VSEBINA – CONTENTS
Predgovor/Foreword
M. Torkar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 319
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Numerical and experimental investigation of the temperature field of a solidifying massive ductile-cast-iron roller
Numeri~ne in eksperimentalne raziskave temperaturnega polja pri strjevanju litega valja iz sive litine s kroglastim grafitom
F. Kavi~ka, K. Stránský, J. Dobrovská, B. Sekanina, J. [tìtina, V. Gontarev . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 321
A fuzzy-based optimal control algorithm for a continuous casting process
Algoritem, ki temelji na mehki logiki, za optimalni nadzor kontinuirnega procesa ulivanja
T. Mauder, C. Sandera, J. Stetina. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 325
Thermal stability of Al-Mn-Be melt-spun ribbons
Temperaturna obstojnost hitro strjenih trakov Al-Mn-Be
G. Lojen, T. Bon~ina, F. Zupani~. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 329
Challenges in the computer modeling of phase change materials
Izzivi v ra~unalni{kem modeliranju materialov s fazno spremembo
L. Klimes, P. Charvat, M. Ostry. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 335
Application of powder metallurgy in the processing of aluminium scraps with high-iron contents
Uporaba postopkov pra{ne metalurgije za predelavo odpadkov na osnovi Al z visoko vsebnostjo Fe
D. Vojtìch, F. Prù{a. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 339
Fatigue-life behaviour and a lifetime assessment of a double-leaf spring using FEM-based software
Utrujanje in ocena dobe trajanja dvolistnatih vzmeti z uporabo MKE-orodja
P. Borkovi}, B. [u{tar{i~, M. Male{evi}, B. @u`ek, B. Podgornik, V. Leskov{ek. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 345
Characterization of defects in PVD TiAlN hard coatings
Karakterizacija defektov PVD TiAlN trdih prevlek
P. Gselman, T. Bon~ina, F. Zupani~, P. Panjan, D. Kek Merl, M. ^ekada . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 351
Digital imaging analysis of microstructures as a tool to identify local plastic deformation
Digitalna analiza posnetkov mikrostruktur kot orodje za ugotavljanje lokalne plasti~ne deformacije
M. Suban, R. Cvelbar, B. Bundara. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 355
Thermodynamic analysis of the formation of non-metallic inclusions during the production of C45 steel
Termodinami~na analiza nastanka nekovinskih vklju~kov pri izdelavi jekla C45
L. Krajnc, G. Klan~nik, P. Mrvar, J. Medved . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 361
Small-angle x-ray scattering spectra of iron-based magnetic fluids
Malokotno sipanje rentgenskega spektra magnetnih teko~in na osnovi `eleza
S. Rugmai, C. Sirisathitkul, K. Chokprasombat, P. Rangsanga, P. Harding, T. Srikhirin, P. Jantaratana . . . . . . . . . . . . . . . . . . . . . . . . . . 369
Microstructure of metal-matrix composites reinforced by ceramic microballoons
Mikrostruktura kompozitov s kovinsko osnovo, oja~ano s kerami~nimi mikrokroglicami
I. N. Orbulov, K. Májlinger . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 375
Optimization of multi-process parameters according to the surface quality criteria in the end milling of the AA6013 aluminum
alloy
Optimizacija multiprocesnih parametrov v odvisnosti od kakovosti povr{ine kot merila pri kon~nem rezkanju aluminijeve zlitine AA6013
H. Durmuþ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 383
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Numerical and experimental analyses of the chemical heterogeneity of a solidifying heavy ductile-cast-iron roller
Numeri~na in eksperimentalna analiza kemijske heterogenosti strjevanja litega te`ko gnetljivega `eleznega valja
K. Stransky, F. Kavicka, B. Sekanina, J. Dobrovska, V. Gontarev . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 389
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 46(4)317–427(2012)
MATER.
TEHNOL.
LETNIK
VOLUME 46
[TEV.
NO. 4
STR.
P. 317–427
LJUBLJANA
SLOVENIJA
JULY–AUG.
2012
Investigation on the aging behaviour of the functionally gradient material consisting of boron carbide and an aluminum alloy
Raziskava pona{anja pri staranju funkcionalnih gradientnih materialov iz borovega karbida in aluminijeve zlitine
B. Sarýkan, E. Balcý, M. Übeyli, N. Camuºcu. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 393
Setting a numerical simulation of filling and solidification of heavy steel ingots based on real casting conditions
Postavitev numeri~ne simulacije polnjenja in strjevanja velikih jeklenih ingotov na podlagi realnih razmer pri ulivanju
M. Tkadleckova, P. Machovcak, K. Gryc, P. Klus, K. Michalek, L. Socha, M. Kovac . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 399
A study of the high-temperature interaction between synthetic slags and steel
[tudij visokotemperaturne interakcije med sinteti~no `lindro in jeklom
K. Gryc, K. Stránský, K. Michalek, Z. Winkler, J. Morávka, M. Tkadle~ková, L. Socha, J. Ba`an, J. Dobrovská, S. Zlá . . . . . . . . . . . . . 403
Development of a model for the internet portal "strength of materials"
Razvoj modela za internetni portal "trdnost materialov"
L. Globa, R. Novogrudska, I. Mamuzi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 407
The influence of the heat-treatment regime on a fracture surface of nickel-based supperalloys
Vpliv toplotne obdelave na povr{ino preloma superzlitin na osnovi niklja
A. Milosavljevic, S. Petronic, S. Polic-Radovanovic, J. Babic, D. Bajic . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 411
Upogibna trdnost korundne keramike: Primerjava razli~nih teoreti~nih porazdelitev na osnovi eksperimentalnih podatkov
Bend strength of alumina ceramics: A comparison of different theoretical distributions on the basis of experimental data
M. Ambro`i~, L. Gorjan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 419
Crevice corrosion of stainless-steel fastening components in an indoor marine-water basin
[pranjska korozija pritrdilnih komponent iz nerjavnega jekla v notranjem bazenu z morsko vodo
M. Torkar, F. Tehovnik, M. Godec. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 423
Predgovor urednika
Opa`amo, da revija Materiali in tehnologije postaja
cenjena v razli~nih strokovnih krogih, saj omogo~a
vpogled v dogajanje na podro~ju kovinskih, polimernih
in anorganskih materialov ter tehnologij s teh podro~ij.
Pove~uje se {tevilo slovenskih avtorjev, pove~uje pa se
tudi interes avtorjev iz tujine. To je dokaz, da gre razvoj
revije v pravo smer, vklju~enost v sistem SCI pa reviji {e
pove~uje dodano vrednost. Revija je pomembna tudi iz
nacionalnega stali{~a. Ve~ina znanstvenih in strokovnih
~lankov je v angle{kem jeziku, v slovenski jezik pa so
prevedeni naslov, povzetek, klju~ne besede, naslovi tabel
in napisi pod slikami. To prispeva k {irjenju slovenske
strokovne terminologije, kar je pomembno za bogatenje
in nadaljnji razvoj slovenskega tehni{kega jezika.
Podobno kot pri drugih revijah, se tudi v reviji
Materiali in tehnologije pojavljajo tematske {tevilke.
Ena takih je bila tudi prva {tevilka v letu 2012, ki je
povzro~ila v dolo~enem krogu bralcev nekoliko
nepri~akovan odziv. Podrobnej{a analiza je pokazala, da
je bila ta {tevilka verjetno izrabljena tudi za pove~anje
samopromocije skupine raziskovalcev, kar je lahko
sporno s stali{~a etike objavljanja ~lankov.
Sedanji sistem ocenjevanja raziskovalcev zelo
upo{teva citiranost kot temeljno merilo kakovosti
raziskovalnega dela in posledi~no tudi financiranja
projektov, premalo pa upo{teva pomembnost dose`kov
za okolje, ki delo financira. Uravnote`eno upo{tevanje
obeh meril bi morala biti, poleg izvirnosti ciljev
raziskovanja, zelo pomembna komponenta ocenjevanja
raziskovalnih projektov, ki so predlo`eni v financiranje
iz javnih sredstev, in kakovosti raziskovalnih skupin.
Prevelika te`a {tevila objav in citatov spodbuja pisce
rokopisov k pridobivanju ~im ve~ citatov, pri ~emer ni
dovolj upo{tevana etika objavljanja. Ve~ina avtorjev
uporablja citiranje in samocitiranje v razumnih mejah.
Mogo~e pa so tudi zlorabe, ki jih uredni{tvo revije MiT
ne more odkrivati, jih pa lahko odkrijejo dobri pozna-
valci tematike, predstavljene v ~lanku, in {e posebej
recenzenti. Prav bi bilo, da bi v takih primerih na odmike
opozorili uredni{tvo MiT, da bi lahko ukrepalo, ko bodo
isti avtorji predlo`ili v objavo nove rokopise.
Kot odgovorni urednik revije menim, da je neeti~no
do avtorjev kakr{no koli {tevil~no omejevanje avto-
citatov in citatov, zato pa naj avtorji poka`ejo primerno
samokriti~nost in naj avtocitate in citate uporabljajo v
razumnem obsegu, kot se to dela pri drugih uglednih
revijah. Pri tem pa morajo ustrezno vlogo odigrati tudi
Foreword of the Editor-in-Chief
The journal Materials and Technology is becoming
increasingly appreciated in various professional circles
because it offers an insight into the metals, polymers and
inorganic materials and technologies used in these fields.
The number of Slovenian authors is increasing, but we are
also seeing an increase in the level of interest from authors
in other countries. This is proof that the journal is
developing in the right direction, and our involvement in
the SCI system increases the journal's added value. The
journal is also important from a national point of view.
Most of the scientific articles are in English, while the
title, abstract, keywords, table headings and figure
captions are translated into Slovene. This contributes to
the spread of Slovenian technical terminology, which is
important for the enrichment and further development of
Slovene.
As is the practice with other journals, the journal
Materials and Technology publishes thematic issues. One
of these was the first issue in 2012, which resulted in a
reaction from a certain circle of readers. A detailed
analysis showed that this thematic issue was used by a
group of scientists for the purposes of self-promotion,
which may have been wrong from the point of view of
ethics in publishing. The current evaluation system
encourages researchers' citations, but the approach of
different authors may vary. Most authors use citations
within reasonable limits, but there are also exceptions.
The existing system for evaluating researchers takes
into account citations as a fundamental measure of the
quality of research work and, consequently, the funding of
projects, but too little emphasis is given to the importance
of the outcome for the environment that funded the work.
A balanced consideration of the two criteria should be, in
addition to the originality of the research goals, a very
important component in the evaluation of research
projects submitted for financing from public funds and for
the evaluation of the quality of research groups. Too much
weight given to the total number of citations encourages
writers of manuscripts to obtain as many references as
possible and not to take sufficient account of publishing
ethics. Most authors use self-citations and citations within
reasonable limits. However, there are also abuses that
cannot be detected by the Editorial Board of the MIT,
although they can be detected by people familiar with the
subject of the article, and especially by referees. In such
cases the problem would be pointed out to the Editorial
Board, to act upon it when the same authors present new
manuscripts for publication.
recenzenti, ki naj bodo v prihodnje pri pregledu in oceni
rokopisov bolj pozorni tudi na ta vidik.
