steel plate as shown in Figure 8. The deformed and un-
deformed bottom steel plate are represented in contours
in Figure 9. It is clear that the bottom steel plate had a
deformation depending on the thermal expansion of the
welding region and the heat fusion from the source to the
steel joints using a superimposed plot option of the
analysis.
5 CONCLUSION
A finite-element analysis of the Submerged Arc
Welding (SAW) process was applied to the T-beam joint
of AH 36 quality steel plates to obtain the temperature
history, thermal gradient, stress distribution and nodal
displacement of the 2D model. According to the analy-
sis, the following time-dependent results were obtained:
The temperature gradient and the thermal history of
model were decreased from the welding region to the
bottom steel plate, as represented graphically and with
contours. The temperature decreased from 1450 °C to
40 °C at the welding zone and increased from 40 °C to
86 °C at the surface of the bottom steel plate.
In this research a 2D model was constructed as 2D
and extruded to 3D. According to the analysis results the
formation of the thermal stress at the welding zone
exceeded the yield strength of material, and at 350 MPa
plastic deformation occurred. When the cooling proce-
dure of the welding finished, the graphical and contour
representation showed that the bottom steel plate of the
joints had a shrinkage effect due to the thermal stress. In
addition, the thermal stress decreased from 350 MPa to
approximately 80 MPa away from the welding zone.
The nodal displacement of the model revealed that
the x and y axes net displacements of the bottom steel
plate were 0.015 mm and 0.081 mm, respectively.
Acknowledgement
The author wishes to thank The Scientific and Tech-
nological Research Council of Turkey (TUBITAK-1501)
for the financial support of the project titled and num-
bered as: Design and developed of thermal straightening
machine using induction heating for welding application
of steel joints and 3110745, respectively. The author
would also like to thank the project team of Ozkan Iron
and Steel Industry Co. Ltd.; Hakan Erçay and Serhat
Turgut for their technical facilities.
6 REFERENCES
1 T. L. Teng, P. H. Chang, W. C. Tseng, Computers and Structures, 81
(2003), 273
2 B. Petrovski, M. Koçak, Mis-Matching of Welds, ESIS 17 (Edited by
K. H. Schwalbe, M. Koçak), Mechanical Engineering Publications,
London 1994, 511
3 F. Jonassen, J. L. Meriam, E. P. Degarmo, Weld J, 25 (1946) 9, 492
4 E. F. Rybicki, P. A. Mcguire, Trans ASME J. Eng Mater Technol.,
104 (1982), 267
5 R. L. Koch, E. F. Rybicki, R. D. Strttan, J. Eng Mater Technol., 107
(1985), 148–53
6 B. L. Josefson, Trans ASME J. Press Vessel Technol., 104 (1982),
245
7 C. Heinze, C. Schwenk, M. Rethmeier, Journal of Constructional
Steel Research, 19 (2011), 1847
8 P. J. Withers, H. K. D. H. Bhadeshia, Mater. Sci. Technol., 17 (2001),
366
9 T. Kannengießer, T. Böllinghaus, M. Neuhaus, Weld World, 50
(2006), 11
10 L. Gannon, Y. Liu, N. Pegg, M. Smith, Marine Structures, 23 (2010),
385
11 S. W. Wen, P. Hilton, D. C. J. Farrugia, Journal of Material Process-
ing Technology, 119 (2001), 203
12 P. Michaleries, A. Debiccari, Welding Journal, 76 (1997), 172
13 K. Ogawa, D. Deng, S. Kiyoshima, N. Yanagida, K. Saito, Computa-
tional Materials Science, 45 (2009), 1031
14 L. Tall, Welding Journal, 43 (1964), 10
15 Y. Ueda, T. Yamakawa, Proceedings of international conference on
mechanical behaviour of materials, 1971, 10
16 R. Karlsson, B. Josefson, ASME Journal of Pressure Vessel
Technology, 112 (1990), 84
248 Materiali in tehnologije / Materials and technology 48 (2014) 2, 243–248
O. CULHA: FINITE ELEMENT MODELLING OF SUBMERGED ARC WELDING PROCESS FOR A SYMMETRIC T-BEAM
Figure 9: Deformed and undeformed model representation of welding
using superimposed plot option of finite-element simulation
Slika 9: Predstavitev modela, deformiranega in nedeformiranega, pri
varjenju s predpostavko mo`nosti simulacije z metodo kon~nih
elementov
A. ALTIN: OPTIMIZATION OF THE TURNING PARAMETERS FOR THE CUTTING FORCES ...
OPTIMIZATION OF THE TURNING PARAMETERS FOR
THE CUTTING FORCES IN THE HASTELLOY X
SUPERALLOY BASED ON THE TAGUCHI METHOD
OPTIMIRANJE SIL ODREZAVANJA S TAGUCHIJEVO METODO
PRI STRU@ENJU SUPERZLITINE HASTELLOY X
Abdullah Altin
Van Vocational School of Higher Education, Yuzuncu Yýl University, 65080 Van, Turkey
aaltin@yyu.edu.tr
Prejem rokopisa – received: 2013-04-23; sprejem za objavo – accepted for publication: 2013-06-11
In this study the effects of the cutting-tool coating material and cutting speed on the cutting forces and surface roughness were
determined using the Taguchi experimental design. For this purpose, the nickel-based superalloy Hastelloy X was machined
under dry cutting conditions with three different cemented-carbide tools. The main cutting force, Fz, is considered to be the
cutting force as a criterion. The mechanical loading and the abrasiveness of the carbide particles have an increasing effect on the
cutting forces. According to the results of the analysis of variance (ANOVA), the effect of the cutting speeds was not important.
Depending on the tool-coating material, the lowest main cutting force was found to be 538 N and the lowest average surface
roughness, 0.755 μm, both at 100 m/min with a multicoated cemented-carbide insert KC9240, whose top layer is coated by TiN.
Moreover, the experimental results indicated that the CVD cutting tools performed better than the PVD and the uncoated cutting
tools when turning the Hastelloy X in terms of the surface quality and the cutting forces with the current parameters.
Keywords: machinability, Hastelloy X, superalloy, cutting force, surface quality
V tej {tudiji so s Taguchijevim eksperimentalnim sestavom dolo~eni vplivi materiala opla{~enega rezilnega orodja, hitrosti
rezanja na silo odrezavanja in hrapavost povr{ine. V ta namen je bila superzlitina na osnovi niklja Hastelloy X obdelana pri
suhih razmerah s tremi razli~nimi orodji iz karbidne trdine. Kot merilo je bila vzeta glavna sila rezanja Fz. Mehansko
obremenjevanje in abrazivnost karbidnih delcev vplivata na pove~evanje sile pri odrezavanju. Skladno z rezultati analize
variance (ANOVA) vpliv hitrosti odrezavanja ni bil pomemben. Odvisno od prevleke na rezalnem orodju je bila ugotovljena
najmanj{a stri`na sila 538 N in najmanj{a povpre~na hrapavost (0,755 μm), oboje pri 100 m/min z ve~plastnim nanosom na
vlo`ku iz karbidne trdine KC9240 z vrhnjim prekritjem iz TiN. Rezultati poskusov so pokazali, da je pri stru`enju Hastelloy X s
stali{~a kvalitete povr{ine, sil pri rezanju pri danih parametrih CVD rezilno orodje bolj{e od PVD in od orodja brez prevlek.
Klju~ne besede: obdelovalnost, Hastelloy X, superzlitina, sile rezanja, kvaliteta povr{ine
1 INTRODUCTION
Nickel-based alloys constitute an important class of
materials that are used under demanding conditions of
high corrosion resistance and high-temperature strength.
These characteristics together with their good ductility
and ease of cold working make them generally very
attractive for a wide variety of applications; nearly all of
which exploit their corrosion resistance in atmospheric,
salt water and various acidic and alkaline media.1 Hastel-
loy X is a nickel-chromium-iron-molybdenum alloy that
has been developed for high-temperature applications
and is derived from the strengthening particles, Ni2(Mo,
Cr), which are formed after a two-step age-hardening
heat-treatment process. With their face-centred cubic
(FCC) structure, the Ni-Cr-Mo-W alloys, known as
Hastelloys, are used for marine engineering, chemical
and hydrocarbon processing equipment, valves, pumps,
sensors and heat exchangers.2 Hastelloy X is chosen by
many for use in furnace applications because it has an
unusual resistance to oxidizing, reducing, and neutral
atmospheres. The resistance to localized corrosion ma-
kes the alloy an attractive material as a general-purpose
filler metal, or weld overlay.3 Hastelloy X is widely used
in the clamshell of a rocket, engine tailpipes, afterburner
components, cabin heaters, and other aircraft parts.4 It
has also been found to be resistant in petrochemical
applications. Hastelloy X is widely used in a number of
industries.5 Wang,6 Richards and Aspinwall,7 Ezugwu
and Wang,8 with Khamsehzadeh9 studied the effect of
applied stress and temperatures generated at the cutting
edge and they were found to influence the wear rate and,
hence, the tool life. Notching at the tool nose and the
depth of cutting region was a prominent failure mode
when machining nickel-based alloys. This is due to a
combination of high temperature, high workpiece
strength, work hardening and abrasive chips. Kramer and
Hartung,10,11 observed that cemented-carbide tools used
for machining nickel-based alloys at a cutting speed of
30 m/min failed due to the thermal softening of the
cobalt binder phase and the subsequent plastic defor-
mation of the cutting edge. Focke et al.12 examined worn
tools, which revealed a layer of "disturbed material"
beneath the crater and the cutting edge. Hastelloy is a
registered trademark name of Haynes International Inc.
The Hastelloy trademark is applied as a prefix for a
range of corrosion-resistant nickel-based alloys
Materiali in tehnologije / Materials and technology 48 (2014) 2, 249–254 249
UDK 669.24:621.941:519.233.4 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)249(2014)
promoted under the name "superalloys" or "high-
performance alloys". Within corrosion applications,
Hastelloy alloys may be chosen as a trade-off between
performance, cost and other technical issues, e.g.,
suitability for welding. Hastelloy alloys are generally
less attractive for use in acids compared to Tantaline
(Tantaline is recognized as the leading performance/price
option).
Nomenclature
v cutting speed (m/min)
f feed (mm/r)
da axial depth (mm)
y tool life (min)
Fm feed (mm/min)
TL total length
2 MATERIALS AND METHOD
2.1 Experiment Specimens
Specimens of Hastelloy X, which has an industrial
usage, were prepared with dimensions of diameter Ø2 ×
40 inches and then used for the experiments. The che-
mical composition and mechanical properties of the
specimens are given in Tables 1 and 2, respectively. As
the contents of the workpiece, chromium (21 %) and
molybdenum (17 %), are high, the material is hard to
machine. The material consists of approximately 50 %
nickel, making the alloy suitable for high-temperature
applications. The specimen was annealed and has a
hardness of Rockwell B 90.
2.2 Machine Tool and Measuring Instrument of Cut-
ting Forces
The machining tests were carried out on a JOHN-
FORD T35 industrial-type CNC lathe with a maximum
power of 10 kW and a rotating speed between 50 r/min
and 3500 r/min. During the dry cutting process, a Kistler
brand 9257 B-type three-component piezoelectric dyna-
mometer under a tool holder with an appropriate load
amplifier was used for measuring three orthogonal cut-
ting forces (Fx, Fy, Fz). This allows direct and continuous
recording and a simultaneous graphical visualization of
the three cutting forces. The technical properties of the
dynamometer and a schematic figure of the experimental
setup are given in Table 3 and Figure 1, respectively.
Table 1: Chemical composition of the workpiece material (Hastelloy
X), w/%
Tabela 1: Kemijska sestava obdelovanca (Hastelloy X), w/%
Ni Cr Mo Fe Co W Mn Al Si C B
50 21 17 2 1 1 0.80 0.50 0.08 0.01 0.01
Table 2: Mechanical properties of Hastelloy X
Tabela 2: Mehanske lastnosti Hastelloy X
Hardness
conductivit
y(HB)
Tensile
strength
(MPa)
Yield
strength
(MPa)
Breaking
extension %
(5 do)
Thermal
(W/m K)
388 1370 1170 23.3 11.4
Table 3: Technical properties of dynamometer
Tabela 3: Tehni~ne lastnosti dinamometra
Force interval (Fx, Fy, Fz) -5…10 kN
Reaction < 0.01 N
Accuracy Fx, Fy ≈ 7.5 pC/N
Accuracy Fz ≈ 3.5 pC/N
Natural frequency f0(x, y, z) 3.5 kHz
Working temperature 0…70 °C
Capacitance 220 pF
Insulation resistance at 20 °C > 1013 
Grounding insulation > 108 
Mass 7.3 kg
2.3 Cutting Parameters, Cutting Tool and Tool Holder
The cutting speeds (50, 65, 80, and 100) m/min were
chosen by taking into consideration the ISO 3685
standard, as recommended by the manufacturers. The
depth of the cut (1.5 mm) and the constant feed rate
(0.10–0.15 mm/r) were chosen to be constant. During the
cutting process, the machining tests were conducted with
three different cemented-carbide tools, i.e., Physical
Vapour Deposition (PVD) coated with TiN/TiCN/TiN;
Chemical Vapour Deposition (CVD) coated with
TiN+AL2O3-TiCN+TiN; and WC/CO. The dimensions of
the test specimens were 2 × 40 inches in terms of
diameter and length. The properties of the cutting tools
and the level of the independent variables are given in
Tables 4 and 5. Surtrasonic 3-P measuring equipment
was used for the measurement of the surface roughness.
The measurement processes were carried out with three
replications. For measuring the surface roughness on the
workpiece during machining, the cut-off and sampling
length were considered as 0.8 mm and 2.5 mm, respec-
tively. The ambient temperature was (20 ± 1) °C. The
resultant cutting force was calculated to evaluate the
machining performance. The following are the details of
the tool geometry CNMG inserts when mounted on the
tool holder: (a) CNMG shape; (b) axial rake angle, 6°;
A. ALTIN: OPTIMIZATION OF THE TURNING PARAMETERS FOR THE CUTTING FORCES ...
250 Materiali in tehnologije / Materials and technology 48 (2014) 2, 249–254
Figure 1: Measurement of the cutting forces and a schematic figure of
the dynamometer unit
Slika 1: Merjenje sil pri rezanju in shematski prikaz dinamometrske
enote
(c) end relief angle, 5°; and (d) sharp cutting edge. The
cutting tool was mounted in the tool holder (PCLNR
2525M12) and used for such cutting tools (CNMG
120404) with an approach angle of 75°. Analysis of
variance (ANOVA) was applied to the experimental
study.
Table 5: Level of independent variables
Tabela 5: Nivoji neodvisnih spremenljivk
Variables
Level of variables
Lower Low Medium High
Cutting speed,
v/(m/min) 50 65 80 100
Feed, f/(mm/r) 0.1–0.15 0.1–0.15 0.1–0.15 0.1–0.15
Axial depth, da/mm 1.5 1.5 1.5 1.5
3 RESULTS AND DISCUSSION
3.1 Cutting forces and surface roughness
After the test specimens were prepared for experi-
mental purposes, they were measured with a three-com-
ponent piezoelectric dynamometer to obtain the main
cutting force. According to Figure 2, increasing the
cutting speed decreases the main cutting force, excluding
the area between 80 m/min and 100 m/min for K313.
The lowest obtained values for the main cutting force at
the cutting speeds of 50 m/min, 662 N, 65 m/min, 622 N,
80 m/min, 601 N, and 100 m/min, 538 N at a 0.1 mm/r
constant feed rate, respectively. The lowest main cutting
force was observed at 100 m/min cutting speed as 538 N.
In Figure 2 the main cutting force depending on the
cutting speed and the uncoating material of the cutting
tool were changed in all the experiments. The cutting
forces and surface roughness according to the experi-
mental cutting parameters are given in Table 6. Accord-
ing to W. Konig, the cutting speed must be increased in
order to reduce the main cutting forces.13 However, in
this study, a decrease was observed in the main cutting
force between 50 m/min and 100 m/min. It is thought
that this case is caused by the good performance of the
cutting tool. The effect of the TiN-coated carbide inserts
was found to be important for the main cutting force, but
the effect of the cutting speeds was not important in the
analysis of variance. The main cutting force decreases in
spite of increasing the cutting speed from 50 m/min to
100 m/min. As a result of the experimental data, an
increase of 100 % in the cutting speed (from 50 m/min to
100 m/min), a decrease in the main cutting force with
K313 (6.5 %), KT315 (10 %) and KC9240 (19 %) was
found with the 0.1 mm/r constant feed rate, and the
decreasing contact surface area caused the main cutting
force to decrease in comparison to the increased cutting
speed. The decrement of the cutting force depends on the
material type, the working conditions and the cutting
speed range.14 The high temperature for the flow region
and the decreasing contact area and chip thickness cause
the cutting force to decrease, depending on the cutting
speed.15
As is widely known, the cutting speed must be
decreased to improve the average surface roughness.16
The scatter plot between the surface roughness and the
cutting speed as shown in Figure 3 indicated that there
was a linear relationship between the surface roughness
and the cutting speed. The results of Figure 3 show that
the average surface roughness decreases by 66 % with an
increasing cutting speed from 50 m/min to 100 m/min
A. ALTIN: OPTIMIZATION OF THE TURNING PARAMETERS FOR THE CUTTING FORCES ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 249–254 251
Table 4: Properties of the cutting tools
Tabela 4: Lastnosti rezilnih orodij
Coating material
(top layer) Coating method and layers
ISO grade of material
(grade) Geometric form
Manufacturer and
code
TiN CVD (TiN, Al2O3, TiCN, TiN, WC) P25-40, M20-30 CNMG120404RP Kennametal KC9240
TiN PVD (TiN, TiCN, TiN, WC) P25-40, M20-30 CNMG120404FN Kennametal KT315
WC-CO Uncoated P25-40, M20-30 CNMG120404MS Kennametal K313
Figure 2: Change of main cutting force in Hastelloy X, according to
the cutting speed: a) at f = 0.1 mm/r, b) at f = 0.15 mm/r
Slika 2: Spreminjanje glavne sile rezanja materiala Hastelloy X pri
hitrostih rezanja: a) f = 0,1 mm/r in b) f = 0,15 mm/r
with the KC9240 cutting tool (0.1 mm/r constant feed
rate).
3.2 Optimization with the Taguchi Method
In this part the optimization of the turning parameters
was carried out in terms of the cutting forces with the
Taguchi analysis. The importance order of the effects of
each control factor on the turning forces was identified.
For this purpose, the factors selected in the Taguchi
experimental design and the levels of these factors are
shown in Table 7. Taguchi’s L18 2*1 3*2 mixed design
was used. In the Taguchi method there are three catego-
ries, i.e., "the smallest is better", "the biggest is better"
and "the nominal is better" for the calculation of the
signal/noise (S/N) ratio. In this research "smallest is
better" was used since the minimum of the cutting force
and surface roughness was intended. In the ith experi-
ment, the S/N ratio can be calculated using the following
equation.17–19
 i i
n
n
Y= − ∑10 110 2
1
log (1)
n is the number of replications and Yi is the measured
characteristic.
Table 7: Cutting parameters and levels
Tabela 7: Parametri rezanja in nivoji
Control parameters Units
Levels
1 2 3
Feed rate (A) m/min 65 80 100
Cutting speed (B) (mm/r) 0.1 0.15
Cutting tool (C) KC313 KT315 KC9240
3.3 Confirmation Experiments
The final step of the Taguchi experimental design
process includes confirmation experiments.18,19 To
achieve this, the results of the experiments were compa-
red with the predicted values using the Taguchi method
and the error rates were obtained. The S/N ratios were
predicted using the following model:18–20
   pred = + −
=
∑m i m
i
k
( )
1
(2)
Moreover, the optimum turning parameters were
obtained for the performance characteristics using the
Taguchi analysis, where m is the total mean of the S/N
ratios, i is the mean S/N ratio at the optimum level and k
is the number of the main design parameters that signifi-
cantly affect the performance characteristics. After pre-
dicting the S/N ratios other than 18 experiments (with
Eq.2), the main cutting force or Fz was calculated using
the following equation:20
Y
S N
pred =
−
10 20
/
(3)
A. ALTIN: OPTIMIZATION OF THE TURNING PARAMETERS FOR THE CUTTING FORCES ...
252 Materiali in tehnologije / Materials and technology 48 (2014) 2, 249–254
Figure 3: According to cutting speed the change of the average
surface roughness in Hastelloy X: a) at f = 0.1 mm/r, b) at f = 0.15
mm/r
Slika 3: Spreminjanje povpre~ne hrapavosti Hastelloya X od hitrosti
rezanja: a) pri f = 0,1 mm/r in b) pri f = 0,15 mm/r
Table 6: Cutting forces and surface roughness according to the
experimental cutting parameters when turning Hastelloy X
Tabela 6: Sile rezanja in hrapavost povr{ine pri ustreznih parametrih
poskusa rezanja pri stru`enju Hastelloy X
E
xp
er
im
en
t
nu
m
be
r
C
ut
ti
ng
to
ol
C
ut
ti
ng
sp
ee
d
(m
/m
in
)
F
ee
d
ra
te
(m
m
/r
)
D
ep
th
of
cu
t
(m
m
)
Fx/N Fy /N Fz /N Ra/μm
1 K313 65 0.1 1.5 366 75 691 1.7
2 K313 80 0.1 1.5 323 67 655 1.599
3 K313 100 0.1 1.5 316 72 658 1.717
4 KT315 65 0.1 1.5 295 132 622 1.605
5 KT315 80 0.1 1.5 281 130 601 1.41
6 KT315 100 0.1 1.5 277 130 598 1.667
7 KC9240 65 0.1 1.5 446 203 715 1.455
8 KC9240 80 0.1 1.5 441 184 694 1.368
9 KC9240 100 0.1 1.5 393 158 538 0.755
10 K313 65 0.15 1.5 398 96 919 3.649
11 K313 80 0.15 1.5 371 90 901 3.462
12 K313 100 0.15 1.5 325 78 854 3.137
13 KT315 65 0.15 1.5 358 177 863 2.669
14 KT315 80 0.15 1.5 356 179 855 1.88
15 KT315 100 0.15 1.5 335 177 830 3.132
16 KC9240 65 0.15 1.5 527 272 966 1.492
17 KC9240 80 0.15 1.5 446 187 696 1.405
18 KC9240 100 0.15 1.5 436 195 697 1.085
where Ypred is the main cutting force or Fz with regard to
the S/N ratio.
3.4 Taguchi Analysis for FZ
Figure 4 shows the main effect plot for the S/N
ratios, giving the effect of cutting parameters on the
cutting force Fz. The smallest main cutting force is ob-
tained with the cutting insert KT315. In this way, to
achieve the minimum cutting forces it is understood that
the KC9240 cutting tool should be used at a 0.10 mm/r
feed rate and 100 m/min cutting speed.
After analysing the effect of the cutting parameters
on the cutting force, in order to find out which Fz cutting
force it effected, a variance analysis was made. Accord-
ing to the results of the ANOVA in Table 8, it is under-
stood that the most effective cutting parameters that
effected the cutting force was 65.99 % and the cutting
speed was 11.14 %.
Table 8: ANOVA results for Fz
Tabela 8: ANOVA-rezultati za Fz
Source DF Sum ofsquares
Mean
square F
Prob >
F
Distribu-
tion %
A 1 181805 181805 57.75 0.002 65.99
B 2 30700 15350 4.88 0.085 11.14
C 2 13213 6607 2.1 0.238 4.80
A*B 2 4024 2012 0.64 0.574 1.46
B*C 2 9391 4696 1.49 0.328 3.41
A*C 4 23792 5948 1.89 0.276 8.64
Error 4 12592 3148 4.57
Total 17 275517 100.00
3.5 Taguchi Analysis for Ra
Figure 5 shows the effect of feed rate, cutting speed
and cutting-tool material on the average surface rough-
ness. According to this figure, in order to obtain the
smallest surface roughness, it is necessary to use the
KC9240 cutting tool at a low feed rate (0.10 mm/r) and a
high cutting speed (100 m/min). Besides this, in order to
find out which important parameter affects the surface
the roughness, the variance analysis was made with this
aim. According to the results of the ANOVA in Table 9,
the cutting parameters which effect the surface rough-
ness, the cutting tool (40.38 %), feed rate, (33.15%) and
15.57 % feed speed and cutting tool’s interaction were
found.
Table 9: ANOVA results for Ra
Tabela 9: ANOVA rezultati za Ra
Source DF Sum ofsquares
Mean
square F Prob > F
Distri-
bution %
A 1 4.1424 4.1424 56.56 0.002 33.15
B 2 0.18817 0.09408 1.28 0.371 1.51
C 2 5.04646 2.52323 34.45 0.003 40.38
A* 2 0.06687 0.03343 0.46 0.663 0.54
B*C 2 0.81483 0.20371 2.78 0.173 6.52
A*C 4 1.94611 0.97305 13.29 0.017 15.57
Error 4 0.29294 0.07323 2.34
Total 17 12.49777 100.00
4 CONCLUSIONS
The goal of this study was to identify the effect of
turning parameters such as feed rate, cutting speed and
cutting tools on the main cutting force and surface
roughness using an analysis of Taguchi. The experimen-
tal design described here was used to develop the main
cutting force and the surface-roughness prediction model
for the Hastelloy X turning operation. The results of this
experimental study can be summarized as follows:
• The main cutting force decreased with increasing
cutting speed and the cutting force increased at
higher feed rates.
• According to the analysis of Taguchi, in order to
obtain the smallest cutting force and surface rough-
ness, it was necessary to use the KC9240 cutting tool
A. ALTIN: OPTIMIZATION OF THE TURNING PARAMETERS FOR THE CUTTING FORCES ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 249–254 253
Figure 5: Mean response graphs of the surface roughness according to
feed rate, cutting speed and cutting tool
Slika 5: Prikaz odziva hrapavosti povr{ine glede na hitrost podajanja,
hitrost rezanja in rezilnega orodja
Figure 4: Mean response graphs of the cutting forces according to the
feed rate, cutting speed and cutting tool
Slika 4: Prikaz odziva sil rezanja glede na hitrost podajanja, hitrost
rezanja in rezilnega orodja
at a low feed rate (0.10 mm/r) and a high cutting
speed (100 m/min).
• The minimum main cutting force is obtained with
CNMG 120404-type multicoated TiN+AL2O3-TiCN+
TiN carbide tools, while the maximum main cutting
force is obtained as 965 N with the CNMG 120404-
type uncoated carbide tools. However, cemented
carbide tools have no significant effect on the main
cutting force when machining Hastelloy X.
• An increasing relation between the cutting speed and
the arithmetic average surface roughness as well as
between the coating number and the average surface
roughness is observed.
• In the case of coated tools, the effect of cutting speed
on the surface roughness is no more pronounced than
the effect of uncoated cemented-carbide inserts.
• The minimum average surface roughness is deter-
mined with CNMG 120404-type multicoated TiN+
AL2O3-TiCN+TiN carbide tools, while the maximum
average surface roughness is observed with CNMG
120404-type uncoated tools. Moreover, uncoated and
multiple-layer coated tools have a significantly
different effect on the average surface roughness.
• It was found that there is a positive correlation bet-
ween the main cutting force and the average surface
roughness.
Acknowledgments
The authors would like to express their gratitude to
the University of Yuzuncu Yýl for the financial support
Under Project No. BAP 2012-BYO-013.
5 REFERENCES
1 Q. Zhang, R. Tang, K. Yin, X. Luo, L. Zhang, Corrosion behavior of
Hastelloy C-276 in supercritical water, CSci., 51 (2009), 2092–2097
2 V. B. Singh, A. Gupta, The electrochemical corrosion and passiva-
tion behavior of Monel 400 in concentrated acids and their mixtures,
Transaction of JWRI., 34 (2000), 19–23
3 Haynes Hastelloy C-22HS Standard Product Catalogue, Haynes
International, Indiana, 2007, 2–16
4 P. C. Jindal, A. T. Santhanam, U. Schleinkofer, A. F. Shuster, Per-
formance of PVD TiN, TiCN, and TiAlN coated cemented carbide
tools in turning, Int. J. Recfrac. Met. Hard Mater., 17 (1999),
163–170
5 Website of trademark owner of Hastelloy C-276. www.hynesintl.com
6 M. Wang, Ph. D. Thesis, South Bank University, London, 1997
7 N. Richards, D. D. Aspinwall, Use of ceramic tools for machining
nickel-based alloys, Int. J. Mach. Tools Manuf., 29 (1989) 4,
575–588
8 E. O. Ezugwu, Z. M. Wang, Performance of PVD and CVD coated
tools when machining nickel-based, Inconel 718 alloy, In: N. Naru-
taki, et al. (Eds.), Third International Conference on Progress of
Cutting and Grinding, 111 (1996), 102–107
9 H. Khamsehzadeh, Behaviour of ceramic cutting tools when machin-
ing superalloys, Ph.D. Thesis, Universtiy of Warwick, 1991
10 J. Barry, G. Byrne, Cutting tool wear in the machining of hardened
steels. Part I. Cubic boron nitride cutting tool wear, Wear, 247
(2001), 139–151
11 B. M. Kramer, P. D. Hartung, Proc. Int. Conf. of Cutting Tool Mat.,
Fort Mitchell, KY, (1980), 57–74
12 A. E. Focke, F. E. Westermann, A. Ermi, J. Yavelak, M. Hoch, Fail-
ure mechanisms of superhard materials when cutting superalloys, in:
Proc. 4th Int.-Am. Conf. of Mat. and Tech, Caracus, Venezuela,
1975, 488–497
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ponents not only by grinding, Ind. Diamond Rev., 3 (1994), 127–132
14 C. Çakýr, Modern metal cutting principles, Vipaº, Bursa 2000
15 Sandvik, Modern metal cutting practical Handbook, Sandvik 1994
16 R. I. King (Ed.), Handbook of high speed mach., techn. Chapman
and Hall, London 1985
17 A. Taskesen, K. Kütükde, Optimization of the drilling parameters for
the cutting forces in B4C-reinforced Al-7XXX-series alloys based on
the Taguchi method, Mater. Tehnol., 47 (2013) 2, 169–176
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AL/SIC P metal matrix composite, International Journal of
Advanced Manufacturing Technology, 55 (2011) 5–8, 477–485
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Nostrand Reinhold, New York, 1990
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ling of GFRP composites, Measurement, Journal of the International
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A. ALTIN: OPTIMIZATION OF THE TURNING PARAMETERS FOR THE CUTTING FORCES ...
254 Materiali in tehnologije / Materials and technology 48 (2014) 2, 249–254
L. B. GETSOV et al.: THERMOCYCLIC- AND STATIC-FAILURE CRITERIA FOR SINGLE-CRYSTAL SUPERALLOYS ...
THERMOCYCLIC- AND STATIC-FAILURE CRITERIA
FOR SINGLE-CRYSTAL SUPERALLOYS OF
GAS-TURBINE BLADES
TERMOCIKLI^NA IN STATI^NA MERILA ZA PORU[ITVE
LOPATIC PLINSKIH TURBIN IZ MONOKRISTALNIH
SUPERZLITIN
Leonid B. Getsov1, Artem S. Semenov2, Elena A. Tikhomirova3, Alexander I. Rybnikov1
1NPO ZKTI, Russia
2St. Petersburg State Polytechnical University, Russia
3Klimov Company, Russia
guetsov@yahoo.com, semenov.artem@googlemail.com
Prejem rokopisa – received: 2013-05-17; sprejem za objavo – accepted for publication: 2013-06-28
The problem of the thermo-mechanical fatigue (TMF) of single-crystal turbine blades has not been fully investigated
theoretically or experimentally. In the present work TMF tests were performed for two single-crystal nickel-based alloys (ZhS36
and ZhS32) with various crystallographic orientations (001, 011, 111) under different temperatures and cycle durations.
The dependence of the failure modes (crystallographic or non-crystallographic) on the loading regimes was analyzed. The
non-linear viscoelastic, elastoplastic and viscoelastoplastic material models with non-linear kinematic hardening were used to
predict the cyclic stress-strain state, ratcheting and creep of the samples. The deformation criterion of damage accumulation was
introduced to the lifetime prediction. A stress analysis of the single-crystal samples, with concentrators (in the form of circular
holes) and without them, was carried out using physical models of the plasticity and creep. These material models take into
account that an inelastic deformation occurs due to a slip mechanism and it is determined with the crystallographic orientation.
The proposed failure model using the deformation criterion allows qualitative and quantitative predictions of the TMF fracture
process in single crystals.
Keywords: gas-turbine blade, single crystal, thermo-mechanical fatigue, damage, crystallographic and non-crystallographic
failure modes
Problem termo-mehanske utrujenosti (TMF) monokristalnih turbinskih lopatic {e ni v celoti raziskan niti teoreti~no niti
eksperimentalno. V tem delu so bili izvr{eni TMF-preizkusi na dveh monokristalnih zlitinah na osnovi niklja (ZhS36 in ZhS32)
z razli~no kristalografsko orientacijo (001, 011, 111) pri razli~nih temperaturah in trajanju ciklov. Analizirana je bila
odvisnost na~ina poru{itve (kristalografska ali nekristalografska) od vrste obremenjevanja. Modeli nelinearne viskoelasti~nosti,
elastoplasti~nosti in viskoelastoplasti~nosti z nelinearnim kinemati~nim utrjevanjem so bili uporabljeni za napovedovanje
cikli~nega stanja napetost – raztezek, nazob~anja in lezenja vzorcev. V napovedovanje dobe trajanja je bilo vpeljano
deformacijsko merilo akumuliranja po{kodb. Izvr{ena je bila analiza napetosti v monokristalnem vzorcu s koncentratorji (v
obliki okroglih lukenj) in brez njih, z uporabo fizikalnega modela plasti~nosti in lezenja. Ti materialni modeli upo{tevajo, da se
pojavi neelasti~na deformacija z mehanizmom lezenja in je dolo~ena s kristalografsko orientacijo. Predlagani model poru{itve z
uporabo deformacijskih meril omogo~a kvalitativno in kvantitativno napovedovanje TMF- procesa preloma monokristala.
Klju~ne besede: lopatica plinske turbine, monokristal, termo-mehanska utrujenost, po{kodba, kristalografski in nekristalografski
na~in poru{itve
1 INTRODUCTION
The use of single-crystal alloys for the manufacturing
of the blades of a gas-turbine engine allows a significant
increase in the gas temperature before a turbine and sets
a number of tasks that should help increase the reliability
of a stress and strength analysis. In the present investi-
gation, the results of the experimental studies of single-
crystal nickel-based alloys, as well as the approaches to
the computation of the stress-strain state and strength of
the structural parts are considered and discussed.
2 MATERIALS AND METHODS OF RESEARCH
Numerous experimental studies were performed on
two single-crystal alloys, ZhS32 and ZhS36 (Table 1)
with different alloying and, most importantly, different
carbon amounts and were designed to expand the infor-
mation given in1,2. The tests of the mechanical properties,
creep and thermal-fatigue resistance at different tempe-
ratures were carried out.
The creep tests were conducted on an installation of
ATS (Applied Test Systems, Inc.) determining the kine-
tics of the accumulated inelastic deformation during the
first stage and in the steady state of the creep. The test
methods for TMF are described in detail in3,4. For the
tests, rigidly clamped samples with polished surfaces
were used, as shown in Figure 1. The tests were con-
ducted in a vacuum that allowed us, during the testing, to
observe the formation of slip bands and crack initiation,
and to determine the rate of the crack propagation on a
polished surface using the magnification of 250-times.
The tests were conducted with various values for the
maximum (Tmax = 900–1100 °C) and minimum (Tmin =
Materiali in tehnologije / Materials and technology 48 (2014) 2, 255–260 255
UDK 539.42:669.018 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)255(2014)
150–700 °C) cycle temperatures. The tests with the retar-
dation times of 2 min and 5 min at Tmax were carried out
for some parts of the samples. Some samples had the
stress concentrator in the form of a central hole with the
diameter of 0.5 mm. The test specimens had different
crystallographic orientations: 001, 011 or 111. To
determine the crystallographic orientation for each
sample Laue’s diffraction patterns were obtained and
three Euler angles and Schmid factors were computed.
On the basis of the results of a crystallographic ana-
lysis, possible directions (angles) of the slip bands on the
specimen surfaces were calculated with the aim to
compare them with the angles, at which the samples
were destroyed. The location of the fracture nucleus was
determined with the results of fractographic studies
using a TESCAN microscope. The comparison of the
experimental data on the orientation of the fracture
surfaces with the results of the crystallographic analysis
and with the results of the finite-element analysis of the
specimens allowed us to define the dependence of the
failure mechanism (crystallographic or non-crystallo-
graphic) on the parameters of the thermal cycle.
3 RESULTS OF EXPERIMENTAL STUDIES
The tests of mechanical properties show that single-
crystal alloys do not have a high plasticity at all the
temperatures (see, for example, Table 2). Low values of
plasticity  are observed for the carbon-free ZhS36 alloy
(as opposed to the carbon ZhS32 alloy) at 500 °C (as
opposed to the cases of high temperatures where T > 900
°C).
Table 2: Mechanical properties of single-crystal alloys with orienta-
tion 001 at 500 °C
Tabela 2: Mehanske lastnosti monokristalne zlitine z orientacijo 001
pri 500 °C
Alloy Rp0.2MPa
Rm
MPa
A
%
Z
%
ZhS36 964967
982
1000
1.3
2.3
5.0
6.9
ZhS32
Schedule t/t 1 850 880 19.5 35.5
Schedule t/t 2 810 1110 13.0 11.7
Figure 2 shows the short-term creep curves of alloy
ZhS32. The curves obtained under the stress of 550 MPa
at 850 °C significantly differ for various samples. The
results of the creep tests for alloy ZhS36 are given in5.
L. B. GETSOV et al.: THERMOCYCLIC- AND STATIC-FAILURE CRITERIA FOR SINGLE-CRYSTAL SUPERALLOYS ...
256 Materiali in tehnologije / Materials and technology 48 (2014) 2, 255–260
Table 1: Chemical compositions of the ZhS32 and ZhS36 single-crystal alloys, w/%
Tabela 1: Kemijska sestava monokristalnih zlitin ZhS32 in ZhS36, w/%
Alloy C Cr Co Mo W Ta Re Nb Al Ti Ni
ZhS32 0.12–0.18 4.3–5.6 8.0–10.0 0.8–1.4 7.7–9.5 3.5–4.5 3.5–4.5 1.4–1.8 5.6–6.3 – Base
ZhS36 0.03 4.03 8.48 1.41 11.50 - 1.94 1.07 5.70 0.91 Base
Figure 2: Creep curves of alloy ZhS32 at (850, 975 and 1050) °C
Slika 2: Krivulje lezenja zlitine ZhS32 pri (850, 975 in 1050) °C
Figure 1: a) Specimen for the thermal-fatigue test, b) with typical cyc-
lic-temperature changes over time in the central point
Slika 1: a) Vzorec za preizkus toplotnega utrujanja, b) z zna~ilnim
nihanjem temperature v srednjem delu
The tests of the ZhS36 alloy show that the conditions
for failure under the thermal cyclic loading depend on
the crystallographic orientation of the single-crystal alloy
and also on the temperatures of the cycle. Unfortunately,
the experiments conducted on the ZhS36 alloy with
orientations 001, 011 and 111 were not numerous
and the obtained results reflect only a trend. However, a
formulation of the failure criterion depends on the failure
mode (crystallographic or non-crystallographic).6 In this
research, we obtained the conditions (a range of ther-
mal-cycle parameters) (Figure 3) for the fracture modes
of the ZhS32 alloy with the orientations close to 001.
An accumulation of unilateral irreversible deformation
(ratcheting) was observed in all the tests (Figure 4) for
both alloys.
4 CRITERIA OF FAILURE UNDER STATIC
LOADING
Single-crystal superalloys, as a rule, are plastic
materials and the possibility of a brittle fracture under
static loading of gas-turbine blades is remote. However,
this issue requires a special investigation. We considered
such a possibility on the basis of two (stress and
deformation) failure criteria.
The effect of stress on deformation capacity * is
determined with the formulas of Mahutov N. A. or
Hancock J. W. and Mackenzie A. C.7:
 