Nadaljnja rast kvalitete revije mora biti cilj nam
vsem, ki smo s to revijo povezani `e vrsto let, tudi ~lanov
uredni{kega odbora in sveta.
Glavni in odgovorni urednik
doc. dr. Matja` Torkar
As the responsible Editor-in-Chief I consider it
unethical to impose a numerical limitation on authors'
citations. Instead, authors should show more self-
restraint and use citations within reasonable limits, as is
customary with other reputable journals. The role of the
referee, who should pay more attention to this aspect
during the reviewing and evaluation of manuscripts, is
also vitally important.
A steady growth in the quality of the journal should
be the goal for all of us who have been connected with
this journal for many years.
Editor-in-Chief
A/Prof. Dr. Matja` Torkar
F. KAVI^KA et al.: NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE TEMPERATURE FIELD OF ...
NUMERICAL AND EXPERIMENTAL INVESTIGATION
OF THE TEMPERATURE FIELD OF A SOLIDIFYING
MASSIVE DUCTILE-CAST-IRON ROLLER
NUMERI^NE IN EKSPERIMENTALNE RAZISKAVE
TEMPERATURNEGA POLJA PRI STRJEVANJU LITEGA VALJA IZ
SIVE LITINE S KROGLASTIM GRAFITOM
Franti{ek Kavi~ka1, Karel Stránský1, Jana Dobrovská2, Bohumil Sekanina1,
Josef [tìtina1, Vasilij Gontarev3
1Brno University of Technology, Technicka 2, 616 69 Brno, Czech Republic
2VSB TU Ostrava, 17. listopadu 15/2172, 708 33 Ostrava, Czech Republic
3University of Ljubljana, A{ker~eva 12, 1000 Ljubljana, Slovenia
kavicka@fme.vutbr.cz
Prejem rokopisa – received: 2011-10-17; sprejem za objavo – accepted for publication: 2012-02-24
The quality of the working rollers used for rolling rails of different profiles is determined by the chemical and structural
composition of the material of the rollers and the production technology The requirements for the quality cannot be ensured
without a perfect knowledge of the course of the solidification, the cooling and the heat treatment of the cast rollers as well as
the kinetics of the temperature field of the casting and the mould. The solidification and cooling of the ductile-cast-iron roller in
metal and non-metal moulds is a very complicated problem of heat and mass transfer with the phase and structural changes
described by the Fourier equation. An original application of ANSYS simulated the forming of the temperature field of the
entire system, comprising the casting, the mold and the ambient. The simulation of the release of the latent heats of the phase or
structural changes is carried out by introducing the thermodynamic enthalpy function. In the experimental investigation of the
temperature field, an original methodology for the measurement of the distribution of temperatures and heat flows in the
roller-mould system had been developed and verified in operation. In the design of the original procedure, there were a number
of problems connected with the large size of the roller and the mould, the uneven dimensional changes of the solidifying roller
and the mould, the installation and the insulation of the thermocouples, the wiring of the thermocouple system, etc. The findings
regarding the kinetics of the temperature field of the roller and the mould, obtained from experimental research, were used to
determine the boundary conditions and for the verification of the numerical simulation program. The calculation of the
temperature field focused on an analysis of the effect of the mould separator on the course of the solidification of the roller. The
results of the mathematical modelling indicate that the distribution of the temperatures and the solidification in the vertical
direction is significantly uneven – and this has an effect on the internal quality of the casting.
Keywords: spheroidized graphite cast iron (SGI), massive roller, solidification, temperature field, numerical model
Kakovost delovnih valjev za valjanje `elezni{kih tirnic razli~nih profilov je dolo~ena s kemijsko sestavo, mikrostrukturo in s
proizvodno tehnologijo. Zahteve za kakovost ni mogo~e zagotoviti brez popolnega poznanja poteka strjevanja, ohlajevanja in
toplotne obdelave ulitih valjev ter kinetike temperaturnega polja ulitka in forme. Strjevanje in ohlajevanje ulitega valja, ki je ulit
iz sive litine s kroglastim grafitom v kovinsko in nekovinsko formo, je zelo zapleten problem z vidika prenosa toplote in mase s
faznimi in strukturnimi spremembami, ki jih opisuje Fourierjeva ena~ba. Z uporabo programskega paketa ANSYS-a je bilo
simulirano temperaturno polje celotnega sistema, ki obsega ulivanje, formo in okolje. Simulacija prenosa latentne toplote faznih
in strukturnih sprememb se izpelje z uporabo termodinami~ne funkcije entalpije. Pri eksperimentalni raziskavi temperaturnega
polja je bila razvita izvirna metodologija za merjenje temperaturnega polja in toplotnega toka v sistemu valj – forma. V
konstrukciji izvirnega postopka so {tevilni problemi vezani z velikostjo valjev in forme, neenako dilatacijo pri strjevanju valja in
forme, namestitvi in izolaciji termo~lenov, napeljavi sistema termo~lenov itd. Ugotovitve glede na kinetiko temperaturnega
polja valja in forme, dobljene z eksperimentalnimi raziskavami, so bile uporabljene za dolo~itev mejnih pogojev ter za potrditev
numeri~nega programa simulacije. Izra~un temperaturnega polja je bil osredinjen na analizo u~inka separatorja forme na potek
strjevanja valja. Rezultati matemati~nega modeliranja ka`ejo, da je porazdelitev temperatur in strjevanja v navpi~ni smeri
neenaka, kar vpliva na notranjo kakovost ulitka.
Klju~ne besede: siva litina s kroglastim grafitom, masivni valj, strjevanje, temperaturno polje, numeri~ni model
1 INTRODUCTION
The solidification and cooling of these rollers – partly
inside a sand mould and partly inside an iron mould – is
a very complicated problem of heat and mass transfer
with both phase and structural changes. An original
application of ANSYS simulated the forming of the
temperature field of the entire system. The introduced
3D model of the temperature field is based on the
numerical finite-element method. Experimental research
and temperature measurements must be conducted
simultaneously with the numerical calculation in order to
make the model more accurate and to verify it1–3.
2 THE ASSIGNMENT AND PREPARATION OF
THE INVESTIGATION
The assignment focused on investigating the transient
3D temperature field of a system comprising the casting,
the mould and the ambient using a numerical model. The
Materiali in tehnologije / Materials and technology 46 (2012) 4, 321–324 321
UDK 669.131.6:536.421.4 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)321(2012)
dimensions of the cylindrical casting and of the iron
mould are given in Figure 1, while the diameter of the
actual roller is 1180 mm with a height of 2100 mm.
This figure illustrates the entire set-up, comprising
two parts of the sand mould for the upper and lower
spindle ends and the iron mould. The working surface of
the iron mould is covered with a separating layer –
hereafter referred to as the "separator" (which is a special
lubricant applied in various thicknesses to the inside
walls of the iron mould and kiln-hardened at 180 °C
prior to casting). The initial temperature of the mould
was 20 °C, while the pouring temperature of the melt
was 1336 °C. The casting was performed from under-
neath with a tangential in-flow. The total time of the
casting of the spheroidal graphite cast iron (SGI) was
175 s.
2.1 Numerical model of the temperature field of the
roller
The mathematical model for the simulation of the
temperature field of the roller-mould system was created
by ANSYS. The simulation of the release of the latent
heats of the phase or structural changes is carried out by
introducing the thermodynamic enthalpy function4,5. The
program also considers the non-linearity of the task, i.e.:
• The dependence of the thermo-physical properties of
all the materials entering the system
• The dependence of the heat-transfer coefficients (on
all the boundaries of the system) on the temperature
of the surface of the casting and mould.
A series of experimental temperature measurements
was conducted for the verification of the model and the
closer determination of the boundary conditions of the
numerical solution of the temperature field. The results
from the experimental measurements were used for the
verification of the model and the correction of the boun-
dary conditions for the numerical solution of the
temperature field. The calculation of the temperature
field dealt mainly with the effect of the separator
between the casting and the iron mould on the solidi-
fication of the roller.
2.2. The results of the simulation of the temperature
field with respect to the separator thickness1
The thicknesses of the separator – for individual
simulations of the solidification and cooling of the roller
– were (0, 5, 10 and 15) mm. The point where the melt
solidifies last is the centre – 2100 mm from the base of
the iron mould. The centre of the body of the roller is on
the axis at a distance of 1050 mm from the same base.
Table 1 contains the solidification times of the roller,
which are calculated from the simulation.
Table 1: Solidification times for various separator thicknesses
Tabela 1: ^as strjevanja pri razli~nih debelinah separatorja
Separator
thickness
Solidification time
Centre of roller Entire roller
(mm) t/h t/s t/h t/s
15 8 187 29 473 8 783 31 618
10 7 845 28 242 8 124 29 246
5 7 030 25 310 7 531 27 111
0 6 510 23 437 6 554 23 593
Figures 2 and 3 illustrate the temperature field along
the longitudinal axis of the system without the separator
and with a 15 mm separator after 7.845 h, i.e., at the time
when the centre of the roller solidifies with a separator
that has a mean thickness of 10 mm (Table 1).
It is possible to observe an enlargement of the
above-mentioned red area along the vertical axis with an
increasing thickness of the separator. Figure 4 shows the
dependence of the total solidification time of the centre
of the roller and the entire roller on the separator
F. KAVI^KA et al.: NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE TEMPERATURE FIELD OF ...
322 Materiali in tehnologije / Materials and technology 46 (2012) 4, 321–324
Figure 2: The temperature field along the longitudinal section (with-
out the separator)
Slika 2: Temperaturno polje po vzdol`nem prerezu (brez separatorja)
Figure 1: The set-up of a vertically cast
Slika 1: Sistem vertikalnega ulivanja
thickness. The relationship between the separator thick-
ness and the solidification time can be described in great
detail using the linear function – as the reliability coeffi-
cient values in Figure 4 indicate for both straight lines.
The separator has proven to be a good insulator. The
solidification time of the roller inside the mould with a
15 mm separator increases by up to 26 % compared to
that without the separator.
3 EXPERIMENTAL MEASUREMENT OF THE
TEMPERATURES INSIDE THE PERMANENT
MOULD AND ON THE SURFACE OF THE
ROLLER
The distribution of the temperatures along the height
of the iron mould and roller (Figure 5) was measured in
three horizontal planes, 100 mm (section I), 1000 mm
(sect. II) and 1950 mm (section III), from the bottom
edge of the iron mould (Figure 1).
The inner surface of the permanent mould, i.e., point
9), was selected for the comparison of the temperatures
from the numerical simulation with the experimentally
obtained temperatures from inside the iron mould.