* . exp
.
= ⋅
−
pr
mean
17
1 5
i (1)

 
 
* =
pr e
1 mean
K i
2
(2)
where pr is the maximum deformation, determined
from the experiments under short-term tension, and Ke
is the characteristic of the material state (in a brittle
state Ke = 1, in a viscous state Ke = 1.2).
We need to consider the effect of stress on the
fracture conditions in terms of the power criterion. Let us
consider the general case of stress: 2 = A11, 3 = A21,
where A1 and A2 can vary from – to 1. Depending on
the relations between 1, 2, 3 and on the ratio of pr/0.2,
there is an area of brittle damage, in which the ultimate
tensile stress is used as the limiting strength characteri-
stic pr for the local stresses.
The above relation can also be written in a different
form. Let us consider the case where 1/i > 1. Using an
approximation of the stress-strain curve in the form of
   i i
p
A
m
0 2. + and the fracture condition according to
the first theory of strength, 1 = pr. q = 1/i, we obtain
q A i
p m
( ). 0 2 + = pr, where the plastic deformation is
defined with the equation:
 
 
i
p
qpr m/ .−⎡
⎣⎢
⎤
⎦⎥
0 2
1
(3)
For k = pr/0.2, using:
 
 


i
p
k q k qpr m pr m/ ( / ).−⎡
⎣⎢
⎤
⎦⎥
=
−⎡
⎣⎢
⎤
⎦⎥
0 2
1 1
1
with k/q > 1, a brittle fracture takes place as k/q = 1.
An analysis of crack initiation in the blades under
static loading (centrifugal force and bending) on the
basis of the stress-failure criterion should include:
1. A thermal finite-element analysis of the steady-state
temperature field in a blade;
2. An elastic finite-element analysis of the stress fields
with the subsequent definitions of q = 1/2 and k =
pr/0.2 at the corresponding temperatures for all the
elements of cooled blades;
3. According to the first strength theory we assume that
pr = separation  F / (1 – ) and verify the absence of
equality for q = k at all the points. For the remaining
cases, we calculate the values of the plastic strain
using equation (3);
4. An estimation of the strength by comparing  i
p
(3)
with the limiting plasticity * (1) or (2).
L. B. GETSOV et al.: THERMOCYCLIC- AND STATIC-FAILURE CRITERIA FOR SINGLE-CRYSTAL SUPERALLOYS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 255–260 257
Figure 4: Ratcheting of alloys ZhS32 and ZhS36 under thermal cyclic
loading
Slika 4: Zob~enje pri zlitinah ZhS32 in ZhS36 pri toplotnem cikli-
~nem obremenjevanju
Figure 3: Map of the fracture mechanisms for alloy ZhS32 under
thermal cyclic loading
Slika 3: Prikaz podro~ij mehanizmov poru{itve za zlitino ZhS32 pri
toplotnem cikli~nem obremenjevanju
An analysis of the crack initiation in the blades under
static loading (centrifugal force and bending) on the
basis of the strain-failure criterion should include:
1. A thermal finite-element analysis of the steady-state
temperature field in a blade;
2. An elastoplastic finite-element analysis with a defini-
tion of the zones of plastic deformation and maxi-
mum values  i
p
max for all the elements of cooled bla-
des;
3. A comparison of the obtained value for  i
p
max with the
limiting characteristic * (1) or (2) at the correspond-
ing temperatures, also taking into account a decrease
under the effect of aging at the operating tempera-
tures and during long exposures;
4. A viscoelastic finite-element analysis of the stress-
relaxation processes with the initial conditions
obtained with the elastoplastic analysis (see 2);
5. A calculation of the equivalent stress. If the value of
(0 – res)/E is lower than, or approximately equal to,
the maximum ductility under creep conditions at a
suitable temperature, determined with the formulas:
p i* . exp
.
= ⋅
−



c
mean
17
1 5
(4)
p
K i
* =
 
 
c e
1 mean
2
(5)
then a brittle fracture under the conditions of stress rela-
xation is possible.
An acceptance of the assumption that fine micro-
cracks are formed in the zone of inelastic deformation.
Determination of the size of the zone of inelastic defor-
mation (plastic and creep) and a comparison with the
limit values corresponding to the beginning of the acce-
lerated crack growth in non-linear fracture mechanics.
5 CRITERIA OF FAILURE UNDER THERMAL
CYCLIC LOADING
For a prediction of a TMF failure of single-crystal
materials, it is reasonable to use the deformation crite-
rion proposed in6. The crack-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+ + + = (6)
The criterion (6) is based on the linear summation of
the damages caused by the cyclic plastic strain:
D
C i
n
i1
1 1
1
=
=
∑ ( )Δ eqp k (7)
the cyclic creep strain:
D
C i
n
i2
2 1
1
=
=
∑ ( )Δ eqc m (8)
the unilaterally accumulated plastic strain:
D3
1
=


r
p eq
pmax (9)
and the unilaterally accumulated creep strain:
D4
1
=


r
c eq
cmax (10)
C1, C2, k, m,  r
p ,  r
c are the material parameters depend-
ing on the temperature and on the crystallographic
orientation. Usually the relations k = 2, m = 5/4,
C1 = ( ) r
p k , C 2
3
4= ( ) r
c m are used.
Different norms of the strain tensor are considered as
an equivalent strain for (6); among them there are: the
maximum shear strain in the slip system with normal
n{111} to the slip plane and slip direction l011:
eq = n{111} · e · l011 (11)
the maximum tensile strain (the maximum eigenvalue of
the strain tensor):
eq = 1 = max arg {det (e – l) = 0} (12)
the von Mises strain intensity:

     
eq dev dev= ⋅ ⋅ =
= − + − + − +
2
3
2
9
2 2 2 3
2
e e
( ) ( ) ( ) (x y y z z x[ ]  xy yz zx2 2 2+ +
(13)
and the maximum shear strain:
[ ] [   

eq 1 3= − = − = −
− −
1
2
1
2 0maxarg det( )
maxarg det(
{ }
{
e
e
l
l ])=0}
(14)
Equivalent strain (11) corresponds to the crystallo-
graphic failure mode, while equivalent strains (12)–(14)
correspond to the non-crystallographic failure mode.
6 RESULTS OF THE FINITE-ELEMENT
SIMULATION
The stress-strain analysis of single-crystal samples
(Figure 1), with a concentrator (in the form of a central
circular hole) or without a concentrator, was carried out
on the basis of the finite-element program PANTO-
CRATOR8 with an application of physical models of
plasticity. These material models take into account that
an inelastic deformation occurs in accordance with the
crystal-slip systems due to a slip mechanism and, there-
fore, the deformation is strongly sensitive to the crystal-
lographic orientation. The elastoplastic and viscoelasto-
plastic material models9,10 with non-linear kinematic and
isotropic hardening, also accounting for the self-hard-
ening of each system and the latent hardening11 between
the slip systems, are used in the finite-element simu-
lations. The application of viscoelastic models leads to
unrealistically overestimated levels of the stress.
The obtained results for the inhomogeneous stress,
strain and damage fields allow us to find the location of
the specimen critical point. The damage field is obtained
with criterion (6) on the basis of the analysis of the
L. B. GETSOV et al.: THERMOCYCLIC- AND STATIC-FAILURE CRITERIA FOR SINGLE-CRYSTAL SUPERALLOYS ...
258 Materiali in tehnologije / Materials and technology 48 (2014) 2, 255–260
strain-field evolution using the experimental data on the
creep and elastoplastic deformation-resistance curves.
The typical damage-field distribution after the first
thermal cycle (20 °C  Tmax = 900 °C  Tmin = 150 °C)
is presented in Figure 5 for sample 5-1 of ZhS36 with
orientation 001.
The results of the finite-element simulations show
that the crystallographic orientation has a significant
influence on the stress-strain state of the samples
(Figure 6), as confirmed also by the experiments.12 The
width of the hysteresis loops and the unilaterally accu-
mulated strain are also very sensitive to the thermal
cyclic regime (Figure 7).
The numbers of the cycles to crack initiation, calcu-
lated on the basis of criterion (6) using different equiva-
lent strains, (11)–(14), are given in Table 3. A satisfac-
tory correlation is observed between the criterion
predictions and the results of the experiments (without a
sufficient statistical representation).
7 CONCLUSIONS
1. In the investigations of I. L. Svetlov, E. R. Golubov-
sky and other researchers a series of tests were
conducted regarding the definition of the thermal
fatigue resistance and short-term creep of the single-
crystal ZhS32 alloy, determining the temperature
range causing the changes in the failure modes.
2. The failure criteria for the single-crystal alloys under
static and thermal cyclic loadings are proposed and
discussed. A satisfactory correlation is observed bet-
ween deformation criterion (6) and the obtained
experimental results. All the considered variants of
equivalent strains (11)–(14) give practically the same
results. The criterion using von Mises strain intensity
(13) offers the most conservative estimation.
3. The finite-element simulations of single-crystal spe-
cimens under thermal cyclic loading were performed
using different material models. The obtained results
indicate an applicability of the proposed deformation
criteria of failure for the single-crystal ZhS32 and
ZhS36 alloys for the temperatures up to 1100 °C.
L. B. GETSOV et al.: THERMOCYCLIC- AND STATIC-FAILURE CRITERIA FOR SINGLE-CRYSTAL SUPERALLOYS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 255–260 259
Figure 7: Influence of the temperature program on the cyclic stress-
strain curve. Results of the finite-element simulations of the speci-
mens with the 001 orientation (the central point of a specimen).
Slika 7: Vpliv temperaturnega programa na krivuljo napetost – raz-
tezek. Rezultati simulacije s kon~nimi elementi za vzorce z orientacijo
001 (sredinski del vzorca).
Figure 5: Damage-field distribution after the first cycle for sample 5-1
with the 001 orientation
Slika 5: [irjenje podro~ja po{kodbe po prvem ciklu pri vzorcu 5-1 z
orientacijo 001
Table 3: Deformation-criterion (6) predictions vs. experimental results for the crack initiation life
Tabela 3: Napovedi merila deformacij (6) v odvisnosti od eksperimentalnih rezultatov za za~etek nastanka razpoke
Orientation Tmax/°C Tmin/°C
Number of cycles to crack initiation
eq = nl (11) eq = 1 (12) eq = i (13) eq = max (14) Experiment
Sample 5-1 900 150 338 275 195 280 435
Sample 5-3 1000 500 218 196 150 172 305
Sample 1-2 900 150 85 73 59 75 190
Sample 2-1 900 150 15 9 10 15 17
Sample 1-1* 900 150 5 4 3 4 10
Sample 2-6* 1000 500 61 44 56 57 10–130
Sample 4-1* 900 150 6 4 4 5 10
*Specimen with a concentrator (the radius of the central hole is 0.25 mm)
Figure 6: Influence of the crystal orientation on the cyclic stress-
strain curve. Results of the finite-element simulations of the thermal
cycles with Tmax = 900 °C and Tmin = 150 °C (the central point of the
specimen).
Slika 6: Vpliv orientacije kristala na cikli~no krivuljo napetost – raz-
tezek. Rezultati simulacije s kon~nimi elementi za toplotne cikle s
Tmax = 900 °C in Tmin = 150 °C (sredinski del vzorca).
Acknowledgements
The study was supported by the Russian Fundamen-
tal Research Program, Project No.12-08-00943. The
authors are also grateful to S. M. Odabai-Fard for help-
ing prepare the paper.
8 REFERENCES
1 E. N. Kablov, E. R. Golubovsky, Heat-resistant nickel alloys,
Mechanical Engineering, Moscow 1998
2 R. E. Shalin, I. L. Svetlov, E. B. Kachanov, V. N. Toloraiya, O. S.
Gavrilin, Single crystals of nickel-base superalloys, Mechanical
Engineering, Moscow 1997
3 L. B. Getsov, N. I. Dobina, A. I. Rybnikov, A. S. Semenov, A. A.
Staroselsky, N. V. Tumanov, Thermal fatigue resistance of single cry-
stal alloy, Strength of Materials, (2008) 5, 54–71
4 L. B. Getsov, A. I. Rybnikov, A. S. Semenov, Thermal fatigue
strength of heat-resistant alloys, Thermal Engineering, 56 (2009),
412–420
5 L. B. Getsov, Materials and strength components of gas turbines,
Rybinsk, Publ. House, Gas Turbine Technology, 1–2 (2011)
6 L. B. Getsov, A. S. Semenov, Failure criteria for polycrystalline and
single crystal materials under thermal cyclic loading, Proc. NPO
CKTI, N296, Strength of materials and resource elements of power,
St. Petersburg, 2009, 83–91
7 L. B. Getsov, B. Z. Margolin, D. G. Fedorchenko, The determination
of elements of engineering safety margins in the calculation of
structures by finite element method, Proc. NPO CKTI, N296,
Strength of materials and resource elements of power, St. Petersburg,
2009, 51–66
8 A. S. Semenov, PANTOCRATOR-finite-element software package,
focused on the solution of nonlinear problems in mechanics, Proc. of
V Int. Conf. Scientific and technical problems of forecasting the
reliability and durability of the structures and methods for their
solution, Publishing House of Polytechnic University, St. Petersburg
2003, 466–480
9 G. Cailletaud, A Micromechanical Approach to Inelastic Behaviour
of Metals, Int. J. Plast., 8 (1991), 55–73
10 J. Besson, G. Cailletaud, J. L. Chaboche, S. Forest, Non-linear
mechanics of materials, Springer, 2010
11 U. F. Kocks, T. J. Brown, Latent hardening in aluminium, Acta
Metall., 14 (1966), 87–98
12 L. Getsov, A. Semenov, A. Staroselsky, A failure criterion for single
crystal superalloys during thermocyclic loading, Mater. Tehnol., 42
(2008) 1, 3–12
L. B. GETSOV et al.: THERMOCYCLIC- AND STATIC-FAILURE CRITERIA FOR SINGLE-CRYSTAL SUPERALLOYS ...
260 Materiali in tehnologije / Materials and technology 48 (2014) 2, 255–260
U. KAV^I^ et al.: UHF RFID TAGS WITH PRINTED ANTENNAS ON RECYCLED PAPERS AND CARDBOARDS
UHF RFID TAGS WITH PRINTED ANTENNAS ON
RECYCLED PAPERS AND CARDBOARDS
UHF RFID-ZNA^KE Z NATISNJENIMI ANTENAMI NA
RECIKLIRANEM PAPIRJU IN KARTONU
Ur{ka Kav~i~1, Matej Pivar2, Miloje \oki}2, Diana Gregor Svetec2,
Leon Pavlovi~3, Tadeja Muck2
1Valkarton Rakek, d. o. o., Rakek, Slovenia
2University of Ljubljana, Faculty of Natural Sciences and Engineering, Ljubljana, Slovenia
3University of Ljubljana, Faculty of Electrical Engineering, Ljubljana, Slovenia
urskavcic@gmail.com
Prejem rokopisa – received: 2013-05-30; sprejem za objavo – accepted for publication: 2013-07-03
The integration of passive RFID tags in different applications is important in order to increase product functionality. The present
research was focused on the optimization of the printing process conditions for the screen printing of passive UHF RFID
antennas for the box tracking in logistics and for newspaper tracking in the retail trade. The antennas were printed on uncoated
and coated recycled papers and coated cardboards. Two different conductive inks were applied with a semi-automatic screen
printer. Drying conditions were varied in order to obtain a good print quality and the appropriate electrical properties of the
conductive printed layer on all the printing substrates. The integration of flip chips was applied to the printed antennas. At the
end of our research, an analysis of the printed antennas and the final UHF RFID tags was carried out. The quality of the printed
antennas was first evaluated by image analysis, after which the electrical properties, such as the impedance and radiation
patterns, were measured. To analyse the quality of UHF RFID tags, the maximum reading length was also determined.
We demonstrated that working UHF RFID tags with screen printed antennas can be realized on substrates with lower quality,
such as uncoated recycled papers. The main influence on the final working tag is the quality of the conductive ink itself.
Keywords: printed RFID antenna, tag integration, recycled paper, cardboard
Vklju~itev pasivnih RFID-zna~k v razli~ne aplikacije je pomembna, da dose`emo dobro funkcionalnost izdelka. Predstavljena
raziskava je bila osredinjena na optimizacijo razmer sitotiskarskega procesa, s katerim so bile natisnjene UHF RFID-antene za
sledenje embala`e in ~asopisov v trgovinskih verigah. Antene so bile natisnjene na nepremazane in premazane reciklirane
papirje in na premazane embala`ne kartone. Uporabljeni sta bili dve prevodni tiskarski barvi, ki sta bili na tiskovni material
naneseni s polavtomatskim sitotiskarskim strojem. Razmere pri su{enju so bile optimizirane, tako da so bile dose`ene dobra
kakovost in primerne elektri~ne lastnosti odtisov na vseh tiskovnih materialih. Na tiskanih antenah je bilo izvedeno
kontaktiranje ~ipov. Na koncu raziskav je bila izvedena analiza natisnjenih anten in kon~no izdelanih zna~k. Kakovost tiskanih
anten je bila ovrednotena s slikovno analizo ter z merjenjem impedance in sevalnega diagrama anten. Za dolo~anje kakovosti
kon~ne zna~ke je bila dolo~ena {e maksimalna razdalja od~itavanja.
Dokazali smo, da je UHF RFID-zna~ke mogo~e tiskati s sitotiskom tudi na nizkocenovne tiskovne materiale, kot je nepremazan
recikliran papir. Glavni vpliv na delovanje kon~ne zna~ke ima kakovost prevodne tiskarske barve.
Klju~ne besede: tiskane RFID-antene, izdelava RFID-zna~ke, recikliran papir, embala`ni karton
1 INTRODUCTION
The future of automatic identification and data cap-
turing in the field of packaging is increasingly dedicated
to radio-frequency identification (RFID), which enables
wireless identification and saves a lot of time (and
money) in comparison to barcode technology. RFID is an
automated identification technology that consists of a
reader, a reader antenna and a tag (which consists of an
antenna and a chip). It uses radio waves to transfer the
information between the reader and the tag at long dist-
ances. The RFID tag can work at different frequencies:
low (LF: 125 kHz or 134 kHz), high (HF: 13.56 MHz) or
ultrahigh (UHF: 860–960 MHz) radio-frequency ranges.
Due to the fact that the higher the frequency, the larger
the distance at which the information can be read, in
packaging identification UHF RFID tags are mostly
used.
The RFID tag can be produced conventionally by an
etching process or can be printed. While the conven-
tional production of RFID tags is still very expensive and
environmentally unfriendly, many researchers are trying
to produce printed RFID tags, where the electronic com-
ponents are printed in-line, roll-to-roll, in the same pro-
cess as the packaging layout itself. The printing, unlike
conventional etching, is an additive process, and is far
more ecological and economical than the subtractive
etching process. This could reduce the amount of
material used for production, i.e. the waste that is a side
product of production, and consequently the total cost
per tag. RFID antennas can be printed with different
printing technologies:1 offset lithography, flexography,
gravure, ink jet, electrophotography and screen printing.
Different printing technologies enable different accuracy,
resolution and ink thickness. Most research has been
carried out using inkjet2–5 and screen6–9 printing techno-
logies. There have also been some using gravure print-
ing,10,11 but less with offset and flexography. Many
research programmes have also been undertaken to test
the performance of RFID printed antennas12–19 and chip
Materiali in tehnologije / Materials and technology 48 (2014) 2, 261–267 261
UDK 621.396.44:621.396/.398:655.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)261(2014)
bonding.20,21 Printed antennas are usually applied to
different foils22,23 or photo papers. There are also some
researches made on the field of printed paper-based
RFID or sensors,17,24–29 but none of them analysed the
printing RFID antennas on recycled paper and card-
board. The research in that field analysed the printing
substrates from the electrical point of view and not from
the graphical point of view. For this reason, in this paper
the application of silver conductive ink to recycled
papers and packaging cardboards is presented.
The goal of our research was to optimize the printing
process and drying conditions for the screen printing of
passive UHF antennas directly on packaging and paper
substrates. The antennas were printed with silver con-
ductive printing inks using a semi-automatic screen
printer. After that the silicon chip was applied to the
antenna. At the end of our research, the analysis of the
printed antennas and final UHF RFID tags was carried
out. The quality of the printed antennas was evaluated by
image analysis and the resistance to abrasion, which is
important if the antennas printed on packaging or on
newspaper are not additionally protected with another
protective layer. At the end the impedance and radiation
patterns were measured. To analyse the quality of the
UHF RFID, the tag reading length was determined by
measuring the received power in watts.
2 EXPERIMENTS
The current investigation involved the selection of the
UHF RFID antenna for the central frequency of 868
MHz (Figure 1), antenna printing, drying optimization,
analysis of antenna printability, mounting the chip onto
antenna and an analysis of RFID tag reading.
2.1 Printing
The UHF RFID antennas were printed with two con-
ductive printing inks: SunChemical30 (CRSN2442, Sun-
Tronic Silver 280, Thermal Drying Silver Conductive
Ink) and DuPont31 (DuPont 5064H silver conductor). As
printing substrates, two coated cardboards and two
recycled papers (one coated and one uncoated) were
used. A RokuPrint semi-automatic screen printer and
monofilament polyester plain weave mesh with 120
L/cm were used (theoretical ink volume 16.3 cm3/m2).
The properties of the printing inks and printing sub-
strates are presented in Tables 1 and 2.
Table 1: Printing inks properties30,31
Tabela 1: Lastnosti tiskarskih barv30,31
Property SunChemical DuPont
Solids 69–71 63–66
Viscosity 2–3 10–20
Sheet resistance: R/m for
25 ìm) 10–32
6 on 125 μm
PET film
Drying
conditions Static box oven
150 °C/
30 min
130 °C/
(10–20 min)
Reel-to-reel 100–130 °C/(30–90 s)
140 °C/
(2 min)
2.2 Print penetration
The print penetration was determined in accordance
with the IGT–W24 and ICP–T17 methods. At the mo-
ment of printing a quantity of ink or varnish is absorbed
by the surface of the paper. This amount is determined
by the absorption of the liquid in the surface recesses
(roughness) and the absorption into the paper pores at
the surface. With the IGT test method the sum of the two
phenomena is determined: the oil absorption or varnisha-
bility. The reciprocal value of this is called print penetra-
tion. A large stain indicates a low roughness/absorption
of the paper.
The final values for the print penetration were calcu-
lated according to Equation 1, where PP is the print
penetration and l is stain length of the print in mm:
PP = 103/l (1)
2.3 Optimization of drying conditions
After printing, the optimization of the drying condi-
tions was determined to achieve the lowest sheet resi-
stance of the prints. Based on experiments, a two-stage
drying process was determined as optimal (Figure 2).
The samples were dried according to the printing sub-
strate and the printing ink. The optimal drying was
U. KAV^I^ et al.: UHF RFID TAGS WITH PRINTED ANTENNAS ON RECYCLED PAPERS AND CARDBOARDS
262 Materiali in tehnologije / Materials and technology 48 (2014) 2, 261–267
Figure 1: Antenna printing form
Slika 1: Tiskana oblika antene
Table 2: Properties of printing substrates
Tabela 2: Lastnosti podlag za tiskanje
Standard Type
Grammage
(g/m2)
Thickness
(ìm)
Roughness (ìm)
(Stylus TR200)
Porosity (mL/min)
(Bendtsen)
Water absorpti-
veness (g/m2)
(Cobb60)
ISO 536 ISO 534 ISO 4287 ISO 5636-3 ISO 535
paper 1 coated recycled paper 60 60 2.226 16.00 not determinable
paper 2 uncoated recycled paper 60 113 5.706 381.00 not determinable
cardboard 1 coated cardboard 245 400 0.463 27.33 5.609
cardboard 2 coated cardboard 350 525 0.482 12.00 4.427
determined as the point where the sheet resistance of the
printed samples became constant and did not get any
lower with longer drying times or higher temperatures.
2.4 Analysis of printed conductive ink layer
After printing and the two-stage drying process, the
analysis of the printed conductive ink layer was carried
out. The print mottle, abrasion and sheet resistance of the
printed ink layer were determined.
2.4.1 Print mottle
The print uniformity (print mottle) was determined
using image analysis (ImageJ software). The print mottle
was determined with a traditional STFI method, by
calculation of the coefficient of variation (CV), where 
is the standard deviation of the grey values and R is the
mean grey value:32
CV/% = /R × 100 (2)
2.4.2 Abrasion
The abrasion was determined on the samples (5.1 cm
× 23 cm) using a Param RT-01 Rub tester instrument
(according to the standard ASTM D 5264). The printed
samples were positioned on the table, and the unprinted
specimen of the same printing substrate was mounted on
the weight, which rubs the printed sample. The test
duration is determined by the number of strokes (a stroke
is one back-and-forth cycle) the sample is rubbed. The
speed of the rub was 106 cycles per minute using a mass
0.9 kg. The abrasion was determined on the receptor
surface of 4.8 cm × 10 cm size using image analysis after
100 (paper) or 500 (cardboard) strokes. The uncoated
(and coated) recycled paper has a rougher surface than
the coated cardboard and consequently the abrasion of
the printing ink on the paper was much higher than on
the cardboard. If 500 strokes were applied to paper, the
abrasion would not be determinable using image ana-
lysis. On the other hand, after 100 strokes on cardboard
there is no evident abrasion. That is why only 100
strokes were applied to paper and 500 to cardboards. The
results present the proportion of the area coated with the
rubbed ink on the receptor surface.
2.4.3 Resistance
The resistance was measured on the test element
presented in Figure 3 between points 1 and 2 using a
DT-890G multimeter. The nominal length was L = 22
mm and width W = 3 mm. The nominal number of
squares was Nsq = L/W = 10.3. The final results are given
as the value of the conductive printed layer sheet resi-
stance R (m).
2.5 RFID tag analysis
2.5.1 Antenna evaluation
The numerical simulation of the antenna impedance
was carried out using a commercial 3-D solver (Ansoft
HFSS). In addition, the radiation patterns in the E (elec-
tric field) and H (magnetic field) planes for the paper and
cardboard printed antennas were measured at an outdoor
antenna-measuring polygon.
2.5.2 RFID tag reading
After the antennas had been printed and analysed, the
strap chip (NXP) with impedance Z = 22 – j195  33 was
assembled onto the printed antennas. Then the analysis
of the RFID tag was performed using an IDS-R902
reader (IDS Microchip, Ljubljana, Slovenia). The
IDS-R902 reader consists of an IDS reader and a Patch
A0025 antenna (Poynting GmbH, Dortmund, Germany).
The reader is based on the IDS-R902 circuit, supports
the ISO18000-6 C or EPC Gen 2 Protocol and measures
the strength of the modulated signal backscattered from
the tag. The reader antenna (gain 4.5 = 6.5 dBi) emits
circularly polarized UHF radiation with a frequency f =
867 MHz. Its output power is 400 mW (+26 dBm). The
reader uses an amplitude shift keying and has a maxi-
mum input sensitivity of 25 pW (–76 dBm). The quality
of the final UHF RFID tag was evaluated by measuring
the power in W (dBm) for every 5 cm by moving the tag
straight from the reader.
U. KAV^I^ et al.: UHF RFID TAGS WITH PRINTED ANTENNAS ON RECYCLED PAPERS AND CARDBOARDS
Materiali in tehnologije / Materials and technology 48 (2014) 2, 261–267 263
Figure 3: Test element for resistance measurement
Slika 3: Preizkusni element za merjenje upornosti
Figure 2: Schematic representation of two-stage (combined) drying process
Slika 2: Shematski prikaz dvostopenjskega su{enja
3 RESULTS AND DISCUSSION
3.1 Print penetration
Print penetration is a measure of the penetration
velocity of the printing ink into the printing substrate. It
represents the properties (roughness, porosity and
absorptiveness) of the printing substrate. Figure 4 shows
that recycled papers have a higher print penetration than
cardboards. Again, the higher penetration can be
detected with uncoated recycled paper, which can be
directly connected to the roughness and print mottle.
3.2 Drying optimization
The drying optimization was determined as the point
where the resistance became constant with a higher
temperature or longer drying. The optimal drying condi-
tions were determined as presented in Table 3. Samples
were first dried in the hot zone and then they were
exposed to the Heat&Press process (Figure 2).
During the drying process, the printing substrate
exposed to high temperatures changes its surface colour,
which is significant when the final product is graphically
designed, such as with packaging, where the colours and
final product’s appearance are very important. The
colour difference (	E) of the printing substrate was
determined for both printing inks (Figure 5). It is clear
that the printing substrates printed with DuPont printing
ink change their colour more than the samples printed
with SunChemical printing ink, due to the higher heating
temperature. Even so, all the values are very low and the
colour change is hard to detect with the eye. The printing
with SunChemical printing ink is also more appropriate
from an ecological point of view, because of the lower
drying temperature.
3.3 Analysis of printed conductive ink layer
3.3.1 Print mottle
In Figure 6 the print mottle for all the printing
substrates is presented. The print mottle is dependent on
the surface properties of the printing substrate. Usually,
it is the result of an uneven ink layer or non-uniform ink
absorption across the printing substrate.34 The higher
roughness (Table 2), print penetration (Figure 4) and
water absorptiveness (Table 2) of the two recycled
papers affect the higher print mottle, as presented in Fig-
ure 6. The difference between the two conductive inks is
also obvious, with the DuPont printing ink having a
slightly higher non-uniformity. The final antenna perfor-
mance is also dependent on the print quality of the
printed conductive lines. If the antenna is printed onto
rough substrates the conductive Ag particle in ink could
not be connected and the final conductivity could be
questionable. In that case the final antenna performance
would not be as good as expected based on the antenna
design and simulation.
U. KAV^I^ et al.: UHF RFID TAGS WITH PRINTED ANTENNAS ON RECYCLED PAPERS AND CARDBOARDS
264 Materiali in tehnologije / Materials and technology 48 (2014) 2, 261–267
Figure 5: 	E of printing substrates after two stages drying process
Slika 5: 	E podlag za tiskanje po dvostopenjskem su{enju
Figure 4: Print penetration of the printing substrates
Slika 4: Tiskarska penetracija na podlagah za tiskanje
Table 3: Drying conditions
Tabela 3: Razmere pri su{enju
Hot zone Heat&Press
Printing substrate Time (s) Temperature (°C) Printing ink Time (s) Temperature (°C)
Cardboard 135 115 SunChemical 10 150
Paper 90 115 DuPont 10 190
Figure 6: Print mottle
Slika 6: Tiskovna neenakomernost
3.3.2 Abrasion
The abrasion of the printed antennas is important,
while prints have to have good conductivity, which is
degraded if the prints’ abrasion is too high. In that case it
is possible to get a inhomogeneous printed conductive
layer and small holes can appear on the surface of the
antenna, because of the printing substrate roughness.
In Figure 7 it is clear that the prints printed on papers
have much higher abrasion (at lower rub) than those
printed on cardboard. It is also evident that the DuPont
printing ink has higher abrasion than the SunChemical
ink. On paper (especially on uncoated recycled paper),
the binder penetrates more quickly and deeply into the
substrate than on coated cardboard, and consequently the
conductive silver particles remain unbounded at the
surface. Moreover, the surface of the paper is rougher
(Table 2) and when rubbed more the particles remain on
the receptor (unprinted specimen), and the abrasion is
higher.
3.3.3 Sheet resistance
The sheet resistance was calculated and is presented
in Figure 8. All the samples achieved a good sheet resi-
stance, i.e., lower than R = 100 m, with the antennas
printed with DuPont printing ink being even lower than
those printed with SunChemical ink. In the inks’ speci-
fication (Table 1), the specified resistances are much
lower than those presented in Figure 8. This is because
the resistances in Table 1 correspond to a thick dry film
25 μm, while the measured thicknesses of the dry film in
our experiment were much lower (the thickness was
determined on a cross-section of the printed samples),
between 6 μm and 10 μm, and consequently the sheet
resistance is higher.
3.4 RFID tag analysis
3.4.1 Antenna evaluation
The radiation patterns in the E- and H-planes were
evaluated by measurements of antennas on two printing
substrates. The uncoated recycled paper (paper 2) and
coated cardboard (cardboard 1) were selected as they had
the highest and the lowest print mottles (Figure 6) and
abrasions (Figure 7) of the four samples. The samples
printed with the SunChemical printing ink were selected
on the basis of lower heating temperatures, lower abra-
sion and smaller print mottle in comparison with the
DuPont printing ink. The results revealed that the
patterns are nearly the same for both measured antennas,
regardless of which printing substrate was used (Figure
9).
The measured radiation patterns have slight irregula-
rities in the measurement, especially at the nulls. This is
U. KAV^I^ et al.: UHF RFID TAGS WITH PRINTED ANTENNAS ON RECYCLED PAPERS AND CARDBOARDS
Materiali in tehnologije / Materials and technology 48 (2014) 2, 261–267 265
Figure 9: Radiation pattern for antennas printed on both cardboard
and on recycled paper and its simulation
Slika 9: Sevalni diagram antene, natisnjene na embala`nem kartonu in
recikliranem papirju ter njena simulacija
Figure 7: Abrasion
Slika 7: Abrazija
Figure 8: Sheet resistance of the prints
Slika 8: Plastna upornost odtisov
Figure 10: Simulated antenna impedance
Slika 10: Simulirana impedanca antene
due to the disturbance to the antenna and the obstruction
of the attached balun that was used for the balanced
antenna to the coaxial cable interface. Without the balun
interference, the radiation patterns should be approxima-
tely the same as the simulated ones shown in Figure 9.
Besides the radiation pattern, the impedance of the
antenna was simulated (Figure 10). The simulated an-
tenna impedance of about 15 + j180  at 868 MHz
ensures a good match to the RFID chip impedance (22 -
j195 ).
3.4.2 RFID tag reading
The NXP chips on strap were attached to the printed
antennas and the largest reading distances with the
corresponding power values were measured in a real
environment on five samples, as presented in Figure 11.
The printed antenna was oriented horizontally to the
transmitting antenna. Figure 12 shows that the power
(an average of 5 measurements) diminishes with the
distance of reading from around 10 nW to almost 100
pW (-50 dBm to almost -70 dBm) at one metre distance.
The largest measured distance where the tags still
worked was 105 cm. It was observed that the power of
the tags printed on paper was a little lower than that of
the tags printed on cardboard.
4 CONCLUSIONS
The analysed influences of drying conditions on the
quality of the printed conductive layers were proven to
be significant, not only in terms of the quality of the
printing but also on the relevant electrical parameters
measured on the printed antennas. A higher drying tem-
perature increases the final conductivity. The paper and
cardboard enable drying to 150 °C or a maximum of 200
°C. A higher drying temperature causes the substrate to
become yellow.
The two-stage drying process with Heat&Press
makes it possible to use a lower drying temperature and a
shorter drying time. Comparing the printing inks, the
DuPont ink has a slightly lower resistance, which is its
only advantage over the SunChemical ink. On the other
hand, the weakness of the DuPont ink is that it requires a
higher drying temperature. SunChemical ink shows
lower print mottle and has lower abrasion. It means that
antennas printed with SunChemical ink are more uni-
form on their surface and more mechanically stable.
The printing substrate, especially the coating, has
significant influences on better ink formation, higher
uniformity and also on better abrasion resistance. This
effect is visible especially on coated recycled paper and
cardboards compared with uncoated recycled paper.
The design of the UHF antenna is appropriate for
both printing substrates – no differences in radiation
pattern were observed between antennas printed on
coated cardboard and those on uncoated recycled paper.
By the end of our research we had shown that work-
ing UHF RFID tags with screen-printed antennas can be
realized on substrates of extremely low quality, such as
uncoated recycled papers. The main influence on the
final working tag is the quality of the conductive ink
itself.
Acknowledgements
The authors express their gratitude to Ralf Zichner
(Fraunhofer), and to the companies Ams R&D and RLS.
Thanks are also due to Avery Dennison for permission to
use their antenna design. Ur{ka Kav~i~ acknowledges
assistance from the European Social Fund ("Operation
part-financed by the European Union, European Social
Fund").
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U. KAV^I^ et al.: UHF RFID TAGS WITH PRINTED ANTENNAS ON RECYCLED PAPERS AND CARDBOARDS
266 Materiali in tehnologije / Materials and technology 48 (2014) 2, 261–267
Figure 11: Measurements of tag readability
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Materiali in tehnologije / Materials and technology 48 (2014) 2, 261–267 267