Figure 6 confirms the very close similarity between
the calculated and measured temperatures of the
permanent mould. It appears that the absolute values and
the histories of the experimental and calculated tempera-
tures in the points of the cylindrical casting and iron
mould, which had been selected for comparison, are very
similar. The determined relationship between the soli-
dification time (of the entire casting or its geometrical
centre) and the thickness of the separator applies. The
simulation of the temperature field of the roller and the
mould (using ANSYS) can therefore be considered to be
successful.
The results of the mathematical model indicate that
the distribution of the temperatures and the course of the
solidification in the vertical direction is relatively
uneven, which affects the internal quality of the casting.
While the body of the roller is solidified along the entire
height, the temperatures are still high in the lower
spindle and could be the cause of the shrinkage porosity
at the point where the spindle enters the roller. That is
why the lower spindle should solidify in a mould with a
higher heat accumulation, i.e., in a mould made of
CT-CrMg or in an iron mould. A sooner topping-up of
F. KAVI^KA et al.: NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE TEMPERATURE FIELD OF ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 321–324 323
Figure 5: Layout of measurement points
Slika 5: Polo`aj merilnih to~kFigure 3: The temperature field along the longitudinal section (with a
15 mm separator)
Slika 3: Temperaturno polje po vzdol`nem prerezu (s separatorjem 15
mm)
Figure 6: Calculated versus measured temperatures (point 9)
Slika 6: Izra~unane temperature v primerjavi z izmerjenimi (to~ka 9)
Figure 4: The influence of the separator thickness on the total
solidification time
Slika 4: Vpliv debeline separatorja na celotni ~as strjevanja
the upper spindle should ensure replenishing of the
mould with melt into the body of the roller in order to
achieve increased quality.
Acknowledgements
This analysis was conducted using a program devised
within the framework of the GACR projects No.
106/08/0606, 106/09/0940, 106/09/0969 and P107/11/
1566.
4 REFERENCES
1 J. Molinek, et al., Optimization of technological parameters of the
gravity-cast rolls for rolling rails. In: Final research report of the
project GACR 106/04/1334, Ostrava, Czech Republic, January 2007
2 J. Dobrovska et al., Two numerical models for optimization of the
foundry technology of the ceramics EUCOR. In: Proceedings and
CO ROM of the 2004 ASME Heat Transfer Fluids Egineering
Summer Conference, Charlotte, North Carolina, USA, 2004, 30–36
3 F. Kavicka et al., A numerical model of the crystallization of pure
aluminium, in Fluid Structure Interaction and Moving Boundary
Problems, WITpress Southampton, Great Britain, 2005, 619–629
4 F. Kavicka et al., Numerical optimization of the method of cooling of
a massive casting of ductile cast-iron. In: Book of Abstracts and CD
ROM of the 13th International Heat Transfer Conference, Sydney,
Australia, August 2006, 27
5 F. Kavicka et al., The optimization of a concasting technology by
two numerical models, Journal of Materials Processing Technology,
185 (2007), 152–159
F. KAVI^KA et al.: NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE TEMPERATURE FIELD OF ...
324 Materiali in tehnologije / Materials and technology 46 (2012) 4, 321–324
T. MAUDER et al.: A FUZZY-BASED OPTIMAL CONTROL ALGORITHM FOR A CONTINUOUS CASTING PROCESS
A FUZZY-BASED OPTIMAL CONTROL ALGORITHM
FOR A CONTINUOUS CASTING PROCESS
ALGORITEM, KI TEMELJI NA MEHKI LOGIKI, ZA OPTIMALNI
NADZOR KONTINUIRNEGA PROCESA ULIVANJA
Tomas Mauder, Cenek Sandera, Josef Stetina
Brno University of Technology, Faculty Mechanical Engineering, Technicka 2, Brno, Czech Republic
ymaude00@stud.fme.vutbr.cz
Prejem rokopisa – received: 2011-10-19; sprejem za objavo – accepted for publication: 2012-02-24
Nowadays, continuous casting is used for processing almost one hundred percent of liquid steel into an intermediate shape. Steel
with a poor structure and many defects is not acceptable to final customers, and therefore the producers must put strong
emphases on its quality. Many serious steel defects can be eliminated by a preceding computer simulation, optimization and a
proper control of the casting process. This paper describes an algorithm that optimizes the control parameters to ensure both a
high production rate and a high quality of the products. These controlled parameters are the casting speed and all the cooling
rates in the secondary cooling zone. The main principle of the algorithm is to get the surface and core temperatures to the
prescribed values corresponding to the required ductility of steel. The whole optimization procedure consists of two separate
parts, i.e., the numerical simulation of the process and the decision-making part based on fuzzy logic for modifying the control
parameters. By incorporating the fuzzy logic into the optimization process, the algorithm has a very robust behaviour and an
easy adaptation for different grades of steel. This algorithm runs in an off-line version and can be used as a preparation tool for a
real casting process where the proper setting of the casting parameters is crucial for achieving high-quality products at an
economical price.
Keywords: fuzzy optimization, heuristic optimization, temperature field, continuous casting
Postopek kontinuirnega ulivanja za pridobivanje vmesne oblike se danes uporablja pri obdelavi skoraj stoodstotnega dele`a
teko~ega jekla. Jeklo s slabo strukturo in {tevilnimi napakami ni sprejemljivo za kon~ne kupce, zato morajo proizvajalci veliko
pozornosti posvetiti kakovosti jekla. S predhodnimi ra~unalni{kimi simulacijami, z optimizacijo in ustreznim nadzorom procesa
ulivanja se lahko odpravijo mnoge resne napake jekla. V tem prispevku je opisan algoritem za optimizacijo parametrov nadzora,
ki zagotavlja visoko stopnjo proizvodnje ter visoko kakovost izdelkov. Med omenjene parametre nadzora spadata hitrost
ulivanja in hitrost ohlajevanja v sekundarni coni ohlajanja. Glavno na~elo algoritma je, da povr{ina in temperatura sredice
dose`eta dolo~ene vrednosti, ki ustrezajo zahtevani razteznosti jekla. Celoten postopek optimizacije obsega dva med seboj
lo~ena dela – numeri~no simulacijo postopka in pa del, namenjen odlo~anju, ki temelji na mehki logiki za spreminjanje
parametrov nadzora. Z integracijo mehke logike v postopek optimizacije algoritem dose`e zelo robustno vedenje in enostavno
prilagoditev razli~nim vrstam jekla. Ta algoritem deluje v razli~ici »off-line« in se lahko uporablja kot pripravljalno orodje za
dejanski proces ulivanja, kjer je ustrezna nastavitev parametrov ulivanja klju~nega pomena za proizvodnjo visokokakovostnih
izdelkov po ekonomi~ni ceni.
Klju~ne besede: mehka optimizacija, hevristi~na optimizacija, temperaturno polje, kontinuirno ulivanje
1 INTRODUCTION
Nowadays, continuous casting represents the main
method worldwide for forming molten steel into semi-
finished products, such as slabs, blooms, and billets. The
quality of final products plays one of the most significant
roles for customers, and thus to satisfy their demands
and maintain the highest possible productivity is a key
goal for each steelmaker.
Quality control is a prerequisite for continuous
casting and cannot be achieved without a proper
knowledge of the main physical influences of the casting
process, i.e., the solidification, the micro and macro
segregation, the crack formation, etc. During the
processing of the steel in the continuous caster, there can
arise many problems with its final quality. The solid
shell is permanently subjected to thermal and mechanical
stresses and this can result in cracks or the breakout of
liquid steel through the solid shell. The sources of the
mechanical stresses are the friction in the mould,
ferrostatic forces, a poor adjustment of the casting speed,
roll gaps, bending and straightening of strand, etc. The
cracks caused by bending and straightening are
influenced by the ductility of steel, and it is well
documented that the ductility of steel is reduced over
specific temperature ranges depending on the chemical
composition1.
Conducting industrial trials is very expensive,
time-consuming and in some cases even impossible,
which makes a computer simulation the only sustainable
option. There exist various ways of computer or mathe-
matical optimization and each of them has its advantages
and drawbacks. Therefore, to ensure the correctness of
the algorithm, it is necessary to validate the simulation
output using real, measured data.
Previous researches concerned with the optimal
control of a continuous casting process were generally
based on simplified 1D or 2D temperature-field models1
and were optimized by mathematical programming or
heuristic methods, e.g., genetic algorithms2, the firefly
Materiali in tehnologije / Materials and technology 46 (2012) 4, 325–328 325
UDK 621.74.047:66.011 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)325(2012)
algorithm3 or by neural networks. Many of these models
are based on simplified assumptions, and therefore they
describe the casting process very roughly and not
satisfactory. We developed our original numerical model
for the temperature field of the real caster and its results
were validated by temperature measurements acquired
by pyrometers. This more precise model simulates the
process controlled by several numerical parameters and
the goal is to find their values such that the resultant
temperature field is optimal. This algorithm is inspired
by our previous research3 and it is enhanced by a fuzzy
logic inference mechanism, which makes its behaviour
more robust and its setting easier to adjust.
2 PROBLEM DESCRIPTION AND AN OUTLINE
OF THE ALGORITHM
The natural inclination of all steelmakers is to cast as
fast as possible but with a preservation of the required
material quality. Thus, the problem of optimal casting
control can be formulated as follows.
The properties of the final material are highly
dependent on its temperature courses emerging during
the casting process. Therefore, we need to adjust the
parameters of a caster in such a way that the final
temperature field is optimal. We can control the process
by the regulation of the casting speed and by all the
cooling rates in the secondary cooling zone.
Before the optimization starts, it is necessary to
define the temperature field, which leads to the optimal
material and mechanical properties. These optimal
temperature courses are defined by experts (casting
operators, material scientists, etc.) and they differ with
the chemical composition of the steel.
To deal with the high variability of the described
problem, the heat and mass phenomenon must also be
taken into account, and therefore the employed nume-
rical model has to cover all the phase and structural
changes. The search for the optimal cooling rates and the
highest possible casting speed is left to the following
fuzzy-logic heuristic algorithm.
To describe the optimal temperature field, the experts
define a set of points along the caster and prescribe their
optimal temperature ranges. Then the algorithm random-
ly generates the values for all the input parameters and
simulates the final temperature field. After comparing
the computed temperatures with the prescribed ones, the
algorithm decides which nozzles need to be adjusted and
how to change the casting speed. Particular decisions for
changing the values are taken by the fuzzy inference
logic. These steps are iteratively repeated until the
required temperatures are successfully reached.