T. KODELJA et al.: TOPMOST STEEL PRODUCTION DESIGN BASED ON THROUGH PROCESS MODELLING ...
TOPMOST STEEL PRODUCTION DESIGN BASED ON
THROUGH PROCESS MODELLING WITH ARTIFICIAL
NEURAL NETWORKS
PROJEKTIRANJE PROIZVODNJE VRHUNSKIH JEKEL NA
PODLAGI MODELIRANJA SKOZI PROCES Z UMETNIMI
NEVRONSKIMI MRE@AMI
Tadej Kodelja1, Igor Gre{ovnik1,2, Robert Vertnik2,3, Bo`idar [arler1,2,4
1Center of Excellence BIK, Solkan, Slovenia
2University of Nova Gorica, Nova Gorica, Slovenia
3[tore Steel d.o.o., Research, [tore, Slovenia
4IMT, Ljubljana, Slovenia
bozidar.sarler@imt.si
Prejem rokopisa – received: 2013-06-14; sprejem za objavo – accepted for publication: 2013-07-03
Application of artificial neural networks for modeling of a complete process path in a steel production – from the scrap steel to
the material properties of semi products – is presented. The described approach is introduced as an alternative to physics based
through process modeling, with the advantage of lower complexity of the software and much lower computing times for
calculating the influence of a specific settings of the process parameters. This new approach can be beneficially used in
designing the production process. This is clearly demonstrated by estimating the influence of 34 alloying elements and process
parameters of 6 process steps on 5 final mechanical properties of spring steel (elongation, tensile strength, yield stress, hardness
after rolling and necking), based on 1879 recorded data sets from the production line in [tore Steel company. The ANN used is
of a multilayer feedforward type with sigmoid activation function and supervised learning. An important feature of this
approach is its dependence on accurate and sufficient data, acquired from the modeled process. Therefore, special care must be
devoted to validation of the obtained model and error estimation. The reliability and other characteristics of the available data
can vary to a great extent in real industrial practice, therefore analysis of the models is a highly customized task that has to be
performed on a case to case basis. A flexible and easily extensible software base has been developed in the scope of the
described work in order to adequately support research, development and practical application of this kind of models.
Keywords: steel production, mechanical properties of steel, artificial neural networks, response approximation, feed forward
networks with back propagation
Predstavimo uporabo umetnih nevronskih mre` za modeliranje celotne procesne poti izdelave jeklenih polizdelkov – od rene do
snovnih lastnosti polizdelkov. Opisani pristop je vpeljan kot alternativa fizikalnemu modeliranju skozi proces s prednostjo
manj{e kompleksnosti programske opreme ter bistveno manj{imi ra~unskimi ~asi za izra~un vpliva specifi~ne nastavitve
procesnih parametrov. Tak{en pristop se lahko s pridom uporablja pri na~rtovanju proizvodnega procesa. To je nazorno
prikazano pri oceni vpliva 34 legirnih elementov in procesnih parametrov 6 procesnih korakov na 5 kon~nih snovnih lastnosti
vzmetnega jekla (raztezek, natezna trdnost, meja te~enja, trdota po valjanju in skr~ek), na podlagi 1879 zabele`enih podatkovnih
setov iz proizvodne linije podjetja [tore Steel. Uporabljena je usmerjena nevronska mre`a s sigmasto aktivacijsko funkcijo in
nadzorovanim u~enjem. Pomembna zna~ilnost tega pristopa je njegova odvisnost od pravilnih in zadostnih podatkov,
pridobljenih iz procesa. Zato se je potrebno posebej posvetiti validaciji pridobljenega modela in oceni napak. Zanesljivost in
ostale zna~ilnosti razpolo`ljivih podatkov, pridobljenih iz realnih industrijskih procesov, se obi~ajno zelo razlikujejo, zato je
analiza tak{nih modelov zelo specifi~na in mora biti narejena od primera do primera. V ta namen je bila izdelana fleksibilna in
enostavno raz{irljiva programska oprema, ki omogo~a primerno podporo raziskavam, razvoju in prakti~ni uporabi tovrstnih
modelov.
Klju~ne besede: izdelava jekla, mehanske lastnosti jekla, umetne nevronske mre`e, aproksimacija odziva, nevronske mre`e s
povratnim raz{irjanjem napak
1 INTRODUCTION
Controlling the final mechanical properties of
products or semi products is very important for steel
production companies. This is a difficult task because
there are a number of sequentially connected processes
where the output of one process is an input to the next
one. Different physics based numerical models can be
used to predict the outcomes, but their development can
be very complicated and time consuming.1,2 Artificial
neural networks (ANN) based models3,4 can be used as
an alternative to these physics based numerical models.
Over the last years, ANNs have been successfully used
across an extraordinary range of problem domains.
Examples can be found in almost all fields of industry as
well as in research areas that show promise for the
future.5 ANNs are already being used in steel production
industries in modeling of blast furnace,6 continuous
casting, steel rolling,7 etc. The first use of ANN in
modeling of the entire production path (also referred to
as “through process modeling”) has been demonstrated
for production of aluminum foil in8. Furthermore, a pre-
liminary study9 was made for complete steel production
path, while in this study, additional parametric studies
and sensitivity tests were added. The main drawback of
Materiali in tehnologije / Materials and technology 48 (2014) 2, 269–274 269
UDK 669.1:004.032.26 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)269(2014)
ANN models over physics based models is the fact that
they can be used based on the specific training data only,
and do not allow generalization to different production
plants. Only the developed methodology is transferable.
In the present paper, we study the possibility of using
ANN-based models as a comprehensive decision support
tool in steel production. We explore the prediction of
important mechanical properties of steel (elongation,
tensile strength, flow limit, hardness and shrinkage)
based on values of influential process parameters that
determine the complete steel production path. The steel
manufacturing process in the [tore Steel company and
the respective available data were considered10 as a basis
for the present study. The manufacturing process path
consists of six individual processes:11,12 steel making,
continuous casting of steel, hydrogen removal, reheating,
multiple stage rolling, and cooling on the cooling bed.
Each of these processes can be independently modeled
by a physics based numerical model.13–21 The state of the
steel (shape, microstructure) of an individual process
influences the downstream processing (subsequent pro-
cesses in the process chain) and thus act as a part of
input data (e.g. defining initial or boundary conditions)
in the model of that process. This is schematically repre-
sented in Figure 1. In the current work we use another
approach where an ANN is used to build a complete
model of the whole production chain. We model the
outcomes after the last process step and relate them to
process parameters defining all processes involved in the
production path. After the model is built, we can explore
the effect of variation of process parameters to the final
material properties, e,g. by changing process parameters
independently in parametric and sensitivity tests and
observing model outputs.
2 MODELING SOFTWARE
A software for construction and use of ANN-based
models has been developed in the scope of this work.
The software was designed to match the challenges and
requirements met when solving this kind of problems. In
particular, it has to provide good flexibility in designing
training strategies, filtering training data, verification of
results, testing different network layouts, integration
with other software, etc. This is crucial when approxi-
mating behavior of steel processing systems with large
number of processing parameters. Data obtained from
such systems is often inaccurate or even corrupted due to
practical limitations in acquisition procedures. Response
sampling can not be planned in advance but is accom-
modated to production schedules in the factory, therefore
information available may be deficient in some regions
of parameter space in order to obtain good response
approximation and therefore verification of results plays
an important role. The software platform has been elabo-
rated in22,23.
The Aforge.Net library is used as ANN framework.24
A convenient characteristics of neural networks is that
approximation can be performed in two separate stages
(Figure 2). In the training stage, the network is trained
by using the sampled response (either measured or
calculated by a numerical model). In the approximation
stage, trained network is used for all subsequent calcu-
lations of approximated response at arbitrary values of
input parameters. This gives the neural networks an
important advantage over other modeling techniques,
since the second stage if very fast as compared to the
first stage. The software takes full advantage of this
feature by separating these stages. This is especially
T. KODELJA et al.: TOPMOST STEEL PRODUCTION DESIGN BASED ON THROUGH PROCESS MODELLING ...
270 Materiali in tehnologije / Materials and technology 48 (2014) 2, 269–274
Figure 2: Approximation with neural networks: training a network
with presented data pairs (top) and calculation of approximated
response with trained network (bottom)
Slika 2: Aproksimacija z nevronskimi mre`ami: u~enje mre`e na
podlagi podatkovnih parov (zgoraj) in izra~un aprokismiranega odziva
z nau~eno mre`o (spodaj)
Figure 1: Steel manufacturing process modeling strategy
Slika 1: Strategija modeliranja procesa izdelave
important when performing extensive analyses of the
considered process on the basis of the developed ANN
models, or when incorporating the models in automatic
optimization procedures.25,26
3 CONSTRUCTION OF THE ANN-BASED
PROCESS MODEL
In the considered production setup from the [tore
Steel company, the complete process is defined by 123
influential parameters (Table 1). There are 24 parame-
ters defining the steel grade, 12 process parameters
defining the continuous casting, 2 parameters the hydro-
gen removal, 4 parameters the reheating furnace, 31
parameters the rolling mill, 43 parameters the continuous
rolling mill, and 7 parameters the cooling bed. On the
output side, five mechanical properties of the final pro-
duct are observed and represent the output values of the
model (Table 2).
Table 1: Process parameters (input)
Tabela 1: Procesni parametri (vhod)
Processes / properties Number of parameters
Composition 24
Continuous casting of steel 12
Hydrogen removal 2
Billet reheating furnace 4
Rolling mill 31
Continuous rolling mill 43
Cooling bed 7
Total 123
Table 2: Material properties (output)
Tabela 2: Snovne lastnosti (izhod)
Final mechanical
properties
Elongation (A)
Tensile strength (Rm)
Yield stress (Rp0.2)
Hardness after rolling (HB)
Necking (Z)
For construction of the models, data was manually
collected from different databases representing produc-
tion of the steelwork in year 2011. Data was first sepa-
rated for two billet dimensions (140 mm and 180 mm)
which undergo considerably different process parame-
ters. In addition, the data had to be filtered by applying a
number of specially designed criteria in order to exclude
corrupted data and overshoots. After these procedures, a
total of 1879 data sets for dimension 140 mm have been
prepared and used in the training procedure.
This data was randomly divided into disjoint training
and verification sets. Training data was then used in
training a feed forward neural network with sigmoid acti-
vation function, in which we iteratively minimize error
of the model on this data by the back propagation algo-
rithm. After the convergence was achieved, the model
was validated on the verification set that was not
involved in the training, in order to estimate its accuracy
(Figure 3).
A number of training procedures with different ANN
architectures and training parameters have been per-
formed in order to find the best settings. Figure 4 shows
convergence of maximum relative training errors for 15
different ANN settings.
Optimal settings (listed in Table 3) were identified
by the convergence curve that reaches the lowest error at
the end of the training procedure.
Table 3: ANN training and architecture settings
Tabela 3: Nastavitve u~enja in arhitekture umetne nevronske mre`e
Training parameters
Learning rate 0.4
Momentum 0.6
Alpha value 1.0
Architecture
Neurons in input layer 34
Neurons in 1st hidden layer 25
Neurons in output layer 5
T. KODELJA et al.: TOPMOST STEEL PRODUCTION DESIGN BASED ON THROUGH PROCESS MODELLING ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 269–274 271
Figure 4: Maximal relative training error convergence for the best 15
trained ANNs
Slika 4: Najve~ja relativna napaka konvergence napake u~enja za 15
najbolj{ih umetnih nevronskih mre`
Figure 3: Training and verification procedure for model construction
Slika 3: Proces u~enja in verifikacije pri izdelavi modela
In order to build the final model, we trained the ANN
with optimal architecture and training parameters from
Table 3. The maximum number of epochs was set to 105.
Training procedures were performed on the HP ProLiant
DL 380 G7 workstation with 2 six core 3.47 GHz Intel
Xenon X5690 processors (6*256 kB L2 and 12 MB L3
cache), with 24GB installed RAM. Trained neural net-
work which gave us the best results was trained in appro-
ximately 13 hours. A remark should be given here, that
training of ANN is indeed a cumbersome and CPU time
consuming task, typically on the same order of a compu-
tational cost of a physics based model. However, when
the ANN is trained, the use of it is typically several
orders of magnitude faster than executing the physical
model.
4 RESULTS
The training procedure results in the artificial intelli-
gence model that relates the modeled output values v to
the input parameters p:
v = v(p) (1)
In the present context, v contains mechanical pro-
perties from Table 2 and p contains process parameters
from Table 1.
The obtained model can be used for a detailed study
of response of final mechanical properties on variation of
process parameters, which gives operators a better
insight into the process and can be used as a valuable
decision support tool. This is endorsed by low computa-
tional times necessary to evaluate a single response once
the model is built, which are around 10–3 s in our case.
For illustration, we show dependence of hardness on
carbon fraction around different points in the space of
model input parameters (Figure 5). We have randomly
selected 5 sets of parameters (points in the parameter
space) from the training data. Then we varied the para-
meter of interest (in our case the carbon mass fraction),
while the other parameters remained fixed. The parame-
ter was varied from the minimum to the maximum value
attained by that parameter within the training data.
It can be seen from Figure 5 that hardness generally
increases with increasing carbon mass fraction, which is
in line with the well-established metallurgical know-
ledge. This is observed for different fixed combinations
of other parameters, while the precise form of the rela-
tion varies significantly with the values of other para-
meters of the model. Since influences of individual
parameters are highly correlated, it is important for some
purposes to study behavior over larger range of process
settings. This facilitates to obtain a deeper insight in the
process. The described approach employing ANN-based
models is ideal for such purpose due to the short calcu-
lation times and exhaustiveness of information that is
provided by such models.
In another illustrative example, we take a different
point of view. Instead of focusing on influence of indi-
vidual parameters, we try to obtain a broader picture of
the comparative influence of different parameters on the
observed outcomes. We first chose a set from the training
data sets close to the center of the interval containing the
measured data. We denote the vector of input parameters
of this set by pc = pi. We then varied one by one each
component of the vector (i.e. the particular composition
or process parameter) while the others were held fixed,
and observed how the modeled quantities change as
result of this variation. More precisely, we considered
the following function of one variable:
u t v p p p t p p
i
ij i j- j+( ) ( , , ..., , , , ..., )
, ...
=
=
c1 c2 c 1 c c1
1 N j Nv p, , ...,=1
(2)
where Np is the number of model parameters and Nv is
the number of output quantities of the model. Each ele-
ment of the parameter vector pc was varied over the
whole interval that the given parameter attained in the
provided industrial data. The variations were then calcu-
lated for each output value (denoted by index i in equa-
tion (2) and for each parameter (index j in equation (2)
and used as a measure of influence of the specific
T. KODELJA et al.: TOPMOST STEEL PRODUCTION DESIGN BASED ON THROUGH PROCESS MODELLING ...
272 Materiali in tehnologije / Materials and technology 48 (2014) 2, 269–274
Figure 6: Influence of process parameters on changes in elongation
(A), ordered from the most influential one on the left, to the least
influential one on the right
Slika 6: Vplivi procesnih parametrov na spremembe v elongaciji (A),
urejeni od najbolj vplivnih na levi do najmanj vplivnih na desni
Figure 5: Steel hardness after rolling as a function of the carbon mass
fraction, calculated by the ANN model on five training sets
Slika 5: Trdota jekla po valjanju kot funkcija masnega dele`a ogljika,
izra~unana z umetno nevronsko mre`o na petih u~nih mno`icah
parameter. The results are shown in Figures 6 to 10 and
are for each parameter calculated by:
( )Δu t u t u tij ij ij( ) max ( ) min ( )= − (3)
where max ( )u tij and min ( )u tij represents maximum and
minimum influence of j-th parameter on i-th output
value.
Table 4 shows 3 most influential parameters for each
material property. From the available parameters that we
use for training the ANN (Figures 6 to 10), different
elements of the composition of the material are the most
influential for all five properties. Process parameters do
not have major influence. The most important parame-
ters obtained from the present ANN response are tempe-
rature of the liquid steel (Tcast) for elongation and
hardness after rolling, temperature difference in the
mould (Dtmould) for tensile strength, cooling water tem-
perature in zone 1 (TZone1) for yield stress and cooling
water flow rate in first spray system (Qsistem1) for
necking. Obviously, the response of the model is not
entirely expected. This indicates that the represented
methodology should be used with care and finally judged
by engineering expert knowledge. It is however true, that
in the present model, several important process parame-
ters are missing due to the lack of data acquisition in the
plant (particularly for rolling), since a new rolling mill
has been installed recently.
Table 4: The 3 most influential parameters for each mechanical pro-
perty
Tabela 4: Trije najbolj vplivni parametri za posami~no mehansko last-
nost
Elonga-
tion (A)
Tensile
strength
(Rm)
Yield
stress
(Rp)
Hardness
after
rolling
(HB)
Necking
(Z)
1 Ni C C Ti V
2 Al Cr Ni C Si
3 Ti
Delta
temperature
in the mould
Mn Ni C
5 CONCLUSIONS
ANN have been used to model a complete production
path in a steelwork. The developed methodology is
essentially a black box modeling approach. Outcomes of
the process can be predicted for arbitrary combination of
process parameters without directly considering the phy-
sical background of the modeled process, but are instead
relying on information about previous realizations of the
process. As an example, a model of production line in
the [tore Steel company was studied, reduced to 34
influential process parameters and with 5 observed pro-
perties of the final product. Several combinations of
models that will include even less influential parameters
will be studied in the future.
A significant advantage of the approach, as compared
to the physics based numerical models, is much lower
complexity of the model. There is no need to calibrate
the model in order to compensate for physical simplifi-
cations and inaccurate knowledge of model constants,
since the model is based on the realistic data gained from
the actual process. Once the model is built, evaluation
times are extremely short, in the order of a millisecond,
T. KODELJA et al.: TOPMOST STEEL PRODUCTION DESIGN BASED ON THROUGH PROCESS MODELLING ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 269–274 273
Figure 10: Influence of process parameters on changes in necking (Z)
Slika 10: Vplivi procesnih parametrov na spremembe vratu te~enja
(Z)
Figure 9: Influence of process parameters on changes in hardness
after rolling (HB)
Slika 9: Vplivi procesnih parametrov na spremembe trdote po valjanju
(HB)
Figure 8: Influence of process parameters on changes in yield stress
(Rp)
Slika 8: Vplivi procesnih parametrov na mejo te~enja (Rp)
Figure 7: Influence of process parameters on changes in tensile
strength (Rm)
Slika 7: Vplivi procesnih parametrov na natezno trdnost (Rm)
compared to several hours or even days that would be
necessary for state-of-the-art physics based models of
the same process. This represents a great advantage in
tasks where large number of evaluations are required,
such as automatic optimization of process parameters or
detailed parametric studies.22,23 This kind of modeling
has therefore a great potential to enable better insight
and understanding of industrial processing, as well as to
serve as a powerful decision support tool. This potential
was indicated in the present paper by clearly presenting
influence on individual process parameters on the out-
comes.
The drawback of the approach is its dependence on
reliable and abundant data that is sometimes hard to
obtain. Great attention must be paid to estimation of
accuracy of the model in imperfect conditions with
regard to the available data.9 This will remain the main
focus of future research, where influence of various fac-
tors on model accuracy will be studied which will even-
tually lead to procedures for reliable prediction of error
bounds, which is crucial for industrial use. This will
incorporate arrangements where controlled acquisition of
training data is possible, e.g. by using physics based
models. In this context, a large portion of work is devo-
ted to building a flexible, modular and scalable software
base to support such work. Finally, it should be noted,
that the presented methodology stimulated more careful
and complete data acquisition of the process parameters
in [tore Steel company, needed for continuation of the
present work and for better process repeatability as such.
Acknowledgment
The Centre of Excellence for Biosensors, Instrumen-
tation and Process Control (COBIK) is an operation
financed by the European Union, European Regional
Development Fund and Republic of Slovenia, Ministry
of Education, Science Culture and Sport. The financial
support of COBIK, Slovenian Research Agency and
[tore Steel company in the framework of the research
program P2-0379 and the project L2-3651 is kindly
acknowledged.
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T. KODELJA et al.: TOPMOST STEEL PRODUCTION DESIGN BASED ON THROUGH PROCESS MODELLING ...
274 Materiali in tehnologije / Materials and technology 48 (2014) 2, 269–274
A. STAMBOLI], M. MARIN[EK: THE PREPARATION OF MAGNETIC NANOPARTICLES BASED ON COBALT ...
THE PREPARATION OF MAGNETIC NANOPARTICLES
BASED ON COBALT FERRITE OR MAGNETITE
PRIPRAVA MAGNETNIH NANODELCEV NA OSNOVI
KOBALTOVEGA FERITA ALI MAGNETITA
Ale{ Stamboli}1, Marjan Marin{ek2
1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2Faculty of chemistry and chemical technology, A{ker~eva cesta 5, 1000 Ljubljana, Slovenia
ales.stambolic@imt.si
Prejem rokopisa – received: 2013-12-19; sprejem za objavo – accepted for publication: 2014-01-02
Cobalt ferrite and magnetite nanoparticles are superparamagnetic materials. Stable suspensions of superparamagnetic
nanoparticles are magnetic fluids. Stable magnetic fluids are prepared by the addition of surfactants to supporting polar or
nonpolar media. We have studied the preparation of magnetic nanoparticles based on cobalt ferrite or magnetite, and stabilized
the obtained product in an aqueous suspension. We found the optimal co-precipitation time for cobalt ferrite at 90 °C to be 2 h.
Under such conditions the product exhibited the lowest amount of amorphous phase (46.2 %), average particle diameter 10.2 nm
and a relatively high proportion of agglomerated particles (30 %). The product is therefore destabilized in an aqueous solution
of polyvinylpyrrolidone K. The optimal co-precipitation time for magnetite at 90 °C is 30 min. After 30 min the product has no
amorphous phase, average particle diameter is 6.8 nm and the proportion of agglomerated particles is low (3 %). The product
is poorly stabilized because the proportion of agglomerated particles increases when the magnetite is re-dispersed in an aqueous
solution of polyvinylpyrrolidone K. Magnetite prepared with the Massart method has a small amount of amorphous phase
(4 %), average particle diameter is 6 nm and a low proportion of agglomerated particles (4.5 %) and is therefore well
stabilized in an aqueous solution of polyvinylpyrrolidone K.
Keywords: cobalt ferrite, magnetite, superparamagnetic nanoparticles, co-precipitation, stabilization of magnetic fluids
Nanodelci kobaltovega ferita in magnetita imajo superparamagnetne lastnosti. Stabilne suspenzije superparamagnetnih
nanodelcev imenujemo magnetne teko~ine. Stabilne magnetne teko~ine pripravimo z dodatkom surfaktantov v nosilni polarni
ali nepolarni medij. Sintetizirali smo magnetne nanodelce kobaltovega ferita in magnetita ter dobljeni produkt stabilizirali v
vodnem mediju. Dolo~en je bil optimalen ~as soobarjanja kobaltovega ferita pri 90 °C, in sicer 2 h. V tem ~asu ima produkt
najni`jo vsebnost amorfne faze (46,2 %), povpre~no velikost delca 10,2 nm in dele` ne`elenih aglomeratov pa je 30 %. Ko se
delce kobaltovega ferita redispergira v vodi z dodanim surfaktantom, nastane destabilizirana suspenzija, katere vzrok je visok
dele` aglomeratov. Optimalen ~as koprecipitacije magnetita pri 90 °C je 30 min. Takrat produkt ne vsebuje amorfne faze,
povpre~na velikost delca je 6,8 nm, dele` aglomeriranih delcev je 3 %. Produkt se nato redispergira v vodnem mediju z
dodanim surfaktantom, kjer je magnetit slabo stabiliziran zaradi pove~anja dele`a aglomeratov med redisperzijo. Magnetit,
pridobljen z metodo po Massartu, vsebuje 4 % amorfne faze, ima povpre~no velikost delca 6 nm, dele` aglomeratov 4,5 %
in je zato odli~no stabiliziran v vodni raztopini polivinilpirolidona K.
Klju~ne besede: kobaltov ferit, magnetit, superparamagnetni nanodelci, koprecipitacija, stabilizacija magnetnih teko~in
1 INTRODUCTION
Magnetic fluids are stable dispersions of magnetic
nanoparticles in an organic or aqueous medium. Magne-
tic nanoparticles are solid-state particles that are affected
by magnetic fields. The term "nano" indicates particles
that are in the size range below 100 nm. Magnetic nano-
particles occur in the form of ferrites with the general
formula MFe2O4, where M is a divalent metal cation
(nickel, cobalt, manganese, zinc or iron). Ferrites crystal-
lize in the spinel crystal structure.1–3
The most important characteristic of magnetic nano-
particles is superparamagnetism. Superparamagnetism
relates to nanoparticles whose magnetic moments are
oriented irregularly and do not show any magnetic
properties in the absence of an applied magnetic field.
However, even when these nanoparticles are exposed to a
small external magnetic field, dipole moments are
formed within the particles. The dipole moments are
directed in accordance with the source of the magnetic
field.4
There are several techniques for the preparation of
magnetic nanoparticles. The most common are: co-preci-
pitation, sol-gel synthesis, hydrothermal synthesis and
synthesis in microemulsions. When magnetic nanoparti-
cles are obtained it is better if the resulting particles are
crystalline, have a narrow size distribution and have the
same shape.1,3
The developing product must be stabilized so it does
not agglomerate due to the attractive forces between the
particles. The techniques of stabilization are electro-
static, steric and electrosteric. Electrostatic stabilization
is the mechanism in which the attractive forces are due to
the Coulomb forces between the charged colloidal parti-
cles. The repulsion between the particles is achieved
with an equally charged electric double layer surround-
ing the particles. Steric stabilization is accomplished by
the adsorption of long-chained molecules called surfac-
Materiali in tehnologije / Materials and technology 48 (2014) 2, 275–280 275
UDK 621.318.1:66.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)275(2014)
tants on the particle’s surface. These surfactants then
form a coating that creates a repulsive force and sepa-
rates one particle from another particle. Electrosteric
stabilization is a combination between electrostatic and
steric stabilization. The repulsion between the particles is
achieved by the adsorption of charged polymeric mole-
cules on the nanoparticle’s surface.4,5
The applications of magnetic fluids are primarily in
engineering and medicine as seals for rotating mecha-
nical parts, as in the heat dissipation system in speakers
and as a contrast agent in imaging with the nuclear
magnetic resonance technique.6,7
2 EXPERIMENTAL
The main goal of this study was to prepare magnetic
nanoparticles based on cobalt ferrite or magnetite, and
stabilize the obtained product in an aqueous suspension.
The product was quantitatively and qualitatively ana-
lysed with a Rietveld analysis and morphologically
described using scanning electron microscopy. The
cobalt ferrite and magnetite were prepared by co-preci-
pitation at 90 °C and pH = 13. The magnetite was also
synthesized with the Massart method, where the product
is instantaneously precipitated at room temperature with
vigorous stirring. The product was stabilized with poly-
vinylpyrrolidone K in an aqueous medium.
2.1 Synthesis of cobalt ferrite and magnetite at ele-
vated temperature
A solution of Co2+ and Fe3+ ions or a solution of Fe2+
and Fe3+ ions was mixed in a laboratory reactor with a
reflux condenser using a magnetic stirrer (Figure 1).
Then, in deionized water dissolved sodium hydroxide
was added, which allows the progress of the synthesis at
pH = 13. The stirring was continued. Immediately after
the addition of sodium hydroxide the dark precipitate
was formed.
The chemical reaction for the cobalt ferrite prepa-
ration is as follows:
8 NaOH(aq) + CoSO4(aq) + Fe2(SO4)3(aq) 
CoFe2O4(s) + + 4 Na2SO4(aq) + 4 H2O (1)
The chemical reaction for the magnetite preparation
is as follows:
8 NaOH(aq) + 2 FeCl3(aq) + FeCl2(aq) 
Fe2+O · Fe2
3+O3(s) + 8 NaCl(aq) + 4 H2O (2)
Next, the oleic acid was added to the reaction mixture
and it was heated up to 90 °C. This temperature was
regulated until the end of the reaction. Following the
synthesis the pH of the precipitate was reduced from 13
to 5.5, after which the mixture was filtered and the
obtained nanoparticles were re-dispersed and stabilized
in an aqueous medium with the addition of a surfactant.
2.2 Synthesis of magnetite using the Massart method
Aqueous solutions of Fe2+ and Fe3+ ions were mixed
in a beaker. During vigorous stirring ammonia was
added. This leads to the formation of a black, magnetic
precipitate. After the reaction, when all the ingredients
are well mixed, the mixture was filtered and the nano-
particles were re-dispersed and then stabilized in an
aqueous medium with the addition of a surfactant.
The chemical reaction of the magnetite prepared with
the Massart method can be described by Eqs. (3) and (4):
NH3(l) + H2O  NH4
+
(aq) + OH
–
(aq) (3)
8 NH4
+
(aq) + 8 OH
–
(aq) + FeCl2(aq) + 2 FeCl3(aq) 
FeO · Fe2O3(s) + 8 NH4Cl(aq) + 4 H2O (4)
2.3 Stabilization of the magnetic fluids
The first level of stabilization was achieved with the
addition of oleic acid during the development of the syn-
thesis. Oleic acid is an unsaturated fatty acid (between
the carbons atoms there is at least one double bond) with
the molecular formula C18H34O2. The oleic acid sur-
rounds the nanoparticles and prevents their further
growth and agglomeration during the synthesis.
The second level of stabilization was realized after
the purification of the product. The resulting cake of
particles was necessary to re-disperse and stabilize in an
aqueous medium. Steric stabilization with polyvinyl-
pyrrolidone [(C6H9NO)n] label K was used. Its molecular
weight is 30000 g/mol. Polyvinylpyrrolidone is freely
soluble in polar solvents, where it creates films that can
serve as a coating for the particles.
A. STAMBOLI], M. MARIN[EK: THE PREPARATION OF MAGNETIC NANOPARTICLES BASED ON COBALT ...
276 Materiali in tehnologije / Materials and technology 48 (2014) 2, 275–280
Figure 1: Laboratory reactor used for co-precipitation at elevated tem-
perature
Slika 1: Laboratorijski reaktor za koprecipitacijo pri povi{ani tempe-
raturi
2.4 X-ray powder diffraction and Rietveld analysis
X-ray diffraction shows whether the desired product
(cobalt ferrite or magnetite) is formed. A subsequent
Rietveld analysis determines the amount of amorphous
phase in a sample.
X-ray powder analysis provides information about
the chemical composition, the degree of crystallinity and
the size or symmetry of the crystal unit cell. With this
method an X-ray beam is irradiated at the sample’s sur-
face at an angle . During the operation the sample is
synchronously rotated by half the speed of the detector
so that the angle between the source of radiation and the
sample is always equal to the angle between the sample
and the detector. At every  the X-rays are apparently
deducted from the samples surface. If the phase shift
between the various parallel X-rays is equivalent to a
multiple of the wavelength (peak of one wave coexists
with the top of the second wave) then the X-rays are
strengthened, otherwise they override.8,9
The samples were recorded on X-ray diffractometer
PANalytical X’Pert PRO using Cu-K1 radiation with a
wavelength of 1.5406 × 10–10 m.
Using the Rietveld method the sample was quantita-
tively analysed. The relative proportions of the individual
crystal phases within the sample are determined.
Measured and calculated powder patterns are compared
using the Rietveld analysis. Also, the contribution of the
individual phases in multiphase samples can be detected,
including amorphous components or an unknown phase
(with the addition of a crystalline or amorphous standard
with a known mass fraction).10
2.5 Scanning electron microscopy (SEM)
The microstructure of the samples was characterized
with scanning electron microscopy (SEM). The micro-
structure of the materials is related to their properties
(mechanical, electrical, optical and magnetic) and thus
their usefulness and purpose. SEM is used to observe the
surface of non-volatile solid samples. This technique
allows the interaction of electrons with the surface atoms
and gives information about the shape, size, grain size
distribution, elemental composition and the proportion of
each element in the sample.9,11
For a nanostructure analysis a Zeiss Ultra Plus
(FE-SEM) scanning electron microscope equipped with
an EDS detector SDD Oxford was used. From SE ima-
ges the morphological characteristics of the prepared
products, especially the particle size, were defined.
A. STAMBOLI], M. MARIN[EK: THE PREPARATION OF MAGNETIC NANOPARTICLES BASED ON COBALT ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 275–280 277
Figure 2: Dependence of the proportion of amorphous phase vs. time
for the synthesis of: a) cobalt ferrite and b) magnetite
Slika 2: Odvisnost dele`a amorfne faze od ~asa sinteze za: a) kobaltov
ferit in b) magnetit
Figure 3: SE images of cobalt ferrite after: a) 0.5 h, b) 2 h and c) 4 h
of co-precipitation
Slika 3: SE-posnetki kobaltovega ferita po: a) 0,5 h, b) 2 h in c) 4 h
koprecipitacije
3 RESULTS AND DISCUSSION
When preparing cobalt ferrite and magnetite at ele-
vated temperature, we were interested in the proportion
of amorphous phase in the sample during the synthesis in
order to determine the optimal synthesis time. For this
purpose we used the Rietveld method.
Figure 2a shows that the proportion of amorphous
phase in the cobalt ferrite is in a high range of about
50 % for all 4 h of synthesis. The reason for such a high
proportion may be an incomplete reaction, old reactants
or very fast precipitation. If the degree of supersaturation
is high, then the precipitation is finished extremely
quickly. Therefore, the particles do not have sufficient
time to organize into a proper structure, which leads to
the formation of an amorphous phase. This mechanism is
quite possible as the average particle size varies very
little after 30 min during the reaction. This means that
the majority of the precipitation has come to an end
much before 30 min. After 30 min of reaction the
amount of amorphous phase is 51.61 % and then it is
slightly decreasing for another hour and a half, when it
reaches a value of 46.19 %. After 2 hours of synthesis
the amount of amorphous phase varies. The minimum
value is reached after 3.5 h, i.e., 46.13 %, but due to
possible method errors, very small values deviation and
economic reasons (long-term progress of the synthesis
results in higher costs) it is enough to conduct the
synthesis of the cobalt ferrite for 2 h or at least one hour.
Because cobalt ferrite has a relatively high content of
amorphous phase, the material has poor superparamag-
netic properties as the particles react quite slowly on the
magnet.
Figure 2b reveals a slight increase in the proportion
of amorphous phase with the time of synthesis for the
magnetite. Since after 30 min of reaction there is no
amorphous phase in the magnetite, we have further ope-
rated with synthesis for just as long. Due to the lower
levels of amorphous phase the superparamagnetic pro-
perties of the material are much better.
A Rietveld analysis was also performed for the mag-
netite synthesized using the Massart method. The pro-
portion of amorphous phase was 4.3 %.
A. STAMBOLI], M. MARIN[EK: THE PREPARATION OF MAGNETIC NANOPARTICLES BASED ON COBALT ...
278 Materiali in tehnologije / Materials and technology 48 (2014) 2, 275–280
Figure 5: Average diameter changes of individual particles, agglome-
rated particles and overall particle for the synthesis of: a) cobalt ferrite
and b) magnetite
Slika 5: Spreminjanje nadomestnega povpre~nega premera posamez-
nih delcev, aglomeriranih delcev in vseh delcev s ~asom sinteze za: a)
kobaltov ferit in b) magnetit
Figure 4: SE images of magnetite after: a) 0.5 h, b) 2 h and c) 4 h of
co-precipitation
Slika 4: SE-posnetki magnetita po: a) 0,5 h, b) 2 h in c) 4 h kopre-
cipitacije
SE images (Figures 3 and 4) were used to determine
the morphological characteristics of the prepared pro-
ducts. The developing cobalt ferrite particles have predo-
minantly a spherical shape, while the magnetite particles
have a cubic shape.
The average size of the cobalt ferrite’s individual
particles after the first hour is approximately 7.3 nm,
which then increases to about 8 nm. The same trend is
also visible for the overall particle (individual and agglo-
merated), where the average size is in the range 9.46 nm
to 10.37 nm (Figure 5a). The proportion of cobalt ferrite
agglomerated particles is about 30 % for all 4 h of syn-
thesis. A slight increase in the size of the individual
particles is due to the co-precipitation mechanism, which
in this case is not so definite. After 0.5 h the proportion
of smallest particles (less than 4 nm) is about 20 % and
after 1 h it dropped to 10 %. The conclusion from the
morphological analysis of the cobalt ferrite particles is
that in a time interval of 0.5 h to 1 h the formation of a
new crystal nucleus is significantly reduced, but the
particles are still growing. After 1 h of synthesis there
are no noticeable changes so the crystal growth must
have decelerated. The same trend was also observed for
the agglomerated particles.
The size of the individual, agglomerated and overall
particles of magnetite is slowly increasing with the syn-
thesis time. The average individual particle grows from
6.84 nm to 8.97 nm and the overall average particle
(individual + agglomerated) grows from 6.84 nm to
10.49 nm (Figure 5b). The overall average particle size
increases much more than the individual particle because
the proportion of agglomerated particles increases
considerably with the synthesis time (from 2.86 % to
21.36 %). When the magnetite was precipitated, typical
mechanisms of co-precipitation were observed. After 30
min there are no agglomerated particles in the system.
With the extension in the synthesis time the amount of
the smallest particles decreases, which implies a decline
in the formation of the crystallization nucleus. However,
the particles are growing more noticeably and conse-
quently agglomerating with an increase of the synthesis
time. The observed behaviour of the Fe3O4 system during
the synthesis where the magnetite particles are growing
slowly, while slowly increasing the proportion of amor-
phous phase with synthesis time, indirectly indicates the
mechanism of formation of the solid phase magnetite in
two stages. The first stage is relatively fast and results in
magnetite nanoparticles with a number of defects higher
than equilibrium. The second stage is substantially slo-
wer. On the initially precipitated Fe3O4 surface new,
smaller magnetite grains with an equilibrium number of
defects should re-precipitate.
The magnetite synthesized by the Massart method
has a low proportion of agglomerated particles (4.6 %)
so the possibility of their stabilization is extensive. The
average individual particle size is 6.1 nm, and the overall
particle size is 6.4 nm.
The last stage of the research was to stabilize the
magnetic fluid with PVP K. The addition of a surfactant
should take care of the steric stabilization in an aqueous
medium.
A range of the mass fractions w = 0.001 % to 1 %
PVP K was added in the polar aqueous medium and 1 %
in decane. Figure 6a shows that cobalt ferrite is not sta-
bilized in an aqueous medium with PVP K, irrespective
of the quantity of added PVP K. Poorly stabilized,
whereas the precipitate is settling, is the cobalt ferrite in
a non-polar decane on the far right of Figure 6a. Despite
everything, the settling of the particles was noticeably
slower in the suspensions with a higher mass fraction of
PVP K and faster in those with a lower mass fraction of
PVP K.
In Figure 6b the suspensions of magnetite are shown.
The stability of the suspensions is evidently improved
compared to the cobalt ferrite, but the suspensions are
still far from ideal. Again, 0.001 % to 1 % of PVP K was
added and once again we observed an improved stability
for a larger addition of PVP K.
A. STAMBOLI], M. MARIN[EK: THE PREPARATION OF MAGNETIC NANOPARTICLES BASED ON COBALT ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 275–280 279
Figure 6: Magnetic suspensions with different mass fractions (w/%)
addition of PVP K for: a) cobalt ferrite, b) magnetite obtained at an
elevated temperature and c) magnetite synthesized with the Massart
method
Slika 6: Magnetne suspenzije z razli~nimi masnimi dele`i (w/%) PVP
K-ja: a) kobaltov ferit, b) magnetit, pridobljen pri povi{ani temperaturi
in c) magnetit, sintetiziran z Massartovo metodo
In the case of the magnetite synthesized using the
Massart method stable suspensions were prepared. In
stable suspensions the particles do not settle and are
distributed throughout the volume (Figure 6c). All the
reconstituted suspensions are stable, regardless of the
amount of added PVP K.
The stability of the suspensions can be related to the
particle size and the proportion of agglomerated parti-
cles. Cobalt ferrite has an average particle size of 7.45
nm, but it also has 26 % of agglomerated particles. A
high proportion of agglomerated particles implies a
larger number of particles with greater mass that settle
smaller particles below and thus destabilize the suspen-
sion. Magnetite has an average particle size of 8.17 nm,
but only about 17 % of agglomerated particles. The num-
ber of magnetite particles with higher mass is much
lower, so the particles settle slowly, but eventually they
all settle. Magnetite synthesized with the Massart
method has a particle size of 6.37 nm and only 4.5 % of
agglomerated particles. These suspensions are very
stable due to the particles’ small size and because the
contact between the particles was prevented by the
addition of the surfactant.
4 CONCLUSIONS
The synthesis of cobalt ferrite operating for 4 h at 90
°C ensures a relatively constant average size of overall
particles equal to about 10 nm during the synthesis. The
proportion of amorphous phase in the samples was
approximately constant throughout synthesis, i.e., high
 50 %. The proportion of agglomerated particles within
the suspension was around 30 %. An optimal synthesis
time of 2 h was determined. After 2 h of synthesis the
average particle size was 10.22 nm, the sample has 46.19 %
of amorphous phase and the proportion of agglomerated
particles was 30.44 %.
The synthesis of magnetite, which lasted for 4 h at 90
°C, exhibited a growth of the average particle size, an
increase in the amount of amorphous phase and the pro-
portion of agglomerated particles during the synthesis.
The optimal time for the magnetite synthesis was 30 min
when the particle size is 6.84 nm, the amount of the
amorphous phase is 0 % and the proportion of agglome-
rated particles is 3 %.
Magnetic suspensions are stable when the particles
are not agglomerated, dispersed all over the liquid vo-
lume and do not settle. For the stabilization of the mag-
netic fluids the surfactant polyvinylpyrrolidone K (M 
30000 g/mol) was used. PVP K is freely soluble in water.
When the nanoparticles are re-dispersed in liquid + PVP
K the magnetite particles synthesized using the Massart
method are more stable than the magnetite particles
synthesized at 90 °C, which are more stable than the
cobalt ferrite particles. Cobalt ferrite has an average
particle size of 7.45 nm, but has 26 % of agglomerated
particles. A high proportion of agglomerated particles
implies a larger number of particles with greater mass
that settle smaller particles below and thus destabilize
the suspension. Magnetite has an average particle size of
8.17 nm, but only about 17 % of agglomerated particles.
Because the number of particles with a higher mass is
much smaller, the particles settle moderately, but even-
tually they all settle. Magnetite synthesized with the
Massart method has a particle size of 6.37 nm and only
4.5 % of agglomerated particles. These suspensions are
very stable due to the particles’ small size and because
the contact between the particles has been prevented by
the addition of surfactant.
5 REFERENCES
1 V. K. Varadan, K. Chen, X. J. Linfeng, Nanomedicine: Design and
Applications of magnetic nanomaterials, nanosensors and nanosys-
tems, Wiley, Chichester 2009, 38–39
2 M. Rem{kar, Nanodelci in nanovarnost, Ministrstvo za zdravje, Urad
RS za kemikalije, Ljubljana 2009, 14–17
3 A. Goldman, Modern Ferrite Technology, Second edition, Springer,
Boston 2006, 51–58, 172–174
4 R. E. Rosensweig, Ferrohydrodynamics, Dover, New York 1997, 7,
32–38, 46, 55–63
5 M. N. Rahaman, Ceramic processing and sintering, Second edition,
Marcel Dekker, New York 2003, 190–191
6 D. Makovec, Magnetne teko~ine in njihova uporaba v tehniki, @iv-
ljenje in tehnika, 58 (2007) 12, 60–64
7 D. Makovec, Uporaba magnetnih nanodelcev v medicini, @ivljenje in
tehnika, 60 (2009) 2, 39–42
8 W. Clegg et al., Crystal stucture analysis: Principles and practice, Se-
cond edition, Oxford University Press, New York 2009, 251–253
9 F. Zupani~, I. An`el, Gradiva, Fakulteta za strojni{tvo, Maribor 2007,
53–59
10 http://www.ki.si/fileadmin/user_upload/datoteke-L09/nzl/Kristalo-
grafija-NZL-POGL4.pdf 29. 6.2012
11 M. Marin{ek, Electron microscopy: Application of scanning electron
microscope for observation and analysis of the solid samples surface,
Literature on the subject of Materials, Faculty of chemistry and che-
mical technology, Ljubljana 1999
A. STAMBOLI], M. MARIN[EK: THE PREPARATION OF MAGNETIC NANOPARTICLES BASED ON COBALT ...
280 Materiali in tehnologije / Materials and technology 48 (2014) 2, 275–280
K. MRAMOR et al.: SIMULATION OF CONTINUOUS CASTING OF STEEL UNDER THE INFLUENCE ...
SIMULATION OF CONTINUOUS CASTING OF STEEL
UNDER THE INFLUENCE OF MAGNETIC FIELD USING
THE LOCAL-RADIAL BASIS-FUNCTION COLLOCATION
METHOD
SIMULACIJA KONTINUIRNEGA ULIVANJA JEKLA POD
VPLIVOM MAGNETNEGA POLJA NA PODLAGI METODE
KOLOKACIJE Z RADIALNIMI BAZNIMI FUNKCIJAMI
Katarina Mramor1, Robert Vertnik2,3, Bo`idar [arler1,3,4
1CO BIK, Tovarni{ka c. 26, 5270 Ajdov{~ina, Slovenia
2[tore Steel, @elezarska c. 3, 3220 [tore, Slovenia
3University of Nova Gorica, Vipavska 13, 5000 Nova Gorica, Slovenia
4Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
katarina.mramor@cobik.si
Prejem rokopisa – received: 2013-12-30; sprejem za objavo – accepted for publication: 2014-01-13
The initial results obtained with the local-radial basis-function collocation method (LRBFCM) for a two-dimensional (2D),
simplified model of continuous casting of steel with an externally applied magnetic field are presented. The multiphysics model
is composed of turbulent-solidification equations (mass, momentum, energy, turbulent kinetic energy, turbulent dissipation rate)
and Maxwell’s equations that are numerically solved for the non-uniform node arrangement. The numerical procedure is
structured using the explicit time stepping and local collocation with multiquadric radial basis functions (MQ RBF) on the
overlapping five-node subdomains. The pressure-velocity coupling follows the fractional-step method (FSM) and the convection
is treated with adaptive upwinding.
The novel LRBFCM has been already verified in several benchmark test cases, such as the natural convection in a cavity with a
magnetic field, the lid-driven cavity, and the flow over the backward-facing step with a transverse magnetic field.
Keywords: continuous casting of steel, turbulent flow, solidification, magnetohydrodynamics
Namen ~lanka je predstavitev prvih rezultatov s poenostavljenim 2D-modelom za kontinuirano ulivanje jekel pri zunanjem
magnetnem polju, izra~unanih z metodo kolokacije z radialnimi baznimi funkcijami. Numeri~ni model zdru`uje Navier-
Stokesove in Maxwellove ena~be, ki jih re{ujemo na neenakomerni porazdelitvi to~k. V numeri~nem postopku je uporabljena
eksplicitna ~asovna shema na petto~kovnih poddomenah, na katerih so uporabljene multikvadri~ne radialne bazne funkcije.
Sklopitev tlaka in hitrosti se re{uje z metodo delnih korakov.
Omenjeno metodo smo poprej verificirali za izra~un razli~nih preizkusov, kot so naravna konvekcija v kotanji z magnetnim
poljem, kotanja z vsiljenim tokom ter tok preko stopnice v kanalu z magnetnim poljem. Dobljeni rezultati ka`ejo dobro
ujemanje z drugimi numeri~nimi postopki, kot je npr. metoda kon~nih volumnov.
Klju~ne besede: kontinuirano ulivanje jekla, turbulenten tok, strjevanje, magnetohidrodinamika
1 INTRODUCTION
The production of continuously cast steel1 has greatly
expanded in recent years. Continuous casting of billets,
blooms and slabs is the most common process of steel
production.2 The continuously growing demand for cast
steel has fueled the need to produce the steel of even
better quality. Although the process of continuous cast-
ing of steel is cost efficient, exhibiting a high yield and
good quality of the products, it can be further improved
by introducing the electromagnetic (EM) field into it.
The EM force, which is a result of an applied magnetic
field, affects the velocity and temperature fields. By
adjusting the magnetic field, the amount of defects in the
material can be significantly reduced.
The magnitudes of the velocity, the temperature and
magnetic fields are all crucial to the final quality of a
product. As all these quantities are difficult or impossible
to measure, numerical models help us to better under-
stand and further improve the process. A number of
different numerical models3 have so far been used in the
simulations of the problem. They include the finite-
volume method (FVM),4–8 the finite-element method
(FEM)9 and some more advanced meshless methods like
the local-radial basis-function collocation method
(LRBFCM).10
2 GOVERNING EQUATIONS
The system of governing equations that describes the
turbulent heat transfer and the fluid flow as well as the
magnetic field in the continuous casting of steel, is based
on the Reynolds time-averaging approach to modeling a
turbulent flow11 and a mixture-continuum model, first
introduced by Bennon and Incropera.12 The system con-
sists of five time-averaged mixture equations:
∇⋅ =v 0 (1)
Materiali in tehnologije / Materials and technology 48 (2014) 2, 281–288 281
UDK 621.74.047:519.61/.64 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(2)281(2014)
[ ][ ]    ! " # #