3 MATHEMATICAL MODEL OF THE
TEMPERATURE FIELD
The developed numerical model of the 2D tempe-
rature field computes the formation of the solidification
and the temperature distribution of the cast strand. This
temperature field is modelled by the Fourier-Kirchhoff
equation1–3, where the velocity component vy(m/s) is
considered only in the direction of casting:
∂
∂
∂
∂
H
k T T v
H
yy
= +eff ( ) ( )Δ (1)
Equation (1) describes a transient heat transfer, where
keff/(W/m K) is the effective thermal conductivity, T/K is
the temperature, H/(J/m3) is the volume enthalpy, /s is
the time and x and y are spatial coordinates. The phase
and structural changes are included in the model by the
latent-heat accumulation method, where the enthalpy is
used as the primary variable and the temperature is
calculated from a defined enthalpy-temperature relation-
ship3:
H c H
f
T
T
= −
⎛
⎝
⎜ ⎞
⎠
⎟∫      ( ) ( ) ( )Δ
∂
∂
s
d
0
(2)
where H/(J/kg) is the latent heat, /(kg/m3) is the
density, c/(J/kg K) is the specific heat capacity and fs is
the solid fraction.
In order to have a well-posed problem, the initial and
boundary conditions must be stated. The boundary
conditions include the heat fluxes in the mould and under
the rollers, the forced convection under the nozzles and
the free convection and radiation in the tertiary cooling
zone. Equation (1) is discretized by the finite-difference
method3,4 using an explicit formula for the time deriva-
tive. The mesh for the finite-difference scheme is
non-equidistant and its nodes are adapted to the real
rollers and nozzle positions. Our numerical model allows
us to apply various enthalpy-temperature functions and
thermal conductivity-temperature curves, and thus the
temperature field can be calculated for various steels just
by redefining the chemical composition.
Thermo-physical parameters such as the thermal
conductivity, the density, the specific heat capacity and
the enthalpy and their temperature dependence are
computed from a specific chemical composition of steel
by using the solidification analysis package IDS.
The numerical model is designed and verified for the
radial slab caster operating in EVRAZ VÍTKOVICE
STEEL, a. s. The caster contains 12 coolant circuits and
the cross-section of the investigated slab is 1550 mm ×
250 mm.
4 DESCRIPTION OF THE ALGORITHM
The algorithm searches for the optimal values of the
control parameters (cooling rates and the casting speed).
The inputs to the algorithm are the required temperature
T. MAUDER et al.: A FUZZY-BASED OPTIMAL CONTROL ALGORITHM FOR A CONTINUOUS CASTING PROCESS
326 Materiali in tehnologije / Materials and technology 46 (2012) 4, 325–328
ranges in the points along the caster and the maximum
metallurgical length. For each cooling circuit, the
maximum cooling rate is defined. The initial value of
each control parameter is randomly chosen from the
permitted range and the model is consequently evaluated
by the numerical simulation. From the computed result,
the algorithm extracts temperatures in the controlled
points and by a comparison with the prescribed values
determines their deviations (errors). Also, the metallur-
gical length (the length from the meniscus to the point
where the steel is fully solidified) is computed. Using all
this information the algorithm infers the modifications
for all the controlled input parameters.
The temperature courses at the controlled points are
dominantly influenced by the two preceding coolant
circuits. The closer circuit has a greater influence.
Therefore, each coolant circuit defines a numerical value
for each controlled point that describes how much it
impacts on the temperature in the controlled point. These
values are expertly estimated (in the range from 0 to 10)
and they are strongly related to the distance from the
controlled points.
In the presented model there are two fuzzy inference
mechanisms – one for the modification of the cooling
rates and one for the control of the casting speed.
The inference rules for the cooling rates have the
following form: "IF error IS adj1 AND impact IS adj2
THEN modification IS adj3". For each controlled point
the adj1 describes the temperature error, the adj2 is the
cooling impact of the given circuit and the adj3 is the
modification inferred for the cooling rate. Table 1 shows
all the rules in the matrix form.
Table 1: The dependences of adj3 on adj1 and adj2. The abbreviations
stand for Very Small, Small, Medium, Big and their fuzzy sets are
equidistantly distributed along the corresponding universes.
Tabela 1: Odvisnost ozna~evalca 3 od ozna~evalca 1 in ozna~evalca 2.
Okraj{ave pomenijo VS – zelo majhen, S – majhen, M – srednji, B –
velik, njihove mehke nastavitve pa so v enakih razdaljah razporejene
po ustreznih splo{nih parametrih.
adj2 / adj1 VS S M B
S VS VS S S
M VS VS S M
B VS S M B
Sometimes, one circuit can get several different
modifications and if it happens the algorithm takes the
one with the highest absolute value. The defuzzification
method is the standard centre-of-gravity function. If the
maximum absolute value of all the temperature errors
does not exceed a particular limit, the algorithm con-
sequently computes the modification for the casting
speed. The reason for introducing the limit is that if the
maximum error is too big (i.e., the solution is far from
the optimum), we have no information about whether the
speed is high or not and at first it is better to stabilize the
process and then to infer the speed modification.
The rules for the modification of the casting speed
(Table 2) are in the form: "IF maximum_error IS adj4
AND metalurgical_length IS tadj5 THEN modification
IS adj6". The adj4 describes the maximum error over all
the controlled points, the adj5 is the metallurgical length
and the adj6 is the inferred modification of the casting
speed.
For a detailed explanation of the fuzzy logic and the
fuzzy inference, see5.
The values of the impacts for each cooling circuit to
the controlled point are chosen to be 8 for the closest
preceding circuit and 2 for the circuit placed before
(except the last controlled point, where the distance to
the second closest circuit is much longer, and therefore
the values are 9 and 1).
Table 2: The dependences of adj6 on adj4 and adj5. The abbreviations
stand for Very Small, Small, OK, Big, Very Big and More, Little
More, Little Less, Less, Nothing and their fuzzy sets are equidistantly
distributed along the corresponding universes.
Tabela 2: Odvisnost ozna~evalca 6 od ozna~evalca 4 in ozna~evalca 5.
Okraj{ave pomenijo VS – zelo majhen, S – majhen, OK, B – velik,
VB – zelo velik in MO – ve~, LM – malo ve~, LL – malo manj, L –
manj, N – ni~, njihove mehke nastavitve pa so v enakih razdaljah
razporejene po ustreznih splo{nih parametrih.
adj4 / adj5 VS S OK B VB
S MO LM LM LL L
M LM N N N LL
5 RESULTS OF EXPERIMENT
The examined grades of steel are the low-carbon steel
S355J0H and the high-carbon steel X210CrW12. The
expertly defined temperature ranges in the controlled
points are depicted in Figure 1 and Figure 2 (the grey
rectangles). The crucial reason for defining these
temperature ranges in this way is that approximately the
first half of caster is curved and therefore to decrease the
mechanical stresses it is better to keep the temperature
T. MAUDER et al.: A FUZZY-BASED OPTIMAL CONTROL ALGORITHM FOR A CONTINUOUS CASTING PROCESS
Materiali in tehnologije / Materials and technology 46 (2012) 4, 325–328 327
Figure 1: Final temperature distributions (S355J0H)
Slika 1: Kon~ne razporeditve temperature (S355J0H)
above a certain level (for S355J0H, 1000 °C, and for
X210CrW12, 800 °C). The value of the maximum
metallurgical length is 20 m (from the practice) and the
casting temperature is 1550 °C. Reaching the tempe-
rature courses on surface ensures that the mechanical and
material qualities of the final material will satisfy the
demands. The optimal values computed by this algo-
rithm (casting speed and heat-transfer coefficients) are
shown in Figure 1 and Figure 2.
Usually, the most important indicator characterizing
the efficiency of iterative optimization algorithms is the
number of evaluations of the model. The computation of
the numerical model is very time-consuming, and
therefore each repetition can significantly prolong the
computations. Our algorithm is able to find the optimal
input parameters in 50 evaluations, on average. The tests
ran several times for different grades of steel and the
number of evaluations never exceeded 65. The compu-
tation time takes a few hours.
6 CONCLUSION
The problem of the optimization of a continuous
casting process can be efficiently solved by the described
algorithm. The algorithm based on heuristics and
incorporating the fuzzy logic behaves in a very robust
way and it is easily adaptable to any grade of steel and
caster geometry. The number of evaluations of the
included numerical model is low, and thereby the
algorithm proves its high efficiency. Further research will
be focused on extending the numerical model to 3D and
a generalization of the used optimization constraints.
Acknowledgement
The authors gratefully acknowledge the financial
support from the projects GA106/08/0606, GA106/09/
0940 and GACR P107/11/1566, funded by the Czech
Science Foundation, project ED0002/01/01 – NETME
Centre and Specific research BD 13101001.
7 REFERENCES
1 J. P. Birat, et al.: The Making, Shaping and Treating of Steel: Casting
Volume: 11th Edition. Alan W. Cramb, Pittsburgh, PA, USA: The
AISE Steel Foundation, 2003, 1000 s
2 C. A. Santos, J. A. Spim, A. Garcia, Mathematical modeling and
optimization strategies (genetic algorithm and knowledge base)
applied to the continuous casting of steel, Engineering Applications
of Artificial Intelligence, 16 (2003), 511–527
3 D. M. Stefanescu, Science and Engineering of Casting Solidification,
Second Edition, New York, Springer Science, 2009, 402
4 T. Mauder, C. Sandera, J. Stetina, M. Seda, Optimization of Quality
of Continuously Cast Steel Slabs by Using Firefly Algorithm, Mater.
Tehnol., 45 (2011) 4, 347–350
5 H. T. Nguyen, E. A.Walker, A First Course in Fuzzy Logic, CRC
Press, 1999, 392
T. MAUDER et al.: A FUZZY-BASED OPTIMAL CONTROL ALGORITHM FOR A CONTINUOUS CASTING PROCESS
328 Materiali in tehnologije / Materials and technology 46 (2012) 4, 325–328
Figure 2: Final Temperature distributions (X210CrW12)
Slika 2: Kon~ne razporeditve temperature (X210CrW12)
G. LOJEN et al.: THERMAL STABILITY OF Al-Mn-Be MELT-SPUN RIBBONS
THERMAL STABILITY OF Al-Mn-Be MELT-SPUN
RIBBONS
TEMPERATURNA OBSTOJNOST HITRO STRJENIH TRAKOV
Al-Mn-Be
Gorazd Lojen, Tonica Bon~ina, Franc Zupani~
University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, Si-2000 Maribor, Slovenia
gorazd.lojen@uni-mb.si
Prejem rokopisa – received: 2011-10-19; sprejem za objavo – accepted for publication: 2012-03-01
As with other kinds of finely dispersed, small particles, icosahedral quasicrystals (IQCs) also have a distinct strengthening
effect, which can be utilised to enhance the mechanical properties of aluminium alloys. In Al–Mn–Be alloys, IQCs already form
at moderate cooling rates, which can be utilised when using some conventional casting processes, like mould or injection
casting. In this case, however, crystalline intermetallic phases are also present and the mechanical properties are inferior to those
of two-phase Al-IQC alloys. Two-phase microstructures are feasible using rapid solidification techniques, e.g., melt spinning.
Further processing often involves technologies (consolidation, extrusion etc.), which include the influence of heat. The alloy
must not be overheated in order to preserve the strengthening effect of the metastable IQC-particles.