#
∂
∂
v
vv v v
t
p
k
f
T+ ∇ ∇ ∇ + ∇ + ∇ −
− ∇ −
−
( ) ( )
( )
L t
L L2
3
1 2
Cf
g T Tref
L
s T m3 ( ) ( )v v F− + − +
(2)
      "


∂
∂
h
t
h T h f h
f h f
v
+ ∇ ∇ ∇ ∇ − −
− + ∇
v v v
v
( ) (
)
S S S
L L L L
L t
t
L∇
⎛
⎝
⎜
⎞
⎠
⎟h
(3)
     #
#

  #
∂
∂
k
t
k k P G
D
+ ∇ ∇ +
⎛
⎝
⎜
⎞
⎠
⎟∇
⎡
⎣⎢
⎤
⎦⎥
+ + −
− + −
v L
t
k
k k
L
L
L
( )1 2
3
− f
Cf
k
(4)


    #
#

 
!#

∂
∂t
E
f
C
+ ∇ ∇ +
⎛
⎝
⎜
⎞
⎠
⎟∇
⎡
⎣⎢
⎤
⎦⎥
+ −
−
v L
t
L
L( )1
2
[ ]
f
c f P c G c f
kk kL
3 1 1 3 2 2" 

  ( )+ −
(5)
where v is the velocity of the mixture,  = S = L is the
density of the mixture, assumed to be constant; t stands
for the time and p for the pressure; μt is the turbulent
viscosity and μL is the dynamic viscosity; k represents
the turbulent kinetic energy and K is the permeability of
the porous matrix. T, g, T, and Tref represent the
thermal-expansion coefficient, the gravitational accele-
ration, the temperature and the reference temperature,
respectively. Fm stands for the Lorentz force, h for the
enthalpy and  for the thermal conductivity. fS, fL, vS, vL,
hS, and hL represent the solid-volume fraction, the
liquid-volume fraction, the velocity of the solid phase,
the velocity of the liquid phase, the enthalpy of the solid
phase and the enthalpy of the liquid phase, while vt is
the turbulent kinematic viscosity. Symbols  and C stand
for the dissipation rate and the morphology constant of
the porous media. Symbols t, k, , c1, f1, c2 and f2
are the closure coefficients of the turbulence model. Pk,
Gk, D and E are the shear production of the turbulent
kinetic energy, the generation of turbulence due to the
buoyancy force, the source term in the k equation and
the source term in the  equation, respectively. The
closure relations defined by Abe, Kondoh and Nagano13
are used in the present work. A detailed description of
the closure coefficients, source terms and damping
functions are given in the paper by [arler et al14. The
Lorentz force is defined as:
Fm = j × B (6)
where j and B are the current density and the magne-
tic-flux density. Maxwell’s equations are used to calcu-
late the current density:
j = (–$% +v × B) (7)
where  is the fluid electric conductivity, % is the fluid
electric potential and B is the externally applied mag-
netic field. The induced magnetic field is assumed negli-
gible in comparison with the applied magnetic field.
2.1 Initial and boundary conditions
The governing equations are strongly coupled. For
the steady solution of a problem it is, therefore, very
important how the initial conditions are chosen in order
to minimize the required iterations to reach the steady
state. Five different boundaries are chosen: the inlet, free
surface, wall, outlet and symmetry of the present model.
A detailed description of the initial and boundary con-
ditions are given in the article by [arler at al14. The
model of the domain is presented in Figure 1 and the
computational domain is depicted in Figure 2.
3 SOLUTION PROCEDURE
The solution of the governing equations is obtained
with the LRBFCM employing the explicit time stepping.
The fractional-step method (FSM)15 is used to solve the
pressure-velocity coupling and an upwinding scheme is
used to stabilize the highly convective situation.16
The solution procedure begins by calculating the
initial Lorentz force (equation (6)). Afterwards, the inter-
mediate velocity is calculated, without a pressure gra-
dient. The pressure is then calculated from the Poisson
equation14 by assembling and solving a sparse matrix.17
The calculated pressure gradient is afterwards used to
correct the intermediate velocities. After the solution of
the velocity field, the equations for the turbulent kinetic
energy and dissipation rate are solved. This is followed
K. MRAMOR et al.: SIMULATION OF CONTINUOUS CASTING OF STEEL UNDER THE INFLUENCE ...
282 Materiali in tehnologije / Materials and technology 48 (2014) 2, 281–288
Figure 1: Simplified 2D model of continuous casting of steel
Slika 1: Poenostavljeni 2D-model kontinuirnega ulivanja jekla
by the solution of the enthalpy equation. The tempera-
ture is calculated from the enthalpy, using the tempera-
ture-enthalpy constitutive relation.17,18 Finally, the turbu-
lent viscosity, velocity, temperature, turbulent kinetic
energy and dissipation rate are updated and the solution
is ready for the next time step.
The LRBFCM is structured in the following way:
Approximation function  is represented on each of the
subdomains as a linear combination of the radial basis
functions (RBFs) as:
q l n l i l n
i
M
l i( ) ( )p p=
=
∑  
1
(8)
where li, li, and M represent the RBF shape functions,
centred in points lpn, the expansion coefficients, and the
number of shape functions, respectively. The most
commonly used RBF is the multiquadric RBF:19,20
l i l ir c ( ) ( )p p= +
2 2 (9)
where c stands for the dimensionless shape parameter,
set to 32 in all the calculations, and the distance bet-
ween the nodes:
l i
i
l i
i
l i
r
x x
x
y y
y
2
2 2
( )
max max
p =
−⎛
⎝
⎜
⎞
⎠
⎟ +
−⎛
⎝
⎜
⎞
⎠
⎟ (10)
is scaled with lximax and lyimax, the scaling parameters of
subdomain l in the x and y directions, respectively (Fig-
ure 3).
A subdomain is formed around each of the calcula-
tion points consisting of the lM – 1 nodes nearest to node
lpn. For the purpose of this work, five-node overlapping
subdomains are used. By considering the collocation
condition of:
l l n i l n ( ) ( , )p = (11)
a linear system of equations is obtained. To solve the
partial differential equations (PDEs), the first and se-
cond derivatives of function l(p) have to be calculated:
∂
∂
∂
∂
j
j l
j
j l i
i
M
l i&