In this investigation the Al-Mn-Be alloy was melt-spun using a free-jet melt spinner. Subsequently, the thermal stability of the
IQCs was explored by annealing the ribbons for 24 h at different temperatures. The samples were examined in the as-cast and
heat-treated conditions using a dual-beam, scanning electron microscope (SEM-FIB), X-ray diffraction (XRD) and differential
scanning calorimetry (DSC). It was discovered that in the as-cast condition, the ribbons had a two-phase microstructure,
consisting of an Al matrix and finely dispersed IQCs. During annealing at temperatures up to 400 °C, the IQCs did not
decompose and the phase composition remained unchanged. Annealing at 500 °C and at higher temperatures caused a
decomposition of the IQCs, and only the crystalline intermetallic phases Al6Mn and Be4AlMn could be found in the Al matrix.
Keywords: quasicrystal, Al-Mn-Be alloy, thermal stability
Ikozaedri~ni kvazikristali (IQC) imajo kakor druge vrste fino razpr{enih drobnih delcev utrjevalni u~inek, ki ga lahko
izkoristimo za pove~anje trdnosti aluminijevih zlitin. V zlitinah Al-Mn-Be kvazikristali nastajajo `e pri zmernih ohlajevalnih
hitrostih, kakr{ne dose`emo pri postopkih litja, kot sta npr. kokilno litje in tla~no litje. Toda mikrostruktura, ki nastane pri teh
postopkih litja, vsebuje tudi kristalne intermetalne faze in trdnostne lastnosti so slab{e kot pri zlitinah z dvofazno mikrostrukturo
(Al-IQC). Dvofazno mikrostrukturo pa lahko dose`emo le z velikimi ohlajevalnimi hitrostmi, tj. s postopki hitrega strjevanja.
Nadaljnja predelava hitro strjenih materialov pogosto poteka s postopki (kompaktiranje, iztiskovanje itd.), ki zahtevajo
segrevanje. Pri tem materiala ne smemo pregreti, ~e `elimo ohraniti utrjevalni u~inek kvazikristalov.
V okviru te raziskave so bili uliti hitro strjeni trakovi Al-Mn-Be s postopkom prostega litja na vrte~e se kolo (angl.: melt
spinning). Temperaturno obstojnost trakov hitro strjene zlitine smo ugotavljali z `arjenjem 24 h pri razli~nih temperaturah.
Mikrostrukture trakov smo analizirali v litem stanju in po toplotni obdelavi z vrsti~nim elektronskim mikroskopom z
elektronskim in ionskim curkom (SEM-FIB), rentgensko difrakcijo (XRD) in z diferen~no dinami~no kalorimetrijo (DSC).
Trakovi so imeli v litem stanju dvofazno mikrostrukturo, sestavljeno iz matice Al in fino dispergiranih kvazikristalov. @arjenje pri
temperaturah do 400 °C ni povzro~ilo razpada kvazikristalov in fazna sestava je ostala nespremenjena. Po `arjenju pri 500 °C ali
vi{jih temperaturah pa je bila mikrostruktura sestavljena iz matice Al in kristalnih intermetalnih faz Al6Mn ter Be4AlMn.
Klju~ne besede: kvazikristal, zlitina Al-Mn-Be, temperaturna obstojnost
1 INTRODUCTION
The existence of icosahedral quasicrystals (IQCs or
i-phase) was discovered by Schechtman et al. in
melt-spun Al-Mn alloys1. Later, the possibility of IQC
formation was proven for dozens of others alloys2,3.
Some elements, beryllium being one of them,
strongly increase the quasicrystalline-forming ability in
Al–Mn alloys4–6. IQCs had already been observed in
different conventionally cast ternary and quaternary
alloys based on the Al-Mn system, and also in several
Al-Mn-Be alloys. In conventionally cast alloys, however,
other crystalline phases are also present4–8 in addition to
the IQCs.
One of the most promising potential applications for
quasicrystals is their utilisation as a strengthening phase
in lightweight aluminium alloys. As early as 1992 Inoue9
reported on Al–Mn–Ce melt-spun ribbons achieving
tensile strengths above 1 GPa. Recently, a hardness of
over 290 HV was reported10 for Al-Mn-Be melt-spun
ribbons, which, based on the well-known empirical
relationship between hardness and tensile strength, also
allows us to expect tensile strengths in the range of
1 GPa. However, extraordinarily high strengths were
only reported for alloys having two-phase (Al + IQCs)
microstructures, consisting of a fine dispersion of
quasicrystalline particles in an Al-rich solid solution
matrix, as yet unattained by conventional casting.
Melt-spun ribbons are rarely directly applicable,
because of their dimensions. Consequently, for the vast
majority of applications the ribbons must be compacted
into much larger pieces of material. The compacting
processes mostly involve heating, so rapidly solidified
materials must exhibit sufficient thermal stability in
order to preserve their advantages.
Materiali in tehnologije / Materials and technology 46 (2012) 4, 329–333 329
UDK 669.71’74’725:621.74.046 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)329(2012)
Unfortunately, IQCs in Al–Mn-based alloys have
been discovered to be metastable. The thermal stability
of the IQCs in a rapidly solidified binary Al-5 wt.% Mn
alloy was reported to be below 1 hour of isothermal
annealing at 500 °C 11, and by means of DSC analysis
with a high heating rate of 40 K min–1, the onset
temperature of thermal decomposition was estimated to
be 472 °C for the mole fraction Al-14 % Mn alloy8. The
thermal stabilities of the conventionally cast ternary
alloys Al-2.5Mn-7.5Be (in mole fractons, x/%) and
Al-5Be-1.5Mn (in mass fractions, w/%) were investi-
gated by Song et al.4 and Chang et al.7, respectively. In
both studies the ingots, initially consisting of ICQs and
hexagonal approximants dispersed in an Al matrix, were
annealed for 100 h at 540 °C. According to Song et al.4
the annealing caused no phase transformations, and in
addition the DSC did not reveal any transformations
below 600 °C. In contrast, Chang et al.7 reported that
upon annealing the IQC-phase dissolved and only
hexagonal approximants remained in the matrix.
Although these results appear to be contradictory, the
differences could possibly be attributed to different
chemical compositions and different cooling rates during
the solidification of the ingots and, in addition, because
in both cases no shorter and no longer annealing times
than 100 h were applied, the actual thermal stabilities did
not necessarily need to be significantly different at all.
In any case, due to much lower cooling-rates,
conventionally cast ingots as well as IQCs also contain
crystalline intermetallics, and the microstructure is much
coarser than in rapidly solidified alloys. As a result, the
conventionally cast material is, already in an as-cast
condition, much closer to the thermodynamic equili-
brium than the rapidly solidified one, which in an as-cast
condition exhibits an extremely fine, two-phase (Al +
IQCs) microstructure. Therefore, the driving force for
changes during annealing is higher in rapidly solidified
materials and, consequently, a lower thermal stability can
be expected. On the other hand, due to the stabilising
effect of Be, the thermal stability of the IQC phase could
be improved in rapidly solidified, ternary, Al-Mn-Be
alloys.
The aim of this work was to evaluate the thermal
stability of IQCs in a rapidly solidified, ternary,
Al-Mn-Be alloy, which has yet to be determined.
2 EXPERIMENTAL
Although with beryllium the critical content of Mn to
produce IQCs is reduced to no more than x = 2.5 % 4, a
higher Mn content was chosen in order to achieve a
higher volume fraction of quasicrystals. The nominal
composition of the alloy was (in mole fractions x) 89 %
Al, 6 % Mn and 5 % Be, Table 1. It was prepared from
Al 99.99, and master alloys of Al–Mn and Al–Be
containing (in mass fractions) 30 % Mn and 5.5 % Be,
respectively. The precursors were vacuum induction
melted and cast into 50-mm round bars. The bars were
cut into appropriate lengths and melt-spun in a 30
M-type Melt Spinner, Marko Materials Inc, USA. The
melt-spinning conditions were as follows: a BN-pro-
tected graphite crucible with a 1-mm orifice, a wheel
speed of 25 m min–1, a casting temperature of 900–950
°C.
Table 1: Nominal and analysed chemical compositions of the melt-
spun ribbons (ICP-AES)
Tabela 1: Imenske in analizirane kemi~ne sestave hitro strjenih trakov,
ulitih s postopkom litja na vrte~e se kolo (ICP-AES)
element x(Al)/% x(Mn)/% x(Be)/%
nominal 89 6 5
analysed 88.6 5.6 5.8
x/% – mole fraction; x/% – molski dele`
The chemical composition of the melt-spun ribbons
was verified using ICP-AES (inductively coupled
plasma-atomic emission spectroscopy), see Table 1.
The microstructures were investigated in the as-cast
condition and after annealing under a protective
atmosphere of Ar. The annealing temperatures were
(200, 300, 400, 500 and 550) °C, where the annealing
time was always 24 h, regardless of the temperature.
Different techniques were applied for the metallographic
characterisation. The scanning electron microscopy was
performed using a dual-beam FEI Quanta 200-3D
microscope (SEM-FIB). The X-ray diffraction was
carried out at the XRD1 beam-line (Elettra, Sinchrotrone
Trieste, Italy) using synchrotron X-rays with a wave-
length of 0.1 nm in the transmission mode. The phase
compositions were identified using Powder Diffraction
File-212. The DSC analyses were performed with a STA
449 Jupiter, from room temperature to 640°C, with
heating and cooling rates of 10 °C min–1 in nitrogen.
3 RESULTS AND DISCUSSION
3.1 As-cast ribbons
The analysed composition is given in Table 1, along
with the nominal composition. The actual composition
only slightly deviated from the nominal one; therefore,
the ribbons were evaluated as being appropriate for the
intended purpose.
The microstructural analysis of the as-cast ribbons
revealed a two-phase microstructure, consisting of fine
IQCs dispersed throughout the Al matrix. Figure 1
shows a backscattered-electron SEM image. The backs-
cattered-electron imaging provided the best phase
contrast. The Al-matrix appears dark. The coarser quasi-
crystals could be identified by their characteristic
dendritic morphology and the finer ones, most of them
being too small to be distinguishable in a SEM micro-
graph, by their shapes as pentagonal dodecahedrons13.
Beside IQCs, no other phase could be found in the
matrix of the as-cast ribbons when using scanning
G. LOJEN et al.: THERMAL STABILITY OF Al-Mn-Be MELT-SPUN RIBBONS
330 Materiali in tehnologije / Materials and technology 46 (2012) 4, 329–333
electron microscopy. In addition, the XRD analysis (the
lowest curve in Figure 2) only confirmed the presence of
the IQC and the Al. Therefore, the as-cast ribbons were
considered to have a two-phase Al + IQC microstruc-
ture.
Upon heating, Figure 3, the DSC analysis of as the
as-cast ribbons showed a phase transformation that
started at approximately 530 °C, reached its peak at
approximately 570 °C and ended at approximately 600
°C. The thermal decomposition of the metastable ICQs
started at 530 °C, as indicated by the absence of a
reverse transformation upon cooling. As is quite
common for DSC, the heating rate was relatively fast.
Therefore, the influence of the transformation kinetics
was unavoidable and the determined transformation
temperatures were shifted to higher values. The
influence of the transformation kinetics reduced with the
prolongation of the annealing time, so it was expected
that during the 24-hours of isothermal annealing,
decomposition would take place at a lower temperature.
Therefore, only one of the temperatures for the 24 h
annealing was selected above the DSC-estimated start of
the decomposition, 550 °C, and one of the lower
temperatures only slightly below, i.e., 500 °C.
3.2 Annealed ribbons
The X-ray diffractograms of the samples annealed at
200 °C, 300 °C, and 400 °C were very similar to the
diffractograms of the as-cast sample, Figure 2. No
distinguishable peaks for the crystalline intermetallic
phases could be recognised and, only those peaks for
IQCs and Al could be observed. The absence of peaks of
the crystalline intermetallics indicates that at tempe-
ratures up to 400 °C over 24 h the decomposition did not
start or at least did not advance enough to be detected by
XRD. Also, in the SEM micrographs of those samples
annealed at temperatures up to 400 °C, Figure 4, no
distinct changes in the average size and morphology of
the larger IQCs could be observed. Compared to the
as-cast condition (Figure 1), the number of smaller
particles slightly diminished, indicating that annealing at
400 °C caused a certain degree of Ostwald ripening.
In the XRD diffractograms of the samples annealed
at 500 °C and 550 °C an absence of IQC peaks could be
observed at first sight. In their place, distinct peaks for
the intermetallic phases Be4AlMn and Al6Mn were
present. The pronounced equality of both XRD diffracto-
grams indicated that at 500 °C the decomposition of the
IQCs had already been completed within the 24 h. The
SEM micrograph of the sample annealed at 550 °C,
Figure 5a, at first sight appeared similar to the
micrographs of samples annealed at temperatures up to
400 °C. The average size of the large particles remained
almost unchanged. However, their morphology changed
G. LOJEN et al.: THERMAL STABILITY OF Al-Mn-Be MELT-SPUN RIBBONS
Materiali in tehnologije / Materials and technology 46 (2012) 4, 329–333 331
Figure 2: X-ray diffractograms. From bottom to top: as-cast ribbon,
ribbons annealed for 24 h at 200 °C, 300 °C, 400 °C, 500 °C and 550
°C.
Slika 2: Rentgenski difraktogrami. Od spodaj navzgor: trak v litem
stanju, trakovi, `arjeni 24 h pri 200 °C, 300 °C, 400 °C, 500 °C in 550
°C.
Figure 3: DSC curve of as-cast sample upon heating; heating rate 10
K min–1
Slika 3: DSC-krivulja traku v litem stanju; hitrost segrevanja 10
K min–1Figure 1: Microstructure of the as-cast ribbon, SEM-image. The
coarser quasicrystals exhibit a characteristic dendritic morphology.
Slika 1: Mikrostruktura traku v litem stanju, SEM-slika. Grobi kvazi-
kristali imajo zna~ilno dendritno morfologijo.
to a more compact shape. According to the XRD
analysis, Be4AlMn and Al6Mn were present in the Al
matrix. As can be seen from the micrographs in Figures
4 and 5a, through the thermal decomposition of the
IQCs, neither the size nor the distribution of the
dispersed particles changed significantly. Obviously, the
sizes and the distribution of the crystalline intermetallic
particles, grown through the thermal decomposition of
the IQCs, were determined by the sizes and distribution
of the IQCs. However, as reported by Bon~ina et al.14, in
the SEM backscattered-electron images the Be4AlMn
appears only slightly brighter than the Al6Mn and,
consequently, cannot be unerringly recognised. There-
fore, cross-sections were cut by the focused ion beam
in-situ in the dual-beam SEM-FIB microscope. In the
FIB secondary-electron image in Figure 5b, both phases
could be clearly distinguished. The Be4AlMn appeared
as very bright, almost white, and the surface of the
Al6Mn appeared as smoothly grey. Consistent with the
peak heights in the XRD diagrams in Figure 5b, it can
be observed that the volume fraction of Be4AlMn was
significantly lower than the fraction of Al6Mn.
4 CONCLUSIONS
The metallographic investigations of the rapidly
solidified (melt-spun) Al-Mn-Be alloy led to the
following conclusions:
1. The microstructure of the as-cast melt-spun (in mole
fractions) 89 % Al-6 % Mn-5 % Be ribbons consisted
of IQCs dispersed in an Al matrix. No quasicrystal
approximants or equilibrium intermetallics could be
detected.
2. The thermal stability of the rapidly solidified 89 %
Al-6 % Mn-5 % Be alloy lasted at least 24 h at
temperatures up to 400 °C. After 24 h annealing at
400 °C, only insignificant coarsening of the IQCs
was perceivable, whilst the phase composition was
preserved.
3. Upon annealing at 500 °C the thermal decomposition
of the IQCs into Al6Mn and Be4AlMn was completed
after no later than 24 h. The sizes and dispersion of
the newly formed Al6Mn–Be4AlMn clusters were
very similar to the sizes and dispersion of the
dendritic IQCs. As a result, due to the thermal
decomposition of the IQCs, the mechanical
properties may not drastically decrease.
4. The rapidly solidified 89 % Al-6 % Mn-5 % Be alloy
can be processed using techniques that involve
elevated temperatures.
Acknowledgments
The research leading to these results received funding
from the European Community’s Seventh Framework
Programme (FP7/2007–2013) under grant agreement N’
226716. This work was also partly financed from the
G. LOJEN et al.: THERMAL STABILITY OF Al-Mn-Be MELT-SPUN RIBBONS
332 Materiali in tehnologije / Materials and technology 46 (2012) 4, 329–333
Figure 5: Ribbon annealed for 24 h at 550 °C: a) SEM back-
scattered-electron image: Be4AlMn and Al6Mn can practically not be
distinguished; b) FIB secondary-electron image of an in-situ prepared
cross-section: all phases identified by XRD can be recognised
Slika 5: Trak, `arjen 24 h pri 550 °C: a) SEM-slika z odbitimi
elektroni: Be4AlMn in Al6Mn prakti~no ni mogo~e razlikovati; b)
FIB-slika in situ pripravljenega prereza: prepoznavne so vse faze,
odkrite z XRD
Figure 4: SEM micrograph of a sample after 24-h annealing at 400
°C: compared to the as-cast sample, no distinct changes can be
observed.
Slika 4: SEM-slika vzorca po 24-urnem `arjenju pri 400 °C: v
primerjavi z vzorcem v litem stanju ni opaziti ve~jih sprememb.
research programme P2–0120, ARRS. The authors also
wish to thank Mr. Rok [ulek, University of Maribor,
Faculty of Mechanical Engineering, for preparing the
alloy billets, Mr. Borivoj [u{tar{i~ PhD, Institute of
Metals and Technology in Ljubljana, for melt spinning,
and prof. Jo`e Medved PhD, University of Ljubljana,
Faculty of Natural Sciences and Engineering, for the
DSC analyses.
5 REFERENCES
1 D. Shechtman, I. Blech, D. Gratias, J.W. Cahn, A Metallic Phase
with Long-Ranged Orientation Order and No Translation Simetry,
Phys. Rev. Lett., 53 (1984), 1951–1953
2 C. Suryanarayana, H. Jones, Formation and characteristics of quasi-
crystalline phases: a review, Int. J. Rapid Solidif., 3 (1987), 253–93
3 K. F. Kelton, Quasicrystals: structure and properties, Int. Mater. Rev.,
38 (1993), 105–37
4 G. S. Song, E. Fleury, S. H. Kim, W. T. Kim, D. H. Kim, Enhance-
ment of the quasicrystal-forming ability in Al-based alloys by
Be-addition, J. of Alloys and Compounds, 342 (2002), 251–255
5 G. S. Song, E. Fleury, S. H. Kim, W. T. Kim, D. H. Kim, Formation
and stability of quasicrystalline and hexagonal approximant phases
in an Al–Mn–Be alloy, J. Mater. Res., 17 (2002), 1671–1677
6 S. H. Kim, G. S. Song, E. Fleury, K. Chattopadhyay, W. T. Kim, D.
H. Kim, Icosahedral quasicrystalline and hexagonal approximant
phases in the Al–Mn–Be alloy system, Philosophical Magazine A, 82
(2002), 1495–1508
7 H. J. Chang, E. Fleury, G. S. Song, W. T. Kim, D. H. Kim, Formation
of quasicrystalline phases in Al-rich Al–Mn–Be alloys, J. Non-cryst.
solids, 334 (2004), 12–16
8 F. Schurack, J. Eckert, L. Schultz, Synthesis and mechanical pro-
perties of cast quasicrystal-reinforced Al-alloys, Acta mater., 49
(2001), 1351–1361
9 A. Inoue, M. Watanabe, H. M. Kimura, F. Takahashi, A. Nagata, T.
Masumoto, High mechanical strength of quasicrystalline phase
surrounded by fcc-Al phase in rapidly solidified Al-Mn-Ce alloys,
Materials transactions JIM, 33 (1992) 8, 723–729
10 F. Zupani~, T. Bon~ina, B. [u{tar{i~, I. An`el, B. Markoli, Micro-
structure of Al-Mn-Be melt-spun ribbons, Mater. charact., 59 (2008)
9, 1245–1251
11 K. Saksl, D. Vojtech, H. Franz, Quasicrystal–crystal structural trans-
formation in Al–5 wt.% Mn alloy, J. Mater. Sci., 42 (2007),
7198–7201
12 Powder Diffraction File-2, Release 2004. International Center for
Diffraction Data, Newtown, USA, 2004
13 F. Zupani~, B. Markoli, Morphologies of Icosahedral quasicrystals in
Al-Mn-Me-(Cu) alloys, In: Puckermann B. E. (editor): Quasicrystals.
Nova science publishers Inc., New York 2011, 195–217
14 T. Bon~ina, B. Markoli, I. An`el, F. Zupani~, Metallographic tech-
niques for the characterization of quasicrystalline phases in
aluminium alloys, Z. Kristallogr., 223 (2008), 747–750
G. LOJEN et al.: THERMAL STABILITY OF Al-Mn-Be MELT-SPUN RIBBONS
Materiali in tehnologije / Materials and technology 46 (2012) 4, 329–333 333

L. KLIMES et al.: CHALLENGES IN THE COMPUTER MODELING OF PHASE CHANGE MATERIALS
CHALLENGES IN THE COMPUTER MODELING OF
PHASE CHANGE MATERIALS
IZZIVI V RA^UNALNI[KEM MODELIRANJU MATERIALOV S
FAZNO SPREMEMBO
Lubomir Klimes1, Pavel Charvat1, Milan Ostry2
1Brno University of Technology, Faculty of Mechanical Engineering, Technická 2896/2, 616 69 Brno, Czech Republic
2Brno University of Technology, Faculty of Civil Engineering, Veveøí 331/95, 602 00 Brno, Czech Republic
yklime17@stud.fme.vutbr.cz
Prejem rokopisa – received: 2011-10-20; sprejem za objavo – accepted for publication: 2012-03-01
Phase-change materials (PCMs) are a well-established category of materials with many possible applications, ranging from the
stabilization of temperature to heat or cold storage. The main principle of PCMs is the utilization of the latent heat of a phase
change for energy storage. Though many pure chemical elements can be used as PCMs, a PCM very often consists of a number
of substances. The main reason for creating a PCM as a mixture of various substances is to achieve a desirable melting
temperature for a particular application. However, these mixed PCMs require accurate and reliable methods for determining
their physical properties, since for numerical modeling the thermal properties of materials and their proper determination
represent a significant issue that considerably affects the accuracy and credibility of numerical simulations and their outcomes.