&
 ( ) ( )p p=
=
∑
1
(12)
where index j = 1,2 is used to denote the order of the
derivative and & = x,y. A detailed explanation of the
solution procedure is given in the paper by Vertnik et
al.17,18 The schematics of the discretization scheme is
shown in Figure 3.
4 NUMERICAL IMPLEMENTATION
The sparse-pressure matrices, used for the solution of
the pressure and the flux-density Poisson equations, are
solved by applying the Pardiso routine and Intel Math
Kernel Library 11. The OpenMP Library is used for the
parallelization. The post processing is performed in
PGPlot, Gnuplot 4.4 and Octave 3.6.1. The results were
calculated on an HP Proliant DL380 G/7 server running
on 64 bit MS Windows.
5 NUMERICAL EXAMPLES
The numerical procedure has so far been verified on
the following benchmark test cases: the lid-driven
cavity,21 the natural convection in a cavity,22 and the
backward-facing step.23 The lid-driven test case was used
K. MRAMOR et al.: SIMULATION OF CONTINUOUS CASTING OF STEEL UNDER THE INFLUENCE ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 281–288 283
Figure 3: Discretization scheme. ', , lxiMAX and lyiMAX represent
the boundary, domain and scaling parameters in the x and y directions,
respectively. The circles represent the boundary nodes, whereas the
black dots represent the domain nodes.
Slika 3: Diskretizacijska shema. ', , lxiMAX in lyiMAX ozna~ujejo
rob, obmo~je ter skalirna parametra v smeri x in y. Robne to~ke so
ozna~ene s krogci, obmo~ne to~ke pa s pikami.
Figure 2: Scheme of the computational domain
Slika 2: Shema ra~unske domene
to verify the coupling between the mass and momentum
equations.21 In the natural-convection case, the
energy-conservation equation was added.22 For the
backward-facing step problem, the boundary conditions
for the inflow and outflow problems were tested. The
natural-convection and backward-facing-step problems
were first tested without a magnetic field and then with
an external magnetic field. The results of all of the test
cases are in good agreement with the reference results,
calculated with a commercial code or obtained from
several published sources. Since the results of our
calculations, mentioned in the above history of test
cases, are in good agreement with the reference results,
the method is subsequently applied to the simplified
problem of continuous casting of steel with a magnetic
field, as defined in the paper by [arler et al.14
5.1 Continuous-casting geometry and material proper-
ties
The simplified 2D continuous-casting model geome-
try is shown in Figure 1 and the elements of the discre-
tization are presented in Figure 2. The computational
domain coincides with a half of the longitudinal section
of the billet, taken to be 1.8 m long and 14 cm wide. The
SEN diameter is 3.5 cm, the mold height is 0.8 m and the
EM-coil height is 10 cm. The material properties of steel
are temperature and steel grade dependent. However, for
the purpose of the present simplified model, constant
values are used. The values are given in Table 1.
Table 1: Simplified material properties of steel
Tabela 1: Poenostavljene snovne lastnosti jekla
Property Value
 7200 kg/m3
 30 W/(m K)
cp 700 J/(kg K)
TS 1680 K
TL 1760 K
hm 250000 J/kg
μ 0.006 Pa s
T 1 · 10–4 K–1
C 1.6 · 108 m–2
The results of the numerical simulation are represen-
ted in the following sections of the paper.
5.2 Convergence of the method and a comparison with
the reference results for the continuous-casting
process
First the convergence of the method was tested for
the node arrangements with 20 220, 50 951, 73 940, 100
089, 131 452, 165 426 non-uniformly arranged nodes.
The convergence was tested for the velocity and the tem-
perature fields. The vertical and horizontal components
of the velocity were compared for three different vertical
cross-sections: just before the application of the mag-
netic field, just after the application of the magnetic field
and at the end of the computational domain, as can be
seen in Figure 4 (left). The temperature was verified at
two different vertical directions, the one at the center of
the domain and the one at the outside wall, as shown in
Figure 4 (right).
As can be seen in Figures 5 to 12, the smallest
appropriate node arrangement is 100 089. The results
were also compared with the results obtained with the
commercial code based on the FVM (Fluent23). The
smallest appropriate amount of the finite volumes in the
commercial code for reaching a reasonable convergence
was 169 169. The agreement of the velocity profiles with
the commercial code (the solid line denoted with F) is
similarly good for the horizontal as well as vertical
velocities. The largest differences occur at the positions
closer to the inlet. At the positions closer to the end of
the computational domain, where a fully developed flow
and a partial solidification take place, the agreement
between the results obtained with the LRBFCM and
FVM is excellent. In the early stages of the flow, slightly
larger differences can be observed. The differences
between the commercial-code results and the results
obtained with the in-house built LRBFCM are
reasonably small. The exact cause for the differences is
K. MRAMOR et al.: SIMULATION OF CONTINUOUS CASTING OF STEEL UNDER THE INFLUENCE ...
284 Materiali in tehnologije / Materials and technology 48 (2014) 2, 281–288
Figure 4: Left: the positions of the velocity-field comparison lines; hI
= –0.8 m, hII = –0.9 m, hIII = –1.8 m. Right: the positions of the
temperature-field comparison lines; vI = 0.0 m (centreline), vII = 0.07
m (surface).
Slika 4: Levo: Polo`aj linij, kjer je primerjano hitrostno polje; hI =
–0,8 m, hII = –0,9 m, hIII = –1,8 m. Desno: Polo`aj linij, kjer je
primerjano temperaturno polje; vI = 0,0 m (sredina), vII = 0,07 m
(povr{ina).
not known; however, several reasons why the results are
not identical might be identified: the solution procedure,
e.g., the energy equation in the commercial code is
solved iteratively and the temperature is obtained directly
from the equation, whereas in our method, the enthalpy
equation is solved first and the temperature is obtained
by solving the enthalpy-temperature relationship. Ano-
ther reason might lie in the differences between the
methods, e.g., in the commercial code a second-order
upwind technique is used, whereas in our case an adap-
K. MRAMOR et al.: SIMULATION OF CONTINUOUS CASTING OF STEEL UNDER THE INFLUENCE ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 281–288 285
Figure 10: Horizontal-velocity profiles at vertical position hIII = –1.8
m, at the end of the computational domain
Slika 10: Profili horizontalne hitrosti na mestu hIII = –1,8 m na koncu
ra~unske domene
Figure 7: Horizontal-velocity profiles at vertical position hII = –0.9
m, at the bottom boundary of the magnetic field
Slika 7: Profil horizontalne hitrosti na mestu hII = –0,9 m na spodnji
meji magnetnega polja
Figure 6: Horizontal-velocity profiles at vertical position hI = –0.8 m,
at the top boundary of the magnetic field
Slika 6: Profili horizontalne hitrosti na mestu hI = –0,8 m na zgornji
meji magnetnega polja
Figure 5: Horizontal-velocity profiles at vertical position hI = –0.8 m,
at the top boundary of the magnetic field
Slika 5: Profili horizontalne hitrosti na mestu hI = –0,8 m na zgornji
meji magnetnega polja
Figure 9: Horizontal-velocity profiles at vertical position hIII = –1.8
m, at the end of the computational domain
Slika 9: Profili horizontalne hitrosti na mestu hIII = –1,8 m na koncu
ra~unske domene
Figure 8: Horizontal-velocity profiles at vertical position hII = –0.9
m, at the bottom boundary of the magnetic field
Slika 8: Profili horizontalne hitrosti na mestu hII = –0,9 m na spodnji
meji magnetnega polja
tive upwind technique is used. Despite the differences in
the formulation, identical results are obtained in very
simple test cases; however, slight differences appear
when the examples become very complicated, like in the
case of continuous casting of steel, where the turbulent-
flow, heat-transfer and magnetic-field equations are
strongly coupled.
The agreement between the temperature fields calcu-
lated with the commercial code and those calculated with
the LRBFCM is better for the outer wall, where the shell
has already solidified, than for the centre of the billet,
where the liquid metal has not yet solidified.
K. MRAMOR et al.: SIMULATION OF CONTINUOUS CASTING OF STEEL UNDER THE INFLUENCE ...
286 Materiali in tehnologije / Materials and technology 48 (2014) 2, 281–288
Figure 12: Temperature layout at position vII = 0.07 m, at the mold
wall
Slika 12: Potek temperaturnega polja na mestu vII = 0,07 m na steni
kokile
Figure 11: Temperature layout at position vI = 0.0 m, at the centre of
the mold
Slika 11: Potek temperature na mestu vI = 0,0 m v sredi{~u kokile
Figure 13: Velocity layout at position vI = –0.8 m, at the top edge of
the applied magnetic field
Slika 13: Potek hitrostnega polja na mestu vI = –0,8 m na zgornjem
robu apliciranega magnetnega polja
Figure 16: Velocity layout at position hIII = –1.8 m, at the centre of
the mold
Slika 16: Potek hitrosti na mestu hIII = –1,8 m na spodnjem robu
apliciranega magnetnega polja
Figure 15: Temperature layout at position hII = –0.9 m, at the centre
of the mold
Slika 15: Potek temperature na mestu hII = –0,9 m, v sredi{~u kokile
Figure 14: Velocity layout at position vI = –0.8 m
Slika 14: Potek hitrostnega polja na mestu vI = –0,8 m
Finally, the velocity and temperature profiles were
compared for different magnetic-field strengths at
different vertical (velocity) and horizontal (temperature)
positions. As can be seen in Figures 13 to 16, velocity
profiles do not change much, when a weak magnetic
field (0.0026 T or 0.026 T) is applied, being only slightly
different from the velocity profiles obtained with no
magnetic field. An application of a weak magnetic field
in the x direction causes the velocity in the x direction to
increase (Figures 13 and 14) and the velocity in the y
direction to decrease (Figures 15 and 16). However, an
application of a strong magnetic field (0.26 T or 2.6 T)
alters the velocity profiles and can even change the
direction of the flow.
The effect of the externally applied magnetic field on
the temperature field is less pronounced. As can be seen
in Figures 17 and 18, the temperature is slightly raised
in the area of the application of the magnetic field.
6 CONCLUSIONS
In this paper, the initial numerical calculations of
continuously cast steel under the influence of a magnetic
field are presented and compared to the commercial
FVM-based computational fluid-dynamics (CFD) code
Fluent.23 The LRBFCM method gives similar results as
the commercial code; it is fully flexible, requiring no
mesh generation, and enabling a straightforward inclu-
sion of different turbulence models and constitutive
equations. In the future, a more realistic magnetic field
will be incorporated into the model and a realistic curved
geometry of the caster will be assumed. Several sensiti-
vity studies, in terms of the magnitude of an externally
applied magnetic field and the position of the coils
producing the magnetic field, will be performed. Finally,
the species-conservation equation will be added in order
to account for the macro-segregation.
Acknowledgements
The research in this paper was sponsored by the
Centre of Excellence for Biosensors, Instrumentation
and Process Control (COBIK) and the Slovenian Grant
Agency under programme group P2-0357 and project
L2-3651. This paper forms a part of a doctoral study of
the first author that is partly co-financed by the European
Union and the European Social Fund. The co-financing
is carried out within the Human resources development
operational programme for years 2007–2013, 1. Deve-
lopmental priorities: Encouraging entrepreneurship and
adaptation; Preferential directives, 1.3: Scholarship
schemes.
7 REFERENCES
1 W. R. Irwing, Continuous Casting of Steel, The Institute of Mate-
rials, London 1993
2 Steel Statistical Yearbook 2013, World Steel Association, Brussels
2013
3 B. [arler, R. Vertnik, A. Z. Lorbiecka, I. Vu{anovi}, B. Sen~i~,
Anwendung eines Stranggieß-Simulationsmodells bei [tore Steel–II,
BHM Berg- und Hüttenmännische Monatshefte, (2013), 1–9, doi:
10.1007/s00501-013-0147-7
4 B. Zhao, B. G. Thomas, S. P. Vanka, R. J. O’Malley, Metall. Mater.
Trans. B, 36 (2005) 6, 801–823
5 M. R. Aboutalebi, M. Hasan, R. I. L. Guthrie, Metall. Mater. Trans.
B, 26 (1995) 4, 731–744
6 D. S. Kim, W. S. Kim, K. H. Cho, ISIJ Int., 40 (2000) 7, 670–676
7 K. G. Kang, H. S. Ryou, N. K. Hur, Numer. Heat Trans. A, 48 (2005)
5, 461–481
8 N. Kubo, T. Ishii, J. Kubota, T. Ikagawa, ISIJ Int., 44 (2004) 3,
556–564
9 C. H. Moon, S. M. Hwang, Int. J. Numer. Meth. Eng., 57 (2003) 3,
315–339
10 B. [arler, R. Vertnik, Comput. Math. App., 51 (2006) 8, 1269–1282
11 D. C. Wilcox, Turbulence Modeling for CFD, DCW Industries, Inc.,
California 1993
12 W. D. Bennon, F. P. Incropera, Numer. Heat Transf. A-Appl., 13
(1988) 3, 277–96
K. MRAMOR et al.: SIMULATION OF CONTINUOUS CASTING OF STEEL UNDER THE INFLUENCE ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 281–288 287
Figure 18: Temperature layout at position vII = 0.07 m, at the mold
wall
Slika 18: Potek temperaturnega polja na mestu vII = 0,07 m na steni
kokile
Figure 17: Temperature layout at position vI = 0.0 m, at the centre of
the mold
Slika 17: Potek temperaturnega polja na mestu vI = 0,0 m v sredi{~u
kokile
13 K. Abe, T. Kondoh, Y. Nagano, Int. J. Heat Mass Transf., 37 (1994)
1, 139–51
14 B. [arler, R. Vertnik, K. Mramor, IOP Conference Series: Mater. Sci.
Eng., 33 (2012) 1, 12012–12021
15 A. Chorin, Math. Comput., 22 (1968) 104, 745–762
16 H. Lin, S. N. Atluri, CMES (Comput. Model. Eng. Sci.), 1 (2000) 2,
45–60
17 R. Vertnik, B. [arler, Int. J. Cast Metal. Res., 22 (2009) 1–4,
311–313
18 R. Vertnik, B. [arler, EABE [online] 2014 [cited 2014-01-02].
Available from World Wide Web: http://dx.doi.org/10.1016/
j.enganabound.2014.01.017
19 M. D. Buchmann, Radial Basis Function: Theory and Implementa-
tions, Cambridge University Press, Cambridge 2003
20 R. Franke, Math. Comput., 38 (1982) 157, 181–200
21 K. Mramor, R. Vertnik, B. [arler, CMES (Comput. Model. Eng.
Sci.), 92 (2013) 4, 327–352
22 K. Mramor, R. Vertnik, B. [arler, CMC (Comput. Mater. & Con-
tinua), 36 (2013) 1, 1–21
23 ANSYS® Fluent, 6.1, ANSYS, Inc.
K. MRAMOR et al.: SIMULATION OF CONTINUOUS CASTING OF STEEL UNDER THE INFLUENCE ...
288 Materiali in tehnologije / Materials and technology 48 (2014) 2, 281–288
E. ALTUNCU, H. ALANYALI: THE APPLICABILITY OF SOL-GEL OXIDE FILMS AND THEIR CHARACTERISATION ...
THE APPLICABILITY OF SOL-GEL OXIDE FILMS AND
THEIR CHARACTERISATION ON A MAGNESIUM
ALLOY
UPORABNOST SOL-GEL OKSIDNIH TANKIH PLASTI IN
NJIHOVA KARAKTERIZACIJA NA MAGNEZIJEVI ZLITINI
Ekrem Altuncu1,2, Hasan Alanyali2
1Sakarya University, Dept. Metallurgical and Materials Eng., Esentepe Campus, 54187 Sakarya, Turkey
2Kocaeli University, Machine-Metal Technology, Hereke Borusan Campus, Vocational School of Asim Kocabiyik, 41800 Kocaeli, Turkey
altuncu@sakarya.edu.tr
Prejem rokopisa – received: 2012-08-09; sprejem za objavo – accepted for publication: 2013-06-07
Magnesium and its alloys are widely used in many industrial applications because of their high specific strength
(strength/density) ratio. However, these applications are still restricted by the relatively poor surface resistance of these
materials. To overcome the inherent drawbacks a useful solution is to deposit a protective coating on the magnesium alloys. The
sol-gel process is an effective method for fabricating oxide films on a magnesium alloy in order to produce a higher corrosion
resistance. The objective of this study is to comparatively investigate the process and properties of repeated sol–gel oxide (ZrO2,
Al2O3) coatings. The coatings were characterized by SEM, XRD, ellipsometry and the effects of the process on their properties
were comparatively analysed.
Keywords: sol-gel, oxide films, magnesium alloy
Magnezij in njegove zlitine se uporabljajo v vrsti industrij zaradi visoke specifi~ne trdnosti (razmerje trdnost – gostota). Vendar
pa je industrijska uporaba omejena zaradi njihove slabe odpornosti povr{ine. Da bi to pomanjkljivost odpravili, je uporabna
re{itev varovalni nanos na povr{ini magnezijevih zlitin. Sol-gel postopek je u~inkovita metoda za izdelavo oksidnih tankih plasti
na magnezijevi zlitini za pove~anje njihove korozijske obstojnosti. Namen te {tudije je primerjava postopka in ugotavljanje
lastnosti ponavljajo~ih se sol-gel nanosov ZrO2 in Al2O3. Nanosi so bili ovrednoteni s SEM, XRD in elipsometrijo, izvr{ena pa
je bila tudi primerjava vpliva procesa na njihove lastnosti.
Klju~ne besede: sol-gel, oksidni nanosi, magnezijeva zlitina
1 INTRODUCTION
Magnesium and its alloys have a high specific
strength, excellent mechanical properties, a good damp-
ing capacity and a high electromagnetic shielding
capability. They can be used for a variety electronic and
many automotive parts.1,2 However, corrosion protection
is one of the main obstacles to the application of magne-
sium alloys in real environments. Conversion coating,
anodizing, plating, laser surface alloying and plasma
electrolyte oxidation have been applied on magnesium
alloys. Material designers attempted to improve the
corrosion resistance of the surface.3–5 However, with
these technologies the protection of magnesium alloys it
is hard to achieve a good cost-to-benefit ratio or there is
a higher energy consumption or environmentally adverse
effects. The sol–gel thin-film deposition method is
described as a practical, environmentally friendly and
cost-effective way to produce magnesium alloy surfaces.
It is well known that Al2O3 thin films have been exten-
sively used as an isolator, oxidation barrier, for anti-
corrosion and for anti-wear.6–8 However, sol–gel ZrO2
thin films, owing to their chemical stability, have
important applications as corrosion-resistant coatings on
metal surfaces.9 Thin (0.2–10 μm), dense, sol–gel oxide
coatings exhibited significant advantages in overcoming
metallic surface-corrosion problems. From the point of
view of synthesis, the sol–gel route offers versatile ways
to synthesize effective coatings with specific properties.
Surface functionality can be optimized by varying expe-
rimental parameters such as the chemical structure, the
composition and the ratio of precursors and complexing
agents, the rate and conditions of hydrolysis, the
synthesis media, the aging and curing conditions, and the
deposition procedure.10–14 The objective of this study is
to comparatively investigate the process and properties
of repeated sol-gel ZrO2 and Al2O3 coatings. The coat-
ings were characterized and the effects of the process on
their properties were analysed.15–17
2 EXPERIMENTAL
AM60 magnesium alloy plates were used as the
substrates. The specimens were cut into pieces of 20 mm
× 20 mm × 2 mm and their surfaces were ground and
degreased ultrasonically in acetone and then dried with
hot air. The deposition steps were as follows: sol prepa-
ration, gelation and heat treatment (Figure 1). Initially,
deionized water was heated in a glass beaker until it
reached 80 °C. Aluminium sec-butoxide (ASB) was
added to the water with continuous mixing. The molar
Materiali in tehnologije / Materials and technology 48 (2014) 2, 289–292 289
UDK 669.721.5:621.793 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(2)289(2014)
ratio of ASB to water was 1 : 25. As hydrolysis took
place, the temperature of the solution rose to 90 °C, at
which it was maintained for 3 min. Then, acetoacetate
(AcAc) was added to the mixture. The molar ratio of
AcAc to water was fixed at 1 : 25. After stirring for 2
min, nitric acid was added to the solution stepwise, with
stirring until it was transformed into a transparent solu-
tion. The pH value of the as-synthesized sol was about 3.
The Al2O3 sol–gel film was deposited onto the magne-
sium substrate via the dip-coating method and with-
drawing at a constant rate of 2 cm/min. The coated
substrate was dried in a clean cabinet at 130 °C for 10
min to produce the alumina gel film. Then, the supported
gel films were subjected to a thermal treatment at 500 °C
for 1 h in an oxygen atmosphere. Zirconium n-propoxide
(ZrNP) was diluted in n-propanol as the source of
zirconia, and then acetylacetone AcAc and deionized
water (diluted with n-PrOH) were added. The molar ratio
of Zr(n-OPr)4 : AcAc : H2O : n-PrOH was 1 : 2 : 2 : 60.
After stirring for 3 h at room temperature, a clear pre-
cursor solution was obtained.16 The ZrO2 coatings were
deposited by dipping the substrate into the sol and
withdrawing it at a constant rate of 5 cm/min. Thin
layers were prepared by repeating the dip coating and
after every deposition there was a heat treatment of the
sample at 500 °C for 1 h in an air atmosphere. The
thicker films were produced by repeating the dipping
process four times, followed by a thermal treatment.
The deposited films were studied by using scanning
electron microscopy (SEM), while the film thicknesses
and the refractive indices were measured using the ellip-
sometry method (Jobin Yvon, UVISEL HR 460). The
structures of the resulting films were examined by graz-
ing-incidence X-ray diffraction (XRD) with Cu K radi-
ation using a thin-film apparatus. The electrochemical
measurements were performed on a EG&G 273A-type
potentiostat, using Pt as an auxiliary electrode, a satu-
rated calomel electrode (SCE) as a reference electrode
and 3.5 % NaCl solution as an aggressive medium. The
potential was scanned from –1.8 V to 0.7 V (vs. SCE) at
a scanning rate of 0.5 mV/s.
3 RESULTS AND DISCUSSION
3.1 Morphology and Thickness
Different viscosities of sols were created by the
addition of the required amounts of nitric acid to the
synthesis mixture of aluminium sec-butoxide and AcAC
in n-propanal, which allowed us to control the rate of the
condensation reactions. Uniform, mesoporous alumina
coatings were obtained for the solvent-withdrawal rates
below 10 cm/min. The limiting withdrawal rate increases
as the solvent evaporation rate increases and it decreases
as the sol viscosity increases. The cross-sectional mor-
phologies indicate a mean thickness of approximately
1.6 μm and 2.3 μm for the repeated sol-gel coating for
E. ALTUNCU, H. ALANYALI: THE APPLICABILITY OF SOL-GEL OXIDE FILMS AND THEIR CHARACTERISATION ...
290 Materiali in tehnologije / Materials and technology 48 (2014) 2, 289–292
Figure 2: Surface morphology of sol-gel coatings using the dip-coating method (SEM micrographs)
Slika 2: Morfologija povr{ine sol-gel nanosa s potapljanjem (SEM)
Figure 1: Flow chart for the preparation of coatings
Slika 1: Postopek priprave nanosa
ZrO2 and Al2O3, respectively (Table 1). The deposition
efficiency is higher in the Al2O3 than in the ZrO2 sol-gel
coating. The porosity level of the alumina coating is
lower than the zirconia coating.
Table 1: Thicknesses of the coatings
Tabela 1: Debeline nanosov
Composition
Average bilayer thicknesses (nm) for 9
measurement
Min Average Max
ZrO2 63.42 65.2 ± 5 68.72
Al2O3 82.33 87.2 ± 4 90.14
It is clear from Figure 2 that both films were smooth,
compact and crack-free, and the thicknesses were
between 60 nm and 95 nm. In addition, a lot of bumps
and small pores were observed in this film.
3.2 Crystal Structure
Figure 3 illustrates the structure of the coatings
formed on the AM60 alloy substrate. Both of the
coatings have an amorphous structure. A small quantity
of crystallite structures were observed, depending on the
heat-treatment temperature and the time. However, the
crystallisation growth was inadequate for both coatings.
At higher temperatures than 600 °C for 1 h the cry-
stallites started to produce strong peaks in the XRD
pattern for the Al2O3 film. The characteristic peaks of
-Al2O3 and -Al2O3 that were observed imply the
existence of crystallized -Al2O3 and -Al2O3, which
were transferred from the alumina sol annealed at 500 °C
in the amorphous coating. The repeated sol-gel coating,
which indicates a more crystallized -Al2O3 phase, is
due to the direct introduction of Al2O3 particles. These
results are similar to those in Wang’s study.18 The
zirconia coatings are made up of m-ZrO2 (monoclinic
crystal structure) and t-ZrO2 (tetragonal crystal structure)
after a heat treatment at 600 °C for 4 h (Figure 3).
Partial spinodal decomposition was observed in the ZrO2
coating.
3.3 Optical Properties
Figure 4 shows the reflective spectra of a glass
substrate, alumina bi-layer coatings with 1.5 and 1.625
refractive indices in the region 300–800 nm. The
refractive index is changing between 1.9 and 2.2 for the
zirconia coating. The magnitude of the refractive index
decreases with the heat-treatment temperature and the re-
fractive indices of the Al2O3 and ZrO2 thin films decrease
with the increasing coating thickness. With an increase
in the crystallinity there is a decrease in the refractive
index of the thin films. The refractive index of the films
depends strongly on their morphology. Up to the heat-
treatment temperature, limiting the amorphous and
crystallized phase, it increases steadily, probably in
relation to an increasing densification of the layer. The
lowering of the refractive index for the crystallized films
is probably related to a lower densification of the films.
The decrease of light intensity is strongly lowered for
amorphous layers.
3.4 Corrosion Properties
In both cases, the polarization curves of the sol-gel-
coated substrates were appreciably different from that of
the bare substrates. The corrosion test showed a smooth
surface with no appreciable delamination or cracking of
the coating on the magnesium alloy substrates. The
open-circuit potential, Eoc, of the sol-gel-coated substra-
tes was significantly lower than that of the bare surfaces
E. ALTUNCU, H. ALANYALI: THE APPLICABILITY OF SOL-GEL OXIDE FILMS AND THEIR CHARACTERISATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 289–292 291
Figure 4: Refractive index (n) versus the wavelength /nm change for
the coatings
Slika 4: Spremembe lomnega koli~nika (n) v odvisnosti od valovne
dol`ine /nm pri nanosih
Figure 3: XRD patterns of the sol–gel coatings
Slika 3: XRD-posnetki sol-gel nanosov
(Figure 5). In addition, a distinct passivation region was
present for the coated substrates. The corrosion-pro-
tection properties of the sol–gel-derived coatings are
strongly dependent on the processing conditions. As
evident from Figure 5, the alumina (Al2O3) film corro-
sion protection is better than the zirconia (ZrO2) films.
4 CONCLUSIONS
The sol-gel-deposited oxide films formed with these
compounds had good adhesion, reflectivity or UV
protection. After the appropriate deposition and heat
treatment the oxide films were formed with very favour-
able properties, such as high adhesion, homogeneity and
density. The sol-gel alumina coatings developed on the
magnesium alloy surface provide superior corrosion
protection. After the heat treatments the XRD patterns
revealed crystalline structures. With a controlled heat
treatment a slight increase in the corrosion resistance
was observed. However, both of the alumina- and zir-
conia-based coatings can be used for optical applications
with suitable heat-treatment conditions. The refractive
index of the alumina film was lower than that of the
zirconia films. This depends on the crystallite structure
and the surface morphology, such as bubbles and the
porosity effect.
Acknowledgements
The authors would like to thank Ýstanbul Technical
University, Metal. and Mat. Eng. Laboratory, JOBIN
YVON Lab. Also, thanks go to the Researchers Céline
Marchand and Michel Stchakovsky, Researcher Stefan
Winter, Metallic Mat. Dept. of Universität Des Saar-
landes for their experimental support.
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structured hybrid organic-inorganic thin films, J. Mater. Chem., 15
(2005), 3598–3627
14 F. Mammeri, E. Le Bourhis, L. Rosez, C. Sanchez, Mechanical pro-
perties of hybrid organic-inorganic materials, J. Mater. Chem., 15
(2005), 3787–3811
15 E. Altuncu, Practicability of a ceramic thin film formation technique
in order to improve the surface properties of magnesium and its
alloys, Kocaeli University, Metallurgical and Materials Engineering,
Master Thesis, 2004
16 H. Li, K. Liang, L. Mei, S. Gu, S. Wang, Oxidation protection of
mild steel by zirconia sol–gel coatings, Materials Letters, 51 (2001),
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free alumina film, Materials Research Bulletin, 42 (2007), 600–608
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E. ALTUNCU, H. ALANYALI: THE APPLICABILITY OF SOL-GEL OXIDE FILMS AND THEIR CHARACTERISATION ...
292 Materiali in tehnologije / Materials and technology 48 (2014) 2, 289–292
Figure 5: Polarization curves for the thin oxide films: ZrO2 and Al2O3
Slika 5: Polarizacijski krivulji za tanek oksidni nanos ZrO2 in Al2O3
D. GOLUBOVI] et al.: TESTING THE TRIBOLOGICAL CHARACTERISTICS OF NODULAR CAST IRON ...
TESTING THE TRIBOLOGICAL CHARACTERISTICS OF
NODULAR CAST IRON AUSTEMPERED BY A
CONVENTIONAL AND AN ISOTHERMAL PROCEDURE
PREIZKU[ANJE TRIBOLO[KIH LASTNOSTI NODULARNE
LITINE, MEDFAZNO KALJENE PO KONVENCIONALNEM IN
IZOTERMI^NEM POSTOPKU
Du{an Golubovi}1, Pavel Kova~2, Borislav Savkovi}2, Du{an Je{i}3,
Marin Gostimirovi}2
1Faculty of Mechanical Engineering, Vuka Karadzica 30, 71213 East Sarajevo, Bosnia and Herzegovina
2Faculty of Technical Science, Trg D. Obradovica 6, 21000 Novi Sad, Serbia
3University of Banja Luka, Faculty of Engineering, V. Stepe Stepanovica 75, 7800 Banja Luka, Bosnia and Herzegovina
pkovac@uns.ac.rs
Prejem rokopisa – received: 2013-03-01; sprejem za objavo – accepted for publication: 2013-07-15
In this paper the heat treatment of ductile iron and the tribological properties of contact pairs (pin and disk) were investigated.
Two types of nodular cast iron, EN-GJS-500-7 and EN-GJS-700-2, austempered using an isothermal and a conventional proce-
dure, were tested. The friction-and-wear test was carried using a PIN on DISC Tribometer and the PQ test. The methodology of
the classical and nodular cast-iron isothermal austempering and the methodology for the examination are also described. From
analysing the results it was concluded that the tribological characteristics depend on the structural characteristic of the nodular
cast iron, which are determined by heat treatment. The tested sample, austempered using a classic approach, gives the best
performance in terms of friction, but it also showed the worst performance in terms of wear. In the same heat-treatment regime
EN-GJS-500-7 is characterized by better tribological characteristics compared to the EN-GJS-700-2. The obtained results can
be used for the proper selection of the type and regime of the nodular cast iron’s heat treatment, with the aim of improving the
exploitation characteristics of the contact pairs.
Keywords: nodular cast iron, heat treatment, friction, wear
^lanek obravnava preizku{anje nodularne litine s toplotno obdelavo in tribolo{ke lastnosti kontaktnih parov (klin in disk). Preiz-
ku{eni sta bili dve vrsti nodularne litine EN-GJS-500-7 in EN-GJS-700-2, medfazno kaljeni po konvencionalnem in izoter-
mi~nem postopku. Preizkus trenja in obrabe je bil izveden s "pin-on-disk" tribometrom in PQ-preizkusom. Prikazana je
metodologija izotermnega pobolj{anja klasi~ne in nodularne litine ter metodologija preiskave. Analiza rezultatov je pokazala, da
so tribolo{ke lastnosti odvisne od zna~ilnosti strukture nodularne litine, ki je dolo~ena z vrsto nodularne litine in toplotno
obdelavo. Preizkusni vzorec, klasi~no medfazno kaljen, izkazuje najbolj{e rezultate glede trenja, hkrati pa najve~jo obrabo. V
istem re`imu toplotne obdelave ima EN-GJS-500-7 bolj{e tribolo{ke lastnosti v primerjavi z EN-GJS-700-2. Dobljeni rezultati
se lahko uporabijo pri izbiri primernega re`ima toplotne obdelave nodularne litine, z namenom izbolj{anja lastnosti kontaktnih
parov med rabo.
Klju~ne besede: nodularna litina, toplotna obdelava, trenje, obraba
1 INTRODUCTION
Modern materials used to create the elements of
machinery, equipment and vehicles (cars, trucks, trac-
tors, etc.) must possess, in addition to a high hardness, a
good toughness and a high fatigue strength, as well as a
more satisfactory resistance to friction and wear.1,2 The
first three properties of metallic materials are determined
by standard methods. Information about them can be
found in the relevant literature and standards, as well as
in the technical documentation of the manufacturer of
the materials.
The resistance to abrasive wear, as the tribological
characteristic of materials, depends not only on their
physical and chemical properties, but also on the condi-
tions under which the contact between the elements of
the tribomechanical system is achieved, as well as the
properties of the other elements in the contact.3 The
measurement of the tribological properties of these two
materials is carried out, usually on tribometers that fulfil
one of the three possible geometries of contact (touch at
a point on the line or on the surface). The paper presents
the results achieved for a linear contact.4,5
In this paper, using a realized tribomechanical system
with a geometric line contact, we examine the basic tri-
bological properties of materials in terms of friction and
wear. Part of the results of the tribological characteristics
of the two types of nodular cast iron, isothermal and the
improved classical procedure, obtained in laboratory
conditions and presented in this study, indicate the
importance of heat treatment on the resistance to friction
and wear.6–8
The use of ductile iron in scientific research has a
strong presence, so the paper9 also considers the micro-
structures and tribological behaviour of different roller
mill specimens, having basically nodular cast iron com-
positions, produced by conventional static and vertical
centrifugal casting methods. Also, in the research work,10
the initiative was taken to improve the surface hardness
as well as the wear resistance of the as-received nodular
Materiali in tehnologije / Materials and technology 48 (2014) 2, 293–298 293
UDK 669.131.6:531.43:620.178.1 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(2)293(2014)
cast iron using the pack-carburizing technique. The
objective of the paper11 is to evaluate the fatigue life of
the nodular cast iron EN-GJS-500-7, which is used for
railway brake discs. In the paper12 the results of the
abrasive and adhesive wear resistance of selected grades
of nodular cast iron with carbides are presented. The
materials of the friction pairs, tested at the stand and sub-
jected to heat treatment and chemical processing in order
to attain specific parameters of their surface layers, were
studied in the work.13 The studies conducted enabled a
determination of the abrasive wear values for the
material samples tested, having entailed the surface-layer
parameters and the factors related to the operation of
actual structural components used in automotive engi-
neering.
Through a literature review we can see what has
already been done in terms of the wear resistance of
ductile iron.14–16 The abrasion wear rates of two-step
austempered ductile cast iron (ADI) were investigated.17
The wear resistance of the ductile cast iron can be im-
proved through a different heat-treatment procedure and
surface engineering techniques, each having some limita-
tions and drawbacks.17,18
A microstructure of lower bainite consisting of acicu-
lar (needle) bainite stable ferrite and carbon-enriched
retained austenite, is provided if the isothermal transfor-
mation is carried out at low temperatures. The micro-
structure of the upper bainite, which consists of ferrite-
tile bainite and stable carbon-enriched retained austenite
is obtained if the isothermal transformation is carried at a
temperature of 390 °C.
With isothermal transformation temperature the dis-
tance between the ferrite tile bainite increases while
reducing the remaining shares, i.e., the volume fraction
of retained austenite was increased.
The impact strength and the fracture toughness can
be strongly reduced. The light microscope, however,
does not show the difference in the microstructure. The-
refore, for a detailed correlation of the mechanical pro-
perties depending on the microstructure it is necessary to
apply a high-resolution transmission electron microscope
(TEM).
For ductile iron of a ferrite base it is strongly recom-
mend to use an isothermal transformation temperature of
350 °C.19
The objectives of the research were to prove that for a
variety of thermal processing on the same material, a
different value for the level of material wear is obtained.
In addition, they want to show that the same level of heat
treatment applied to different materials, for some sizes
that represent the level of friction and wear, are of great
importance and influence, whereas for some they are not.
2 EXPERIMENTAL PROCEDURE
2.1 Material
In all the experiments we used a hardened carbon
steel (C40E) pin of guaranteed chemical composition
and a hardness of 52 HRC. The test program includes
two types of nodular cast iron, EN-GJS-500-7 and
EN-GJS-700-2, austempered by an isothermal and con-
ventional procedure.19,20
The melting of the material was carried out in a net-
work frequency induction furnace pot with an acid
coating capacity of 2.5 t in a foundry. After this the
desulfurization was carried out in a batch nodulation
closed pot using a "sandwich" method with a simulta-
neous modification at a temperature of 1500 °C. A
secondary modification was made before the actual
casting. The as-prepared metal was poured into a total 11
Y tube to BS 2789.
All the Y tubes were covered by two tube sections,
25 mm diameter and length 250 mm. Attention was also
given to two tubes by tensile tests in the irons JUS
C.J2.022 and three tubes with V-notch toughness testing
JUS C.A4.004.
The isothermal austenitising was performed in a
muffle furnace in a protective atmosphere at 900 °C for
90 min prior to preheating at 520 °C for 60 min. The
austempering was carried out in a solar bath (made
specifically for this purpose) at a temperature of 390 °C
with a hold time of 30 min at a speed of movement for
the salts of 0.6 m/s.
Their basic characteristics are listed in Table 1.19 We
present the chemical composition, the conventional
austempering regimes, the isothermal austempering and
the hardness of the tested parts.21
Figure 1 shows the metallographic structure of the
nodular cast iron used in the test. It is important to men-
tion that the EN-GJS-500-7 is ferrite-pearlite structure
D. GOLUBOVI] et al.: TESTING THE TRIBOLOGICAL CHARACTERISTICS OF NODULAR CAST IRON ...
294 Materiali in tehnologije / Materials and technology 48 (2014) 2, 293–298
Figure 1: Microstructures of: a) EN-GJS-500-7 and b) EN-GJS-700-2
Slika 1: Mikrostruktura: a) EN-GJS-500-7, b) EN-GJS-700-2
based (50 % ferrite and 50 % pearlite and more than 90
% of graphite-type-K size 3) and the EN-GJS-700-2 is
predominantly pearlitic-structure based. During the
structural test of the nodular cast iron the etching was
performed with 2 % HNO3 and increased 500-times.
Figure 1 shows that the black parts have a pearlitic
structure, and the white parts are a softer ferrite struc-
ture. In the middle are the visible graphite nodules sur-
rounded by a highly visible ferrite structure.4
Figure 2 shows the austempered ductile iron micro-
structure. The mechanical properties depend on the
structure of the material. One pearlitic ductile cast iron
has, before the isothermal treatment, an improved micro-
structure consisting of pearlite, and in it are the graphite
nodules. After the isothermal improvement the material
microstructure consists of bainite and retained austenite.
The amount of retained austenite is relatively high (aged
between 20 % and 40 %). The austenite contributes to
the high toughness and the ductility of the austempered
ductile iron.
Irons that were improved during a low-temperature
isothermal transformation have a structure of lower
bainite and about 400 HB hardness (HRC about 43), and
are suitable for first gear and applications that must be
resistant to high pressures.
Irons that were improved at higher temperatures of
isothermal transformation have an upper bainite structure
and a hardness from 260 HB to 350 HB (HRC 27 to 38),
and have a particularly high ductility, impact strength
and fatigue strength.
Figure 3 shows a diagram of the classical austempe-
ring procedure of nodular cast iron.14 The process of
austempering conventional materials was achieved by
heating up to a temperature of 900 °C, where it is held
for about 90 min and then by rapid cooling (water or air)
to room temperature. After this it is heated to a tempera-
ture of 520 °C and is then cooled to room temperature.
Figure 4 shows a diagram of the discs’ isothermal
austempering, made of an EN-GJS-500-7 ferrite-pearlite
base and an EN-GJS-700-2 made of a very pearlite base.
The austempering is done by heating the disks to a tem-
D. GOLUBOVI] et al.: TESTING THE TRIBOLOGICAL CHARACTERISTICS OF NODULAR CAST IRON ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 293–298 295
Figure 2: Microstructure of austempered ductile iron: a) EN-GJS-
500-7 and b) EN-GJS-700-2
Slika 2: Mikrostruktura izotermno pobolj{ane nodularne litine: a) EN-
GJS-500-7, b) EN-GJS-700-2
Figure 3: Diagram of conventional discs’ austempering
Slika 3: Diagram klasi~nega pobolj{anja plo{~ic
Table 1: The chemical composition and heat-treatment regimes of nodular cast iron EN-GJS-500-7 and EN-GJS-700-2
Tabela 1: Kemijska sestava in re`imi toplotne obdelave nodularne litine EN-GJS-500-7 in EN-GJS-700-2
Material of Disk
Heat Treatment Hardness
HB Legend StructureTa/°C Tp/°C / t/min
EN-GJS-500-7
900
390/30 355 EN-GJS-500-7-30
ferrite-pearlite
520/60 302 EN-GJS-500-7-k
EN-GJS-700-2
520/60 320 EN-GJS-700-2-k
mostly pearlitic
390/30 365 EN-GJS-700-2-30
Chemical composition, w/%
C Si Mn Mg P S Cu Ni
EN-GJS-500-7 3.85 2.9 0.076 0.035 0.02 0.004 1.5
EN-GJS-700-2 3.76 2.35 0.51 0.02 0.004 1.48 1.5
perature of 900 °C, at which it is held for about 90 min,
and then fast cooled to a temperature of 390 °C.
The discs of EN-GJS-500-7 and EN-GJS-700-2 were
kept for 90 min and 30 min at the temperature of 390 °C.
In this way we get a bainite structure (upper or lower
bainite) that enables improved toughness, and in this
case the wear resistance (abrasion of discs).
2.2 Method
Figure 5 provides a schematic view of the contact
pair (pin and disk) and an image using the Tribometer
TPD-93 where the coefficient of friction is measured for
the inline contact.5 With this equipment it is possible to
measure the friction force, the normal force, the coeffi-
cient of friction, the acoustic emission temperature of the
oil and the contact temperature. The characteristics of
the tribometer are: normal load 1–500 N, sliding speed
0.1–5 m/s, the motor power A is 1.5 kW, the motor
power B is 0.37 kW, and the overall dimensions are 1
200 mm × 600 mm × 1 300 mm.
By measuring the normal and friction forces at diffe-
rent levels of load and sliding speed v = 1.3 m/s, the
friction coefficients for a given combination of contact
pairs are determined. In the given range of load Fn = 1–3
N the differences in friction coefficients are less than
5 %, so the comparison is the heat-treated nodular cast
iron EN-GJS-500-7 and EN-GJS-700-2 derived for the
adopted level of load Fn = 2 N. This force was set on the
tribometer.
To determine the PQ index we used the PQ 2000 Par-
ticle quantifier shown in Figure 6. According to the pro-
cedure methodology the value of the PQ index is directly
proportional to the quantity of the wear products (greater
than 5–10 μm) that are contained in the oil used for
lubrication of the contact zone, during the disk sliding
along the pin. For this purpose the used oil was Polar
INA 55-K.
The PQ index is the average value obtained from
several measurements. The average measurements of the
PQ index for a sample of oil containing products in the
research conducted was 10. In most cases the measure-
ment error did not exceed 10 %.
The duration of the contact in the experimental ope-
rations was about 30 min. The friction force was
measured at the beginning and at the end of the actual
contact (tp = 1 min, tk = 29 min). There was a line contact
D. GOLUBOVI] et al.: TESTING THE TRIBOLOGICAL CHARACTERISTICS OF NODULAR CAST IRON ...
296 Materiali in tehnologije / Materials and technology 48 (2014) 2, 293–298
Figure 5: a) Schematic view of the contact pair (pin and disk) and b)
image of Tribometer TPD-93
Slika 5: a) Shematski prikaz kontakta para (pin in disk) in b) tribo-
meter TPD-93
Figure 6: PQ 2000 Particle quantifier
Slika 6: Kvantifikator delcev PQ 2000
Figure 4: Diagram isothermal discs’ austempering
Slika 4: Diagram izotermi~nega pobolj{anja plo{~ic
between the pin and the disc. The measurement of the
PQ indices in a sample of oil in which there are wear
products was performed 3 times on 10 samples.
The PQ index was determined based on the quantity
of wear products per mg of oil produced during the
sliding of one element of the tribomechanical system
over the other for t = 30 min. The different values of the
PQ index indicate the amount of wear product for the
samples in contact in the lubricant are a direct conse-
quence of the greater or lesser intensity of the wear of
the contact pair. The test covered above thermally aus-
tempered discs of nodular cast iron EN-GJS-500-7 and
EN-GJS-700-2. The wear of the pins for 30 min was
carried out under the same contact conditions and based
on samples of oil particle quantifiers over the levels of
the PQ indices.
3 RESULTS AND DISCUSSION
3.1 Experimental results
3.1.1 According to the experimental results
In order to determine the influence of the heat-treat-
ment type on the tribological properties of ductile iron,
on two types of ductile iron EN-GJS-500-7 and
EN-GJS-700-2, the test of the friction and wear accord-
ing to a pre-defined methodology of experimental
research was conducted. Figure 7 shows the histogram
displaying values of the friction coefficients at the line of
contact for the disc and pin, obtained from the derived
values of forces, recorded by the measuring instru-
mentation. The histogram shows that the heat treatment
of the ductile iron EN-GJS-500-7 has a small (approxi-
mately 2 %), practically negligible, impact on the value
of the friction coefficient. Simultaneously, the EN-GJS-
700-2 ductile iron isothermally improved by about 8 %
higher coefficient of friction than the same nodular cast
iron improved using the classical procedure. On the other
hand, the coefficient of friction in a ductile iron
EN-GJS-700-2 increased by 15–20 % more than the
values that were obtained for ductile iron EN-GJS-500-7.
Figure 8 shows a histogram of the PQ index values
for the observed contact pairs (pin and disk), the load
level of 2 N and the slip velocity v = 1.3 m/s. The
histogram displays the values of the PQ index, different
than the coefficient of friction, which shows that the
treatment regimens have a very large impact on the
aspects of wear. In both nodular irons the PQ index is
higher for iron improved by the classical procedure,
especially for EN-GJS-500-7 (approx. 115 %). On the
other hand, the ductile iron EN-GJS-500-7, improved the
isothermal and classic approach, has a significantly
lower PQ index for iron EN-GJS-700-2 (approx. 10–90
%).
Figure 9 shows a comparative review of the tribolo-
gical characteristics of the tested discs, in terms of
friction and wear, i.e., by the indicators Pt = μr/μx and Ph
= PQr/PQx. Here, μr and PQr are the values of the friction
coefficient and the wear index of the reference contact
pairs (it was a material austempered using the classical
procedure), whereas μx and PQx are the values of the
investigated contact pair. In doing so, for the reference a
contact couple with the lowest coefficient of friction and
wear is chosen, which receives the same indicator value
for the maximum of 100 %.
D. GOLUBOVI] et al.: TESTING THE TRIBOLOGICAL CHARACTERISTICS OF NODULAR CAST IRON ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 293–298 297
Figure 9: Histogram Pt, Ph index of the tested materials with different
heat treatments
Slika 9: Histogram Pt, Ph-indeksov preizku{enih materialov z razli~no
toplotno obdelavo
Figure 8: Histogram values of PQ index
Slika 8: Histogram prikazuje dobljene vrednosti PQ-indeksov
Figure 7: Histogram of the obtained friction coefficients for the tested
materials
Slika 7: Histogram prikazuje dobljene koeficiente trenja preizku{anih
materialov
3.2 Discussion
By analysing the results obtained, it can be concluded
that ductile iron, austempered using the classic approach,
has different tribological properties in terms of friction
and wear.6,7,14–17 The tribological characteristics signi-
ficantly depend on the type of nodular cast iron, and the
heat treatment.
By comparing the tribological characteristics of the
tribomechanical system elements, made of some kind of
ductile iron, some coefficients of friction lead to the con-
clusion that the thermal treatment has no practical
impact. Differences in the tribological characteristics of
the same elements, however, are very large if the tribo-
logical characteristics are determined by the size of their
wear, which occurs after a specified duration of the con-
tact.
The difference in the coefficients of friction is slight-
ly higher when comparing the two studied materials. The
coefficient of friction of the ductile iron EN-GJS-700-2
is slightly larger than the coefficient of friction of the
material EN-GJS-500-7.
The PQ index is higher in the material EN-GJS-
700-2. Both classical materials’ heat-treated PQ index is
greater than the isothermally processed material. Based
on this we can conclude that the higher wear is in the
conventionally heat-treated material. The isothermal heat
treatment helps to reduce the wear of these materials
The indicator of the friction coefficient between the
investigated disks Pt is higher for the classical heat-
treated material. Also, the indicator Pt is significantly
lower for the material EN-GJS-700-2, which suggests
that the friction is less for this material. The values of the
parameters Ph are higher for the materials EN-GJS-
500-7, whether it is a classic procedure or isothermally
processed. It is confirmed that the materials that are
isothermally processed have less contact wear than the
classically heat-treated materials.
4 CONCLUSIONS
The tested sample, a disc made of EN-GJS-500-7
austempered using the classic approach, gives the best
performance in terms of friction, but it also showed the
worst in terms of wear. The disc made of a hardened
EN-GJS-500-7 and isothermally austempered has excel-
lent features from the point of view of friction and wear.
In the same heat-treatment regime EN-GJS-500-7 is
characterized by better tribological characteristics than
the EN-GJS-700-2.
The laboratory studies of the improved tribological
properties of nodular cast iron can be used during the
selection of materials and heat treatment in order to
reduce the friction and wear in the contact pairs’ explo-
itation conditions. Further research is planned to follow
the tribological properties of selected materials in
specific industrial conditions.
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iron wear resistance through local surface reinforcement, Wear,
274–275 (2012), 267–273
19 D. Je{i}, Tribological Properties of Nodular Cast Iron, Monography,
Journal of the Balkan Tribological Association, Sofia, 2000, 125
20 K. Brandenburg, Machining Austempered Ductile Iron, Manufactur-
ing Engineering, 128 (2002), 5
21 A. R. Ghaderi, M. N. Ahmadabadi, H. M. Ghasemi, Effect of gra-
phite morphologies on the tribological behavior of austempered cast
iron, Wear, 255 (2003) 1–6, 410–416
D. GOLUBOVI] et al.: TESTING THE TRIBOLOGICAL CHARACTERISTICS OF NODULAR CAST IRON ...
298 Materiali in tehnologije / Materials and technology 48 (2014) 2, 293–298
M. B. DJURDJEVIC et al.: QUANTIFICATION OF THE COPPER PHASE(S) IN Al-5Si-(1–4)Cu ALLOYS ...
QUANTIFICATION OF THE COPPER PHASE(S) IN
Al-5Si-(1–4)Cu ALLOYS USING A COOLING CURVE
ANALYSIS
UPORABA ANALIZE OHLAJEVALNE KRIVULJE ZA OCENO
KOLI^INE BAKROVIH FAZ V ZLITINAH Al-5Si-(1-4)Cu
Mile B. Djurdjevic1, Srecko Manasijevic2, Zoran Odanovic1, Natalija Dolic3,
Radomir Radisa2
1IMS Institute, Bulevar Vojvode Misica 43, 11 000 Belgrade, Serbia
2Lola Institute, Kneza Viseslava 70a, 11 000 Belgrade, Serbia
3University of Zagreb, Faculty of Metallurgy, Aleja narodnih heroja 3, 44 103 Sisak, Croatia
srecko.manasijevic@li.rs
Prejem rokopisa – received: 2013-04-02; sprejem za objavo – accepted for publication: 2013-06-18
The aim of this paper is to demonstrate that it is possible to characterize the development and quantify the area percentage of
Cu-enriched phases in Al-5Si-(1-4)Cu alloys by applying a cooling-curve analysis. It is shown that several distinct Cu-enriched
phases are manifested as peaks on the first derivative of the cooling curve. The total area percentage of the Cu-enriched phases
is defined as the ratio of the area between the first derivative of the cooling curve and the hypothetical solidification path of the
Al-Si-Cu eutectic to the total area between the first derivative of the cooling curve and the base line. These calculations, based
on the cooling curve analyses, are compared with the image-analysis and chemical-analysis results in order to verify the
proposed method. There is a good correlation between the measured and calculated values for the area of the Cu-rich phase in
Al-5Si-(1–4)Cu alloys.
Keywords: aluminum alloys, thermal analysis, cooling-curve analysis, image analysis
Namen tega ~lanka je predstaviti mo`nost ocene nastanka in koli~insko dolo~iti podro~ja s Cu bogatih faz v zlitinah
Al-5Si-(1-4)Cu z analizo ohlajevalne krivulje. Pokazano je, da se ve~ lo~enih, s Cu bogatih faz ka`e v obliki vrhov v prvem
odvodu krivulje ohlajanja. Skupni dele` obmo~ja s Cu bogatih faz je dolo~en kot razmerje povr{in med prvim odvodom krivulje
ohlajanja in hipoteti~ne poti strjevanja Al-Si-Cu-evtektika ter celotno povr{ino prvega odvoda krivulje ohlajanja in osnovno
linijo. Izra~uni, ki temeljijo na analizi ohlajevalnih krivulj, so bili primerjani z analizo slik in rezultati kemijske analize, da bi
potrdili predlagano metodo. Obstaja dobra korelacija med izmerjenimi in izra~unanimi vrednostmi podro~ij s Cu bogatih faz v
zlitini Al-5Si-(1–4)Cu.
Klju~ne besede: aluminijeve zlitine, termi~na analiza, analiza ohlajevalne krivulje, analiza slik
1 INTRODUCTION
The automotive industry makes frequent use of the
Al-Si-Cu series of aluminum alloys. In order to ensure
that cast components have good mechanical properties,
their as-cast microstructures must be closely monitored.
Two eutectic microconstituents are primarily responsible
for defining the microstructures of Al-Si-Cu series
alloys: Al-Si and Al-Cu. Both of these eutectics can be
detected on a thermal-analysis (TA) cooling curve, or
more precisely, on its first derivative. The solidification
of Al-Si-Cu series alloys and the formation of
Cu-enriched phases can be described, according to many
authors, as follows:1–4
1. A primary -aluminum dendritic network forms
between 580–610 °C. The exact temperature depends
mainly on the amounts of Si and Cu in an alloy. This
leads to an increase in the concentration of Si and Cu
in the remaining liquid.
2. Between 570–555 °C (the Al-Si eutectic tempe-
rature), a eutectic mixture of Si and -Al forms,
leading to a further localized increase in the Cu
content of the remaining liquid.
3. At approximately 540 °C, Mg2Si and Al8Mg3FeSi6
phases begin to precipitate.
4. At approximately 525 °C, a "massive" or "blocky"
Al2Cu phase (containing approximately w = 40 %
Cu) forms together with -Al5FeSi platelets.
5. At approximately 507 °C, a fine Al-Al2Cu eutectic
phase forms (containing mass fractions approxima-
tely 24 % Cu). If the melt contains more than 0.5 %
Mg, an ultra-fine Al5Mg8Cu2Si6 eutectic phase also
forms at this temperature. This phase grows from
either of the two previously mentioned Al2Cu phases.
A metallographic analysis of the TA test samples,
presented in Figure 1, combined with an X-ray micro-
analysis has confirmed that Cu-enriched phases appear
with three main morphologies: the blocky type, the
eutectic type and the fine eutectic type.3,5,6
The Al-5Si-(1–4)Cu alloys are characterized by the
presence of the two eutectics (Al-Si and Al-Si-Cu) that
are primarily responsible for the mechanical properties
of these alloys. Both eutectic temperatures can be
detected on a TA cooling curve, or more precisely, on its
first derivative. The eutectic-formation temperatures can
Materiali in tehnologije / Materials and technology 48 (2014) 2, 299–304 299
UDK 669.715:536.7 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(2)299(2014)
help to define the maximum temperature, to which
castings can be exposed during a solution treatment (i.e.,
by defining the temperature, at which incipient melting
will take place). Unfortunately, the total amount of the
Cu-enriched phases present in an as-cast part can, so far,
only be measured using a metallographic analysis. This
information is critical because these Cu-rich phases play
a significant role in the heat-treatment process and can
have a negative influence on the mechanical properties of
the Al-5Si-(1–4)Cu alloys. The goal of this paper is to
demonstrate that it is possible to quantify and charac-
terize the development of the Cu-enriched phases in the
Al-5Si-(1–4)Cu alloys using the TA system. This esti-
mation is verified using quantitative metallography (an
image analysis (IA)) and a chemical analysis (optical
emission spectroscopy (OES)).
2 EXPERIMENTAL PROCEDURES
Three Al-Si-Cu alloys with the chemical composi-
tions presented in Table 1 were produced. Their che-
mical compositions were determined using the OES.
Liquid test samples with the masses of approximately
300 g were poured into thermal-analysis steel test cups.
The weight of a steel test cup was 50 g. Two K-type ther-
mocouples were inserted into the melt and the tempera-
tures between 700–400 °C were recorded. The tip of a
thermocouple was always kept at the constant height, 15
millimeters from the bottom of the crucible. The accu-
racy of a thermocouple was ± 0.5 °C. The data for TA
was collected using a high-speed data-acquisition system
linked to a personal computer. The cooling conditions
were kept constant during all the experiments and the
cooling rate was approximately 6 K min–1. The cooling
rate was calculated as the ratio of the temperature diffe-
rence between the liquidus and solidus temperatures to
the total solidification time between these two tempera-
tures. Each TA trial was repeated three times. Conse-
quently, a total of nine samples were gathered. In all the
cases, the masses of the thermal-analysis test samples
were virtually identical.
The samples for the microstructural analysis were cut
from the TA test samples, close to the tips of the thermo-
couples. The cross-sections of the specimens were
ground and polished on an automatic polisher using
standard metallographic procedures. The samples were
observed with a scanning electron microscope (SEM)
using the magnifications between 200-times and
5000-times. Qualitative and quantitative assessments of
the chemical compositions of the Cu-enriched phases
were done using an energy dispersive spectrometer
(EDS). The area fractions of the Cu-enriched phases
were calculated using image-analysis software linked to
a microscope, under a magnification of 500-times.
Twenty-five analytical fields were measured for each
sample and the final area fraction was expressed as the
mean value.
3 RESULTS AND DISCUSSION
3.1 Thermal-analysis results
Three representative TA cooling curves obtained for
the Al-5Si-1Cu, Al-5Si-2Cu and Al-5Si-4Cu alloys are
presented in Figure 2. The cooling rate for all three
curves was approximately 6 K min–1. Figure 3 shows
that the increasing Cu amount of the melt lowers all the
M. B. DJURDJEVIC et al.: QUANTIFICATION OF THE COPPER PHASE(S) IN Al-5Si-(1–4)Cu ALLOYS ...
300 Materiali in tehnologije / Materials and technology 48 (2014) 2, 299–304
Figure 1: SEM micrographs (BSE images) with the characteristic
morphologies of Cu-enriched phases found in the investigated alloys:
a) the blocky (#1) and eutectic types (#2), b) the fine eutectic type
(#3)6
Slika 1: SEM-posnetka (BSE-posnetka) z zna~ilno morfologijo s Cu
bogatih faz v preiskovih zlitinah: a) kockasta (#1), evtektik (#2), b)
drobni evtektik (#3)6
Table 1: Chemical compositions (mass fractions, w/%) of the synthetic alloys
Tabela 1: Kemijska sestava (masni dele`i, w/%) sinteti~nih zlitin
Alloy Si Cu Fe Mg Mn Zn Ni Al
Al-5Si-1Cu 4.85 1.03 0.09 0.14 0.01 0.01 0.007 residual
Al-5Si-2Cu 5.01 2.06 0.10 0.26 0.01 0.01 0.007 residual
Al-5Si-4Cu 4.89 3.85 0.09 0.16 0.01 0.01 0.009 residual
characteristic solidification temperatures (TLIQ , TCOH ,
TEUT
Al -Si and TEUT
Al -Si -Cu) except the solidus temperature that
is almost constant for all the investigated alloys.
The first derivatives of the cooling curves are pre-
ented in Figure 4. It is apparent that the shapes of the
first derivative curves strongly depend on the Cu amount
in the melt. The Cu-rich area is particularly affected by
different Cu amounts.
The numbers and shapes of the peaks visible in the
Cu-enriched region of the first-derivative curves show a
strong relationship with the amount of Cu present in the
alloy. It can also be observed in Figure 5 that an increase
in the Cu amount increases the solidification time of the
Cu-rich eutectic phase. The precipitation temperature of
the Cu-enriched phases decreases when Cu increases
from mass fractions 1 % to 4 %. The Cu-enriched phase
represented by the first peak on the cooling curve in Fig-
ure 5 (5 % Si, 1 % Cu in the alloy) began to precipitate
at 542.7 °C and the Cu-enriched phase represented by
the second peak precipitated at 503.2 °C. For the alloy
with 5 % Si and 2 % Cu, three peaks precipitated at
various temperatures, (530.4, 505.4 and 498.1) °C,
respectively. The increasing amount of Cu to 4 % (5 %
Si) further changes the shapes of the Cu-enriched phase
peaks (Figure 5). The precipitation temperatures were
also altered. The Cu-enriched phase represented by the
first peak of the Al-Si5-Cu4 alloy begins to precipitate at
514.4 °C, while the second peak appears at 507.2 °C.
The increasing Cu amount from 1 % to 4 % slightly
increased the total solidification time from 1167 s (for
the Al-5Si-1Cu alloy) to 1211 seconds (for the Al-5Si-
4Cu alloy), increasing also the total solidification tempe-
rature interval of the Cu-rich phase(s) from 31.4 °C (for
the Al-5Si-1Cu alloy) to 65.4 °C (for the Al-5Si-4Cu
alloy).
These results of the experiments (Figures 2 to 5)
indicate that the Cu-enriched phases precipitate at diffe-
rent temperatures depending on the amount of Cu pre-
M. B. DJURDJEVIC et al.: QUANTIFICATION OF THE COPPER PHASE(S) IN Al-5Si-(1–4)Cu ALLOYS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 299–304 301
Figure 4: First derivatives of the Al-5Si-(1–4)Cu cooling curves
Slika 4: Prvi odvod ohlajevale krivulje zlitin Al-5Si-(1-4)Cu
Figure 2: Cooling curves of the investigated Al-5Si-(1–4)Cu alloys
Slika 2: Ohlajevalne krivulje preiskovanih zlitin Al-5Si-(1-4)Cu
Figure 5: First derivatives of the Al-Si5-Cu(1–4) cooling curves
related to the Cu-enriched region
Slika 5: Prvi odvod ohlajevalnih krivulj Al-5Si-Cu(1-4) glede na z
bakrom bogato podro~je
Figure 3: Impacts of different Cu amounts on the characteristic tem-
peratures of Al-5Si-(1–4)Cu alloys
Slika 3: Vpliv razli~nih vsebnosti Cu na zna~ilne temperature v zliti-
nah Al-5Si-(1-4)Cu
sent in the particular Al-Si5-Cu(1-4) alloy. The nucle-
ation temperature of the Cu-enriched phases can be
accurately read from the first derivatives of the cooling
curves and used to define the maximum temperatures
that the castings can be exposed to during the conven-
tional solution-treatment process. However, before the
solution-treatment routines can be "tailored" to specific
alloys and applications, it is also necessary that the
volume fractions of the Cu-enriched phases are known.
This data enables the researchers to predict the mecha-
nical properties of the castings and design components
according to the predetermined specifications and
requirements. To date, a volume-fraction assessment has
only been possible through a metallographic analysis.
3.2 Metallography, the cooling curve and image-ana-
lysis results
Light optical microscopy (LOM) observations com-
bined with the IA showed that the area fractions of the
Cu-enriched phases increased with additions of Cu. A
Cu increase from 1 % to 4 % caused the area fraction of
the Cu-enriched phases to increase from about 0.55 % to
about 2.42 % (Table 2).
Table 2: Comparison of the Cu-enriched-phase area fractions detected
by the IA system and determined with the TA
Tabela 2: Primerjava dele`a podro~ij s Cu bogatih faz, ugotovljenih z
IA-sistemom in dolo~enih s TA
Alloy
Area of
Cu-rich phase,
(TA) %
Area of
Cu-rich phase,
(IAS) %
w(Cu)/%
Al-5Si-1Cu 0.90 0.55 1.03
Al-5Si-2Cu 2.55 1.65 2.06
Al-5Si-4Cu 4.30 2.42 3.85
An additional SEM observation, combined with an
X-ray spot microanalysis for the investigated alloy
(Al-5Si-4Cu) was performed to identify the morpholo-
gies and stoichiometries of the observed Cu-enriched
phases. This analysis confirmed the earlier assertion that
Cu-enriched phases appear with three main morpholo-
gies: the blocky type, the eutectic type and the fine
eutectic type (Figure 6). The quantitative X-ray micro-
analysis of the revealed stoichiometries of the Cu phases
(Table 2) is presented in Figure 6.
It should be noted that a complete evaluation of the
morphologies and the corresponding stoichiometries of
the Cu-enriched phases is beyond the scope of the pre-
sent paper. Quenching experiments will be necessary to
establish the crystallization sequences of the Cu-enriched
phases and the corresponding stoichiometries with
respect to the TA results.
The imperfect agreement between these two measu-