The thermal properties of PCMs are usually obtained by differential scanning calorimetry (DSC), based on temperature and heat
measurements of the investigated and reference materials. There are several approaches to modeling materials with phase
changes. In this paper, the focus is on the enthalpy approach and the effective heat capacity method. Both techniques, which
utilize the results from a particular DSC measurement, allow a treatment with the latent heat, and as a result the desired heat or
cold storage may be effectively simulated. The presented methods were implemented using the results from DSC measurements
in order to simulate a solar air collector with a PCM. The obtained results are presented and discussed. The paper also concerns
the problems associated with the uncertainty of material properties and their influence on numerical simulations.
Keywords: phase change material (PCM), enthalpy method, effective heat capacity method, solar air collector
Materiali s faznimi premenami (PCM) so uveljavljena kategorija materialov z mnogimi podro~ji uporabe: od stabilizacije
temperature do shranjevanja toplote ali hladu. Glavni princip PMS je uporaba latentne toplote pri fazni premeni za shranjevanje
energije. ^eprav je mogo~e uporabiti kot PCM mnoge ~iste kemijske elemente, se pogosto uporablja me{anica razli~nih
elementov. Glavni razlog za uporabo me{anice za PCM je zagotavljanje `elene temperature tali{~a za dolo~en namen. Vendar pa
me{anica PCM-ov zahteva natan~ne in izvedljive metode za dolo~anje fizikalnih lastnosti. Numeri~no modeliranje termi~nih
lastnosti materialov zahteva zanesljivo dolo~ene termi~ne lastnosti, kar bistveno vpliva na natan~nost in zanesljivost simulacij in
njihovih rezultatov. Termi~ne lastnosti PCM-materialov navadno dolo~amo z DSC-metodo, ki temelji na merjenju temperature
in toplote preiskovanega in referen~nega materiala. Obstaja ve~ na~inov, kako modelirati materiale s faznimi premenami. V tem
~lanku je glavna pozornost posve~ena entalpijskemu pribli`ku in metodi efektivne toplotne kapacitete. Obe tehniki, ki
uporabljata rezultate DSC-meritev, omogo~ata uporabo latentne toplote, kar omogo~a u~inkovito simuliranje shranjevanja
toplote ali hladu. Predstavljeni metodi sta bili uporabljeni z uporabo rezultatov iz DSC-meritev za simulacijo son~nega
kolektorja s PCM-materiali. Predstavljeni so rezultati z diskusijo. ^lanek obravnava tudi problem nezanesljivosti podatkov o
lastnostih materialov in njihov vpliv na numeri~no simulacijo.
Klju~ne besede: materiali s faznimi premenami (PCM), entalpijska metoda, metoda u~inkovite toplotne kapacitete, son~ni
kolektor
1 INTRODUCTION
Due to the increasing consumption of energy, the
demands for thermal comfort and the need to decrease
temperature fluctuations and energy losses, new
construction and thermal storage materials have been
developed over the past 10 years. Phase change materials
(PCMs) are substances that are used to absorb, store and
release energy. This is accomplished by utilizing the
latent heat of the phase change from the liquid to the
solid state and vice versa, since the amount of energy
included in a phase change is usually considerably larger
than the sensible heat1. Nowadays, PCMs are being used
in a wide range of applications in architecture and civil
engineering, from heat or cold storage to temperature
stabilization and balancing.
PCMs are usually composed of several substances
rather than being a pure chemical element. The reason is
that a particular application requires a specific range of
phase change according to the operating conditions and a
desired melting temperature can be achieved with a
mixture of several components with a certain chemical
composition.
The optimal layout, usage, integration, and design of
PCMs in buildings, constructions and other equipment
cannot be carried out without using numerical and com-
puter simulations, optimization methods and experi-
mental measurements. In order to obtain accurate and
credible results, it is essential to properly determine all
the physical properties of the used materials and to
precisely describe the operating conditions2. For this
purpose differential scanning calorimetry (DSC) is very
often used.
Materiali in tehnologije / Materials and technology 46 (2012) 4, 335–338 335
UDK 669.017.3:536.65 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)335(2012)
There are several methods and approaches as to how
the materials embracing phase change and related
phenomena can be modeled. In this paper the attention is
focused on the enthalpy approach and the effective heat
capacity method that enable the simulation of the heat
transfer in materials with phase changes. As the inputs
and parameters (e.g., the physical properties of a
material) of the simulation the results from a DSC
measurement may be utilized.
2 DIFFERENTIAL SCANNING CALORIMETRY
Differential scanning calorimetry is a thermo-analyti-
cal method and is widely used to assess and construct
phase diagrams and to determine the thermophysical
properties of materials, e.g., the crystallization and
melting temperatures, heat capacity and the enthalpy3.
The principle of differential scanning calorimetry is a
measurement of the difference in the amount of heat
needed to raise the temperature of the sample and the
reference materials. The aim is to keep the temperature
of both the sample and the reference materials at the
same level during the experiment, while the temperature
increases linearly. In order to precisely determine the
properties and parameters of a sample material, the
thermophysical properties of the reference material have
to be known. The particular outcome of a DSC measure-
ment for a PCM is shown in Figure 1.
3 ENTHALPY METHOD
The numerical modeling of PCMs requires solving
the heat transfer problem in order to obtain the transient
temperature distribution for the examined material. The
heat transfer is described by the heat equation that is a
partial differential equation of the second order:
∂
∂
T
t
T− =Δ 0 (1)
where T/K is the temperature, t/s is the time and
/(m2 s–1) is the thermal diffusivity of the material.
The enthalpy method4,5 introduces the volume
enthalpy H/(J m–3), which includes the phase change
phenomenon, to the equation since the enthalpy
represents the sum of the sensible heat and the latent heat
of the phase change. The function of the volume
enthalpy is defined as follows:
H T c L
f
T
T
( ) = −
⎛
⎝
⎜ ⎞
⎠
⎟∫   f s d
∂
∂0
(2)
where /(kg m–3) is the density, c/(J kg–1 K–1) is the
specific heat, Lf/J is the latent heat of the phase trans-
formation and fs is the solid fraction. By substituting the
enthalpy relation (2) into equation (1) we can obtain the
reformulated equation describing the unsteady heat
transfer in a material with phase changes:
∂
∂
H
t
T= ∇ ∇( ) (3)
where /(W m–1 K–1) is the thermal conductivity. As can
be seen, in equation (3) there are two unknown quanti-
ties – the enthalpy and the temperature. However, they
are coupled together through the enthalpy-temperature
(H-T) relationship that can be obtained from a DSC
measurement. An example of the H-T relationship for a
particular material with a phase change range from
25 °C to 27 °C is shown in Figure 2.
In order to solve equation (3) for a particular
application the corresponding initial and boundary
conditions have to be provided. To solve the system
numerically, an appropriate numerical method is needed
to assemble a system of discretized equations. For this
purpose many numerical methods can be utilized, e.g.,
the finite-difference method, the finite-volume method,
the finite-element method or the mesh-free methods6. For
simplicity, let us assume the finite-difference method.
Thus, the partial derivatives in equation (3) are replaced
by their finite-difference approximations and the original
system is reformulated into a system of linear algebraic
equations:
H H f T Ti j k
n
i j k
n
i j k
n
i x j x k z
n
, , , , , , , ,( , , ...)
+
± ± ±= +
1
Δ Δ Δ (4)
However, due to the presence of unknown values of
both the enthalpy and the temperature, the explicit
scheme of the finite-difference method for the discreti-
zation in time has to be used. The temperature field is
calculated in two successive steps: in the first step the
enthalpy H(n + 1) in time t + 1 is calculated from the
known enthalpy and temperature values Ht, Tt in time t
according to equation (4). After that the temperature
T(n + 1) in time t + 1 is calculated from the enthalpy H(n + 1)
according to the enthalpy-temperature relationship
obtained, for instance, from a DSC measurement.
Repeating the described procedure over all the time
nodes leads to the targeted transient temperature field.
L. KLIMES et al.: CHALLENGES IN THE COMPUTER MODELING OF PHASE CHANGE MATERIALS
336 Materiali in tehnologije / Materials and technology 46 (2012) 4, 335–338
Figure 1: Results of a DSC measurement of a PCM
Slika 1: Rezultati meritev DSC s PCM
4 EFFECTIVE CAPACITY METHOD
The effective capacity method is based on the idea of
incorporating the phase change phenomenon into the
heat capacity calculations4,5. This physical property
represents the amount of heat needed for the matter to
increase its temperature.
It means that the effective heat capacity ceff/(J kg–1 K–1)
has to include the latent heat of the phase change.
Therefore, the function of the effective heat capacity
increases and decreases sharply with an apparent peak
when the material undergoes the phase change. An
example of the effective heat capacity function corres-
ponding to the enthalpy curve is shown in the same
Figure (Figure 2).
In order to calculate congruous and trustworthy
results it is crucial to properly determine the effective
heat capacity. For this purpose differential scanning
calorimetry is one of the tools and techniques available.
By introducing the effective heat capacity, equation
(1) becomes:
 c
T
t
Teff
∂
∂
= ∇ ∇( ) (5)
In equation (5) there is only one unknown variable,
the temperature T, and therefore a wider class of
discretization approaches can be used. One of the most
significant features is that the implicit scheme can be
utilized. Hence, using the temperatures T n from the
foregoing time step for the determination of the effective
heat capacity in a successive time n + 1, the discretized
equation for each node i, j, k can be written in the form:
f T T Ti j k
n
i x j x k z
n
i j k
n( , , ...), , , , , ,
+
± ± ±
+ =1 1Δ Δ Δ (6)
The system of algebraic equations can be assembled
by using equation (6) applied to all the nodes of the
discretized domain and an appropriate numerical method
has to be used to solve it.
5 COMPARISON OF METHODS
The enthalpy method is based on the introduction of
an additional quantity to the governing equation in order
to include the phase change. From the computing point
of view the volume enthalpy function introduces difficul-
ties. Namely, (a) the desired quantity, i.e., the
temperature, cannot be computed directly and the
computation is split into two steps with an intermediate
calculation, (b) the temperature search from a known
enthalpy is demanding and requires a powerful search
algorithm and (c) without additional techniques the
explicit scheme, which is conditionally stable, must be
used. The advantages of the enthalpy method are a better
numerical stability when an explicit discretization
scheme in time is used for both methods and its better
accuracy in comparison with the effective heat capacity
method4,5.