rements can be explained with two factors: First, the IA
measurements do not take into account the small Si
crystals that cannot be resolved with the LOM or the Si
that is dissolved in the aluminum matrix. Second,
because the cast samples are heterogeneous and due to
the fact that only a finite number of regions were eva-
luated using the IA, these measurements may not be
representative of all the test samples.
A determination of the total Cu-enriched-phase area
fraction with metallography is a time-consuming and
laborious procedure; therefore, it cannot be used as an
on-line measurement tool, or as a method of controlling
the casting quality in a foundry environment.
The TA approach developed by Kierkus and Soko-
lowski5 was used in this work for determining the area
fractions of individual phases that precipitate during
M. B. DJURDJEVIC et al.: QUANTIFICATION OF THE COPPER PHASE(S) IN Al-5Si-(1–4)Cu ALLOYS ...
302 Materiali in tehnologije / Materials and technology 48 (2014) 2, 299–304
Figure 6: SEM micrographs of the characteristic morphologies of Cu-enriched phases and their EDX elemental maps
Slika 6: SEM-posnetki zna~ilne morfologije s Cu bogatih faz in njihova elementna EDS-analiza
solidification of Al-Si-Cu alloys. In their work, the
integrated area of the Cu-enriched phases is defined as
the ratio of the area between the first derivative (FD) of
the cooling curve and the hypothetical solidification path
of the Al-Si-Cu eutectic (the hatched area in Figure 7) to
the total area between the first derivative of the cooling
curve and the base line (BL). The rationale of this
assumption is based on:5
1. The IA results, which permit one to postulate that the
solidification of the Al-Si eutectic continues until the
solidus temperature is reached.
2. The total latent energy evolved during the alloy
solidification is the sum of the energy released by all
of the phases involved in the process.
This concept is briefly demonstrated in Figures 7 and
8, which present the FD of the cooling curve and the BL
curve. The area between the two curves, from the liqui-
dus state (TLIQ ) to the solidus state (TSOL ), is pro-
portional to the latent heat of the solidification of the
alloy. If the two aforementioned assumptions are correct,
then the regression line between the arbitrarily selected
state (TNUC
Al -Si -Cu) and the solidus state (TSOL ) is a part of
the solidification path of the Al-Si-Cu eutectic (the
hatched area). Therefore, it is evident that the area
between the path (T TNUC
Al -Si -Cu
SOL− ) and the FD of the
cooling curve should be proportional to the latent heat of
the solidification of the Cu-enriched phases. The
proportionality is constant in both cases; the total latent
heat of the alloy solidification and the latent heat of the
solidification associated with the Cu-enriched phases are
the "apparent specific heat" of the alloy.
A comparison of the total area fraction of the
Cu-enriched phases determined using the IA with the
integrated area (the hatched area in Figure 7) of the
Cu-enriched phase of each alloy tested shows that the
two measurements are almost perfectly correlated (Fig-
ure 9).
The imperfect agreement between these two measu-
rements can be explained with two factors: First, the IA
measurements do not take into account the small Si
crystals that cannot be resolved with the LOM or the Si
that is dissolved in the aluminum matrix. Only TEM
investigations under a very high magnification would be
able to reveal the presence of ultra-fine Al-Cu eutectics.
Second, because the cast samples are heterogeneous and
because only a finite number of regions were evaluated
using the IA, these measurements may not be precisely
representative of all the samples.
The results of the Cu-enriched-phase determinations
are presented in Table 2 and in Figure 9. A high corre-
lation observed on the regression plots (Figure 9) shows
that it is possible to estimate the volume fraction of the
Cu-enriched phases from the TA analysis experiments
without resorting to the IA.
4 CONCLUSIONS
A comprehensive understanding of the melt quality is
of a paramount importance for the control and prediction
of actual casting characteristics. The thermal analysis is
M. B. DJURDJEVIC et al.: QUANTIFICATION OF THE COPPER PHASE(S) IN Al-5Si-(1–4)Cu ALLOYS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 299–304 303
Figure 8: Relationship between IA and TA measurements and the
chemical compositions of the investigated alloys
Slika 8: Odvisnost med IA- in TA-meritvami ter kemijsko sestavo pre-
iskovanih zlitin
Figure 9: Relationship between IA and TA measurements and the
chemical compositions of the investigated alloys
Slika 9: Odvisnost med IA- in TA-meritvami ter kemijsko sestavo
preiskovanih zlitin
Figure 7: Part of the first-derivative curve (FD) related to the Cu-rich
phase5
Slika 7: Del prvega odvoda krivulje (FD) glede na s Cu bogate faze5
an already used tool for the melt-quality control in an
aluminum casting plant. It has been used routinely for
assessing the master-alloy additions to an aluminum
melt. In addition, its application can be extended to
quantify the total volume fraction of the Cu-enriched
phases of the Al-Si-Cu aluminum alloys. Future work
should confirm that an on-line quantitative control of the
Cu-enriched phases is also possible for the other series
of Al-Si alloys using TA.
5 REFERENCES
1 L. Bäckerud, G. Chai, J. Tamminen, Solidification Characteristics of
Aluminum Alloys, Vol. 2: Foundry Alloys, AFS/ScanAluminum,
Oslo 1990
2 C. H. Caceres, M. B. Djurdjevic, T. J. Stockwell, J. H. Sokolowski,
The effect of Cu content on the level of microporosity in
Al-Si-Cu-Mg Casting Alloys, Scripta Materialia, 40 (1999), 631–637
3 M. B. Djurdjevic, T. Stockwell, J. Sokolowski, The effect of stron-
tium on the microstructure of the Al-Si and Al-Cu eutectics in the
319 aluminum alloy, International Journal of Cast Metals Research,
12 (1999), 67–73
4 H. W. Doty, A. M. Samuel, F. H. Samuel, Factors controlling the
type and morphology of Cu-Containing phases in the 319 aluminum
alloy, 100th AFS Casting Congress, Philadelphia, Pennsylvania,
USA, 1996, 1–30
5 W. T. Kierkus, J. H. Sokolowski, Recent advances in cooling curve
analysis: A new method of determining the "base line" equation,
AFS Transactions, 66 (1999), 161–167
6 M. B. Djurdjevic, W. Kasprzak, C. A. Kierkus, W. T. Kierkus, J. H.
Sokolowski, Quantification of Cu enriched phases in synthetic 3XX
aluminum alloys using the thermal analysis technique, AFS Transac-
tions, 24 (2001), 1–8
M. B. DJURDJEVIC et al.: QUANTIFICATION OF THE COPPER PHASE(S) IN Al-5Si-(1–4)Cu ALLOYS ...
304 Materiali in tehnologije / Materials and technology 48 (2014) 2, 299–304
G. AKPÝNAR et al.: INVESTIGATION OF INDUCTION AND CLASSICAL-SINTERING EFFECTS ...
INVESTIGATION OF INDUCTION AND
CLASSICAL-SINTERING EFFECTS ON POWDER-METAL
PARTS WITH THE FINITE-ELEMENT METHOD
PRIMERJAVA VPLIVA INDUKCIJSKEGA IN
KONVENCIONALNEGA SINTRANJA NA DELCE KOVINSKEGA
PRAHU Z UPORABO METODE KON^NIH ELEMENTOV
Göksan Akpýnar, Can Çivi, Enver Atik
Celal Bayar University, Engineering Faculty, Mechanical Engineering Department, 45040 Manisa, Turkey
goksanakpinar105@hotmail.com
Prejem rokopisa – received: 2013-04-29; sprejem za objavo – accepted for publication: 2013-06-13
Induction sintering provides large time and energy savings because the components heat up rapidly and the sintering time is
lower than in classical sintering in a furnace. Therefore, induction sintering is an important alternative to classical sintering. In
this study, mechanical properties of induction-sintered Fe-based components including Cu and carbon (graphite) were compared
with those sintered in a classical furnace. For this purpose, microstructure photographs of both samples were taken. A tensile
analysis of the sintered powder-metal samples was carried out with the finite-element method, and the micro-stress values were
found to change depending on the amount and distribution of the porosity.
Keywords: powder metallurgy, sintering, induction sintering, classical furnace, microstructure analysis, finite-element method
Indukcijsko sintranje omogo~a velike prihranke pri ~asu in energiji, saj se komponente ogrejejo hitro in je ~as sintranja kraj{i,
kot pri klasi~nem sintranju v pe~eh. Zato je indukcijsko sintranje pomembna alternativa klasi~nemu sintranju. V tej {tudiji so
bile primerjane mehanske lastnosti indukcijsko sintrane komponente z Fe-osnovo in dodatki Cu ter grafita s komponentami,
sintranimi v klasi~nih pe~eh. V ta namen so bili napravljeni posnetki mikrostrukture obeh vzorcev. Izvr{ena je bila analiza
nateznih preizkusov sintranih kovinskih vzorcev z metodo kon~nih elementov. Ugotovljeno je bilo, da so vrednosti
mikronapetosti odvisne od koli~ine in porazdelitve poroznosti.
Klju~ne besede: metalurgija prahov, sintranje, indukcijsko sintranje, klasi~na pe~, analiza mikrostrukture, metoda kon~nih
elementov
1 INTRODUCTION
Powders with different compositions are pressed and
then sintered with the powder-metallurgy (P/M) method.
Sintering is one of the most important issues of powder
metallurgy because it causes a significant increase in the
strength of the pressed powders. The sintering process is
generally performed in sintering furnaces. It is done in a
protective atmosphere of batch or continuous furnaces.1
In addition, rapid sintering methods such as induction
sintering, microwave sintering, plasma sintering, laser
sintering and discharge sintering are important alter-
natives to conventional sintering methods.2 Sintering and
additional heat treatments of powder mixtures cause the
microstructure to meet the performance requirements.3
Mixtures of elemental iron and graphite powders are
commonly used for P/M applications. A small amount of
copper powder is always added to further strengthen the
sintered alloys owing to its relative ease of dissolving
and diffusing in an iron matrix upon sintering.4
Almost all low-alloy steel powders contain copper.
The mass fractions of copper varies between approxima-
tely 1 % and 8 % depending upon the desirability of end
products. A small amount of copper is added to provide
strength by age hardening, while the purpose of higher
concentrations is to promote liquid-phase sintering
causing a faster densification and homogenization.5 An
addition of carbon to iron powder increases the sintering
kinetics as it dissolves into the iron lattice, changing the
melting point, surface tension and viscosity of the iron
melt formed. Small areas of martensite and tempered
martensite are also formed. 6
The most important feature of the induction-heating
system is a rapid heating of the material because heating
occurs directly on the metal parts. In general, induction
sintering is used for surface heating of materials.7 If the
frequency increases, eddy currents will occur on the
region close to the surface.8 The heat transfer is 3.000
times better than in the other heating systems.7 This
allows a much faster completion of the warm-up process,
reducing the time spent for this period and, thus,
shortening the sintering time.
In addition, sintered ferrous P/M components have
emerged as attractive candidates to replace wrought
alloys in many applications due to their low cost, high
performance and the ability to be processed to the near-
net shape. Sintered materials are typically characterized
by the residual porosity after sintering, which is quite
detrimental to the mechanical properties of these mate-
rials.9–17 The nature of the porosity is controlled with
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several processing variables such as green density,
sintering temperature and time, alloying additions, and
the particle size of the initial powders.10 In particular, the
fraction, size, distribution and morphology of the
porosity have a profound impact on mechanical behavior.
Alloying elements such as copper, nickel and graphite
affect the sintering parameters leading to the formation
of a heterogeneous internal structure. Thus, the hetero-
geneous nature of the microstructures of P/M steels will
certainly play a role in the onset and evolution of damage
under an applied stress.11–17 Under a monotonic tensile
loading, the porosity reduces the effective load-bearing
cross-sectional area acting as a stress-concentration site
for the strain localization and damage, decreasing both
strength and ductility.11 Interconnected porosity causes
an increase in the localization of the strain on the relati-
vely smaller sintered regions between the particles, while
isolated porosity results in a more homogeneous defor-
mation. It is also not uncommon for the porosity distri-
bution in a material to be inhomogeneous. In this case,
the strain localization will take place at the "pore
clusters". Thus, for a given amount of porosity, the
interconnected porosity is more detrimental, reducing the
macroscopic ductility to a greater extent than the isolated
porosity.17
Porosity affects the mechanical properties of mate-
rials. Many studies have been conducted on this topic. N.
Chawla and X. Deng17 investigated the effects of mecha-
nical properties, the shape and size factors of the poro-
sity of sintered Fe–0.85Mo–Ni steels. They systemati-
cally examined the effect of porosity on the tensile and
fatigue behaviors of the Fe–Mo–Ni steel. The steels of
three densities were studied: 7.0 g/cm3, 7.4 g/cm3 and 7.5
g/cm3. A quantitative analysis of the microstructure was
performed to determine the pore-size distribution and the
pore shape as functions of the sintered density. Holmes
and Queeney18 proposed that the relatively high stress
concentration at pores, particularly the surface pores, is
responsible for the localized slip leading to a crack initi-
ation. Christian and German19 showed that the fraction of
porosity, pore size, pore shape and pore spacing are all
important factors controlling the fatigue behavior of P/M
materials. In general, more irregular pores exhibit a
higher stress than perfectly round pores.10 Polasik et al.15
showed that small cracks nucleate from the pores during
the fatigue and coalesce to form a larger crack leading to
a fatigue fracture. Here, the heterogeneous nature of the
microstructure played an important role by contributing
to the crack tortuosity. Crack arrest and crack deflection
were observed due to microstructural barriers such as
particle boundaries, fine pearlite, and nickel-rich
regions.15
In this study, the microstructures of classically sin-
tered and induction-sintered metal-powder parts with a
medium/low frequency (30 kHz) obtained with the expe-
rimental studies were compared. The effects of the sin-
tering time on the mechanical properties were identified
with image processing and the finite-element method.
The micro-stresses around the internal spaces in the
microstructures were investigated.
2 MATERIALS AND METHODS
In this study, the Högenas ASC 100.29 iron powder
(2 % Cu, 0.5 % graphite and 1 % Zn Stereat lubricant by
mass) was used. Powder-metal bushings were produced
by Toz Metal Inc. with a dual-axis press under a 600
MPa pressure. The sieve analysis of the iron powder is
shown in Table 1.20 Induction and classical sintering
mechanism are indicated in Figure 1. Powder-metal
bushings with the dimensions of (16/14 mm × 36 mm
are shown in Figure 2.
Table 1: Sieve analysis of the metal powder20
Tabela 1: Sejalna analiza kovinskega prahu20
Iron powder Sieve analysis (%)
ASC 100.29
< 45
μm
45–150
μm
150–180
μm
> 180
μm
23 69 8 0
The powder-metal bushings were sintered in an
environment atmosphere in an electric-resistance furnace
G. AKPÝNAR et al.: INVESTIGATION OF INDUCTION AND CLASSICAL-SINTERING EFFECTS ...
306 Materiali in tehnologije / Materials and technology 48 (2014) 2, 305–312
Figure 1: a) Induction-sintering mechanism, b) classical resistance furnace
Slika 1: a) Naprava za indukcijsko sintranje, b) klasi~na uporovna pe~
for 30 min at 1120 °C and they were also sintered by
induction sintering for 8.4 min and 15 min at 1120 °C in
the environment atmosphere. Induction sintering was
carried out in a heat-resistant glass in a copper coil. The
microstructural mechanical properties of these samples
sintered for different periods and in different furnaces are
compared with each other.
The induction sintering was carried out in a heat-resi-
stant glass in a 36 mm diameter copper coil. The con-
veyor belt system is suitable for a mass production. The
sintering temperature of 1120 °C was recorded on a
pyrometer with laser and it was kept constant with the
induction-mechanism unit. The sintered powder-metal
bushings were cut and the microstructure images of the
specimens were investigated using a Nikon Eclipse
LV100 microscope. The cross-sections of the steel
specimens were ground, polished and etched with a 2 %
Nital solution (2 % HNO3 and 98 % alcohol). The ima-
ges of the polished surfaces of the cross-sections were
taken.
The microstructures of the samples (40 μm) were
processed by image processing. A tensile stress was
applied to the samples with the finite-element method
and the micro-stress values were obtained for the
samples. The solution was made with an adoption of the
mechanical properties of the steel containing 0.6 %
graphite.
3 RESULTS
3.1 Microstructural analysis
A microstructural investigation was applied to the
sintered bushings after polishing the surface with alumi-
na and acid etching with a 3 % Nital solution. The
microstructural photos are shown in Figures 3 to 5.
G. AKPÝNAR et al.: INVESTIGATION OF INDUCTION AND CLASSICAL-SINTERING EFFECTS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 305–312 307
Figure 5: Microstructure of an induction-sintered powder-metal
bushing (sintered at 1120 °C for 15 min), light microscope (LM), a
100-times magnification
Slika 5: Mikrostruktura indukcijsko sintrane kovinske pu{e (sintrano
pri 1120 °C za 15 min), svetlobni mikroskop, pove~ava 100-krat
Figure 3: Microstructure of a classically sintered powder-metal
bushing (1120 °C/30 min in the furnace), light microscope (LM), a
100-times magnification
Slika 3: Mikrostruktura klasi~no sintrane kovinske pu{e (1120 °C/30
min v pe~i), svetlobni mikroskop, pove~ava 100-krat
Figure 4: Microstructure of an induction-sintered powder-metal
bushing (sintered at 1120 °C for 8.4 min), light microscope (LM), a
100-times magnification
Slika 4: Mikrostruktura indukcijsko sintrane pu{e iz kovinskega prahu
(sintrano pri 1120 °C za 8,4 min), svetlobni mikroskop, pove~ava
100-krat
Figure 2: Sintered bushings
Slika 2: Sintrane pu{e
3.2 Image processing and FEM Analyses of micro-
structure pictures
The mechanical properties of the materials are shown
in Table 2. The porosity values obtained from the image
analysis of the samples are shown in Table 3. The maxi-
mum and minimum micro-stresses of the static tensile
strength acting horizontally (1 direction) on the internal
pores were found and compared with each other. The
results of the study using the finite-element method are
shown in Figures 6 to 10.
Table 2: Mechanical properties of the iron-based sintered material
with 0.6 % graphite added and the loading conditions of the samples
Tabela 2: Mehanske lastnosti sintranega materiala z dodatkom 0,6 %
grafita in razmere pri obremenitvi vzorcev
Poisson’s ratio
(an approximation)
Thermal-expansion
coefficient (1/K)
Fx – Edge load,
1-direction (N/m²)
0.3 11.8 E–6 20 E6
4 DISCUSSION
It is well known that porosity decreases the Young’s
modulus of a material.10 We use the approach of Rama-
krishnan and Arunachalam (R–A)21 to model the effect
of the porosity on the Young’s modulus. The Young’s
modulus of a material, E, with a given fraction of
porosity, p, is given by:
E = E0 [(1 – p)
2 / (1 + )Ep)] (1)
where E0 is the Young’s modulus of a fully dense steel
(obtained by extrapolating the experimental data to the
zero porosity, yielding a value of approximately 200
GPa), and )E is the constant in terms of the Poisson’s
ratio of a fully dense material, v0:
)E = 2 – 3v0 (2)
For a fully dense steel, the Poisson’s ratio is approxi-
mately 0.3. This is supported by the analysis of Rama-
krishnan and Arunachalam,21 who compared the bulk
modulus of porous materials with the spherical-versus-
angular-pore geometry using FEM. An analytical solu-
tion was made to show that, depending on the density
and porosity, the samples of the microstructures were
G. AKPÝNAR et al.: INVESTIGATION OF INDUCTION AND CLASSICAL-SINTERING EFFECTS ...
308 Materiali in tehnologije / Materials and technology 48 (2014) 2, 305–312
Figure 7: a) Microstructure of a powder-metal bushing classically sin-
tered for 30 min in the furnace, a finite-element model of microstruc-
ture images with normal stress (MPa), b) microstructure of a powder-
metal bushing induction sintered for 8.4 min, a finite-element model
of microstructure images with normal stress, c) microstructure of a
powder-metal bushing induction sintered for 15 min, a finite-element
model of microstructure images with normal stress
Slika 7: a) Mikrostruktura klasi~no sintrane kovinske pu{e 30 min v
pe~i; model mikrostrukture z metodo kon~nih elementov z normalno
napetostjo (MPa), b) mikrostruktura indukcijsko sintrane kovinske
pu{e (8,4 min); model mikrostrukture z metodo kon~nih elementov z
normalno napetostjo, c) mikrostruktura indukcijsko sintrane kovinske
pu{e (15 min); model mikrostrukture z metodo kon~nih elementov z
normalno napetostjo
Figure 6: Finite-element boundary conditions of the real-microstruc-
ture image of a powder-metal part
Slika 6: Robni pogoji za analizo z metodo kon~nih elementov na
realni mikrostrukturi sintranega kovinskega dela
Figure 8: Induction-sintered sample 15 min, the maximum stress
(MPa) in the area of the porosity of the microstructure
Slika 8: Indukcijsko sintran vzorec 15 min, najve~ja napetost (MPa)
na obmo~ju poroznosti v mikrostrukturi
affected by micro-stresses. The Young’s modulus values
of the samples are given in Table 3.
Although the finite-element analysis was used to
study the mechanical behaviors of powder-metallurgy
materials,16,22–24 the pores are generally modeled as
perfect spheres. But, at critical values of strain the
imperfections cause localization of plastic flow.24 In the
R-A model a single spherical pore is surrounded by a
spherical matrix shell, causing an intensification of the
pressure on the pore surface due to the interaction of the
pores in the material.25 The material behavior is con-
trolled by the microstructure of steel, in particular, the
nature of the porosity.17
In this study we used two-dimensional microstruc-
tures as the basis for the finite-element simulations of the
samples induction sintered for 8.4 minutes and 15 min,
and the samples classically sintered for 30 min in a fur-
nace. Figure 6 shows the actual microstructure version
of the uniaxial loading, boundary conditions and the
mesh. A quadratic triangular mesh modeling was deemed
appropriately. A finer mesh was used in the regions of
pore clusters. In order to yield accurate simulation
results, we used an entire picture of the microstructure
simulation. The 2D analysis presented here shows the
qualitative effects of the pore microstructure on the
localized plastic strain and stress initiation around the
pores. The porosity values of the samples, the maximum
and minimum normal stresses, the Young’s modulus, the
total displacement, thermal-expansion coefficient values,
G. AKPÝNAR et al.: INVESTIGATION OF INDUCTION AND CLASSICAL-SINTERING EFFECTS ...
Materiali in tehnologije / Materials and technology 48 (2014) 2, 305–312 309
Figure 9: Finite-element investigation of the microstructures of the samples with both deformed and undeformed shapes: a) powder-metal
bushing classically sintered for 30 min in the furnace, b) powder-metal bushing induction sintered for 8.4 min, c) powder-metal bushing induction
sintered for 15 min, d) bulk material
Slika 9: Preiskava mikrostrukture z metodo kon~nih elementov vzorcev v deformiranem in nedeformiranem stanju: a) klasi~no sintrana kovinska
pu{a (30 min v pe~i), b) indukcijsko sintrana kovinska pu{a (8,4 min), c) indukcija sintrana kovinska pu{a (15 min) in d) osnovni material
Table 3: Porosity, stresses and total-displacement values obtained from the image-processing analysis
Tabela 3: Poroznost, napetosti in skupen pomik, dobljeni iz analize slik
Samples
Porosity from
image analysis
(%)
Density
(kg/m³)
Maximum
stress around a
pore (MPa)
Minimum
stress around a
pore (MPa)
Total
displacement
(μm)
Young’s
modulus (GPa)
Average of the samples
induction sintered for 8.4 min 3.5411 7270 794.128 –234.780 2.31 E–2 185.684
Average of the samples
induction sintered for 15 min 3.1846 7352 581.068 –473.677 3.186 E–2 187.077
Average of the samples
classically sintered for 30 min 2.4004 7474 708.883 –228.570 2.774 E–2 190.177
Bulk sample 0 7860 34.171 17.967 1.691 E–2 200.000
the Poisson’s ratio and the edge 1-direction loads are
shown in Tables 2 and 3.
The normal stress around the pores, the total displa-
cement, the deformed and undeformed shapes of the
microstructures of the samples are shown in Figures 7
and 9. These figures show that the porosity was caused
by an inhomogeneous deformation. The modeling also
shows that the plastic-strain intensification begins at the
tips of the irregular pores in the microstructure. This
means that the more irregular the porosity, the more
damage can be seen in the microstructure. For a regular
deformation, as well as the porosity shape, the distribu-
tion of porosity in a structure is also important. Vedula
and Heckel9 investigated the mechanisms of a damage of
flat and angular pores in a microstructure. They observed
the forming of local shear bands at the ends of the pores,
creating unstable tensions around the angular pores.
A mesh view and the finite-element boundary con-
ditions of a real-microstructure image of a powder-metal
part is shown in Figure 6. The solutions for each sample
were compared by applying the same boundary condi-
tions. Also, the local plastic strain acting around the
micropores was found to be a result of the two-dimen-
sional analysis. Each sample of the tensile surface is
taken to have a value of F1 = 20 N/m2 for the stress-edge
1-direction load.
The stress equation of the modeling is as follows:25
F(ij – *ij) – h(h) = 0 (3)
where ij is the symmetric stress index, *ij is the first
cycle of the yield surface, lh is the scalar function of the
plastic strain and h(h) refers to the amount of expan-
sion of the yield surface.
The stress concentrations at the tips of irregular pores
in the microstructure are shown in Figure 7. Also, for
the area around the pores of the sintered samples, the
normal-stress FE-analysis results are given.
The maximum stress around the pores of the sample
induction sintered for 8.4 min was 794.1 MPa. For the
sample induction sintered for 15 min, a relatively lower
value of 581.5 MPa for the maximum stress was
obtained. On the basis of these results, it can be con-
cluded that the maximum-stress value around the pores
decreases with an increase in the density. As you can see
in Figure 8, the maximum stress in the area of porosity
depends not only on the density but also on the pore
shape. The maximum stress for the microstructure of the
sample induction sintered for 15 min is also shown in
Figure 8. However, we have a difficulty here: although
the sample induction sintered for 15 min showed the
lowest tensile stress, the maximum stress value in the
opposite direction is –473.7 MPa, which is higher than
the other values. This result shows that the pore shape of
a microstructure plays a key role in the micro-tensile
stress.
The non-deformed and deformed shapes of the sam-
ples were analyzed. Despite having the lowest density,
the induction-sintered sample 8.4 min showed a more
uniform deformation than the other porosity samples.
According to these results, smaller and regular pores
contribute to a uniform deformation. When comparing
the amounts of deformation in the porosity samples and
the bulk sample, the porosity samples are found to be
more deformed than the bulk sample. It can be said that
the deformation amount of the samples depends on the
pore shape as well as on the density.
The strain curves of the tensile surfaces of the sam-
ples are given in Figure 10. These strain curves appear
to be quite different. The reasons for this difference are
the rate, the shape and the pore density of the samples.
As expected, the bulk sample has a symmetric defor-
mation curve. When the inner tensile surface of the pores
G. AKPÝNAR et al.: INVESTIGATION OF INDUCTION AND CLASSICAL-SINTERING EFFECTS ...
310 Materiali in tehnologije / Materials and technology 48 (2014) 2, 305–312
Figure 10: Total displacement curves of the tensile surfaces: a) pow-
der-metal bushing classically sintered for 30 min in the furnace, b)
powder-metal bushing induction sintered for 8.4 min, c) powder-metal
bushing induction sintered for 15 min, d) bulk material
Slika 10: Skupni premik natezno obremenjene povr{ine: a) klasi~no
sintrana kovinska pu{a (30 min v pe~i), b) indukcijsko sintrana kovin-
ska pu{a (8,4 min), c) indukcija sintrana kovinska pu{a (15 min) in
d) osnovni material
increases, the surface-deformation-curve peak increases
as well.
5 CONCLUSIONS
In this study, Fe-based powder-metal bushings were
sintered with the classical-furnace and induction-sinter-
ing mechanisms. The microstructures of the classically
sintered and induction sintered powder-metal bushings
with a low to medium frequency (30 kHz) were com-
pared. The effects of the sintering time on the mecha-
nical properties were investigated with image processing
and the finite-element method. The results show the
following:
• The stresses that occurred around the pores in the
microstructures of the samples were investigated
numerically, showing how the stresses and displace-
ment of the pores related to the sintering methods and
parameter changes. Besides, it was also found that
the mechanical properties of porous materials and the
bulk material are quite different.
• Numerical results are shown in the Table 3. The
maximum and minimum stresses for the samples
classically sintered for 30 min are 708.9 MPa and
–228.6 MPa, respectively. The maximum and mini-
mum stresses for the sample induction sintered for
8.4 min are 794.1 MPa and –234.8 MPa, respectively.
The maximum and minimum stresses for the samples
induction sintered for 15 min are 581.5 MPa and
–473.7 MPa, respectively, while the maximum and
minimum stresses for the bulk samples are 34.2 MPa
and 18 MPa, respectively.
• The pore sizes decrease with the increasing sintering
time as illustrated in Figure 7. With the increasing
induction-sintering time, large pores become relati-
vely small, small pores disappear and the sintered
density increases. On the other hand, when the induc-
tion-sintering time in the microstructures increases,
lower and more homogenized tensile stresses occur
around the pores.
• When looking at the values for the porosity obtained
with the image analysis shown in Table 3, the mini-
mum porosity value (2.4 %) is found for the samples
classically sintered in the furnace for 30 min and, as
expected, this porosity causes a smaller displacement
than the other porosities. It is seen that the sample
induction sintered for 8.4 min has a smaller displace-
ment than the one induction sintered for 15 min. A
more regular deformation is also shown in Table 3
and Figure 8. According to this result, smaller and
more regular pores of the sample induction sintered
for 8.4 min are thought to cause a more regular defor-
mation.
• The porosity samples were also compared to the bulk
sample. It was seen that considerable internal stresses
were formed around the pores of the porosity sam-
ples. This means that the material is exposed to ten-
sion, and the micro-stress concentrating around the
pores can cause damage in a much shorter time. As a
precaution, smaller and more regular pores should be
formed in the microstructure and the microstructures
of the materials should be concentrated.
• The microstructure-based FEM modeling showed
that smaller, more regular and more clustered pores
cause a more regular displacement and a reasonable
micro-stress. So, the micro-stress and micro-strain
depend on the pore shape and the loading condition
as well as on the pore density.
• In our previous study, it was found that the strength
values of the samples sintered with induction were
increased by increasing the sintering time.26 Due to
more uniform and smaller pores in our current study,
the micro-stress values of the sintered samples
decreased. This also proves that the strength of the
samples increases with a decrease in the micro-stress
values.
• Another aim of this study was to investigate how the
pores affect the micro-stress and micro-deformation.
It was found that the strength of a porous material
depends on the shape, the size and the density of the
pores.
Acknowledgments
We would like to thank Toz Metal Inc. and Mr. Aytaç
Ataº for providing the metal powder and pressing the
powder-metal bushings.
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