The effective heat capacity method calculates the
temperature field directly without intermediate steps and
it allows the use of an implicit discretization scheme that
is unconditionally stable6. However, without additional
techniques an assumption is required that the effective
heat capacity is calculated using the temperature from a
previous step. In the case that the effective heat capacity
is a function of an unknown time, this makes the system
of equations considerably more difficult to solve, and
therefore the possibility of using an arbitrary time step
(owing to the implicit scheme) is strictly confined.
Moreover, when a problem with a large number of nodes
(a 3D case with a large and intricate geometry) the
computational demands of solving the system of
equations becomes more arduous and time consuming.
6 RESULTS AND DISCUSSION
Both discussed methods were implemented in order
to simulate and analyze the solar air collector, which is a
solar thermal device that is used for heating in buildings.
An issue with solar thermal systems is the general need
for thermal storage in order to balance the heat demands
over a certain time period. Therefore, the utilization of
L. KLIMES et al.: CHALLENGES IN THE COMPUTER MODELING OF PHASE CHANGE MATERIALS
Materiali in tehnologije / Materials and technology 46 (2012) 4, 335–338 337
Figure 3: The results of the numerical simulation based on the
effective heat capacity for the solar air collector
Slika 3: Rezultati numeri~ne simulacije na osnovi efektivne kapacitete
za son~ni kolektor zraka
Figure 2: The enthalpy temperature relationship and the effective heat
capacity function
Slika 2: Sprememba entalpije v odvisnosti od temperature in efektivne
toplotne kapacitete
phase change materials and their feature of energy
storage can overcome the mentioned problem.
The TRNSYS software, which is a simulation tool
designed for the analysis of transient systems, was used
for modeling the solar air collector. Owing to the phase
change material, in the case of the enthalpy method, the
numerical model of heat transfer in a PCM has been
implemented in MATLAB and both software packages,
TRNSYS and MATLAB, were coupled together. In the
case of the effective heat capacity method the numerical
model of the PCM was implemented in the form of a
stand-alone module (a DLL library) for TRNSYS in the
programming language C++.
Both the software implementations utilized a result
obtained from a DSC measurement of the PCM with the
phase change temperature ranging from 30 °C to 38 °C
and the latent heat of the phase change being 70 kJ kg–1.
The simulations performed with the described
methods produced comparable results. Figure 3 shows
the outcomes of the simulation (meteorological data for
24 h in June, using the effective capacity method, the
finite-difference discretization and the explicit scheme,
the time step 2 seconds) for the so-called front-pass solar
air collector that was split longitudinally into five
separate sections. As seen in Figure 3 the courses of the
PCM temperatures show the phase change in the given
temperature ranges, the absorption and the release of the
latent heat.
However, even though the material properties are
properly determined by the DSC there are still quantities
under uncertainty that may considerably influence the
result of the simulations. A typical example is the heat
transfer coefficient (HTC), since it can randomly
fluctuate according to various factors. The HTC belongs
to the class of operating conditions that can significantly
affect the behavior of the system through the initial and
boundary conditions. In order to properly solve the
systems with parameters under uncertainty, additional
techniques and methods, e.g., stochastic modeling and
stochastic optimization, must be considered.
Moreover, the presented methods and their imple-
mentations can be utilized in a wide range of
applications, e.g., for steel solidification and optimal
control in the steel industry7.
7 CONCLUSION
The presented paper introduces two possible
approaches to modeling materials undergoing phase
changes. Both techniques assume some particular
thermophysical properties obtained, e.g., from the
differential scanning calorimetry method. The concepts
of the enthalpy approach and the effective heat capacity
are described and the advantages, drawbacks and the
usage are discussed.
The implemented methods were successfully tested
on the simulation of the solar air collector with a PCM
absorber. The obtained results suggest that the described
approaches of numerical modeling and their implemen-
tations are valuable tools for solving problems in
optimizing the design, usage and the integration of
materials with phase changes.
Acknowledgement
The presented research was supported by the project
OC10051 of the Czech Ministry of Education and by the
BUT project BD13102003 for young researchers. The
main author, the holder of the Brno PhD Talent Financial
Aid sponsored by Brno City Municipality, also gratefully
acknowledges the financial support.
8 REFERENCES
1 F. Kuznik, D. David, K. Johannes, J. J. Roux, A review on phase
change materials integrated in building walls, Renewable and
Sustainable Energy Reviews, 15 (2011), 379–391
2 F. Kuznik, J. Virgone, J. J. Roux, Energetic efficiency of room wall
containing PCM wallboard: A full-scale experimental investigation,
Energy and Buildings, 40 (2008), 148–156
3 G. H. Zhang, C. Y. Zhao, Thermal and rheological properties of
microencapsulated phase change materials, Renewable Energy, 36
(2011), 2959–2966
4 D. M. Stefanescu, Science and Engineering of Casting Solidification,
2nd ed., Springer, 2009
5 M. Muhieddine, E. Canot, R. March, Various approaches for solving
problems in heat conduction with phase change, International Journal
On Finite Volumes, 6 (2009) 1
6 Handbook of numerical heat transfer. Editors W. J. Minkowycz, E.
M. Sparrow, J. Y. Murthy, 2nd ed., Wiley, 2006
7 J. Stetina, F. Kavicka, The influence of the chemical composition of
steel on the numerical simulation of a continuously cast slab, Mater.
Tehnol., 45 (2011) 4, 363–367
L. KLIMES et al.: CHALLENGES IN THE COMPUTER MODELING OF PHASE CHANGE MATERIALS
338 Materiali in tehnologije / Materials and technology 46 (2012) 4, 335–338
D. VOJTÌCH, F. PRÙ[A: APPLICATION OF POWDER METALLURGY IN THE PROCESSING OF ALUMINIUM SCRAPS ...
APPLICATION OF POWDER METALLURGY IN THE
PROCESSING OF ALUMINIUM SCRAPS WITH
HIGH-IRON CONTENTS
UPORABA POSTOPKOV PRA[NE METALURGIJE ZA PREDELAVO
ODPADKOV NA OSNOVI Al Z VISOKO VSEBNOSTJO Fe
Dalibor Vojtìch, Filip Prù{a
Department of Metals and Corrosion Engineering, Institute of Chemical Technology, Technická 5, 166 28 Prague 6, Czech Republic
dalibor.vojtech@vscht.cz
Prejem rokopisa – received: 2011-10-20; sprejem za objavo – accepted for publication: 2012-03-12
The Al-14Si-8Fe and Al-22Si-8Fe-1Cr (w/%) alloys were prepared via melt centrifugal atomization and compaction at a
pressure of 6 GPa. The resulting alloys were porosity-free with good contacts between the particles. The structures were very
fine and consisted of primary (Al) dendrites, eutectic silicon and the 	-Al5FeSi phase. The Vickers hardness values of the
Al-14Si-8Fe and Al-22Si-8Fe-1Cr alloys were (170 ± 10) HV5, (185 ± 16) HV5, respectively. Compressive strength values
were 670 MPa and 720 MPa, respectively. Surprisingly, both alloys exhibited a very good plasticity due to the refined structure.
The alloys were annealed at 300–400 °C for 100 h to reveal their thermal stability. When the thermal stability of the
powder-metallurgy Al-14Si-8Fe and Al-22Si-8Fe-1Cr alloys was compared to that of the casting-piston Al-12Si-1Cu-1Mg-1Ni
alloy, the former was found to be superior.
Keywords: aluminium, powder metallurgy, rapid solidification, iron
Zlitini Al-14Si-8Fe in Al-22Si-8Fe-1Cr (w/%) sta bili pripravljeni s centrifugalno atomizacijo in hladnim enoosnim stiskanjem s
tlakom 6 GPa. Izdelane zlitine so bile brez por z dobrim stikom med delci. Zlitini sta imeli zelo fino mikrostrukturo, ki je
vsebovala primarne (Al)-dendrite, Si-evtektik in fazo 	-Al5FeSi. Trdota zlitine Al-14Si-8Fe je bila (170 ± 10) HV5 in (185 ±
16) HV5 zlitine Al-22Si-8Fe-1Cr. Tla~na trdnost zlitin je bila 670 MPa oziroma 720 MPa. Presenetljivo sta imeli obe zlitini tudi
odli~no plasti~nost zaradi rafinirane strukture. Zlitini smo `arili 100 h pri 300–400 °C, da bi ugotovili njuno termi~no stabilnost.
Ugotovili smo, da je termi~na stabilnost izdelanih P/M-zlitin Al-14Si-8Fe in Al-22Si-8Fe-1Cr superiorna v primerjavi z
ventilsko litino Al-12Si-1Cu-1Mg-1Ni.
Klju~ne besede: aluminij, metalurgija prahov, hitro strjevanje, `elezo
1 INTRODUCTION
Aluminium alloys exhibit a relatively high strength-
to-weight ratio and good corrosion resistance. This
makes them attractive for automotive and aerospace
industries. In the recycling of aluminium, various grades
of the Al-containing scraps are processed. In addition to
water, oils, plastics and other non-metallic contaminants,
the main metallic impurity in the Al scraps is iron (Fe).
The iron contents depend on the actual scrap grade that
can exceed the mass fraction (w/%) 10 %. Unfortunately,
iron very negatively influences the mechanical properties
of Al alloys, because it forms hard and brittle inter-
metallic phases in the structure, as seen in Figure 1.
Therefore, the iron contents in most of the commercial,
wrought, Al-based alloys must be minimized to a few
tenths of w/%. Only die-casting Al-Si alloys can contain
approximately 1 % of Fe, but the negative effect of the
intermetallic phases is balanced with the relatively high
cooling rates and fine structures of the castings. In the
processing of the Fe-containing Al scraps, the iron
contents can be lowered, for example, by magnetic
separation before melting or by melting at carefully
controlled temperatures to prevent dissolution of Fe in
the molten aluminium. These techniques increase the
investment and processing costs. For this reason, the
Fe-containing Al scraps are significantly cheaper than
the Fe-free Al scraps. For example, the price of the scrap
Materiali in tehnologije / Materials and technology 46 (2012) 4, 339–343 339
UDK 621.762:669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)339(2012)
Figure 1: SEM micrograph of the as-cast Al-14Si-8Fe alloy contain-
ing large 	-Al5FeSi plates (light) deteriorating the mechanical pro-
perties
Slika 1: SEM mikroskopski posnetek litine na osnovi Al-14Si-8Fe, ki
vsebuje velike 	-Al5FeSi (svetle) plo{~ice, vzrok za slab{e mehanske
lastnosti
containing up to 10 % of Fe is only one third of that of
the Fe-free scrap.
In addition to the classic ingot and casting metallurgy
of the Al-based alloys, these alloys can also be processed
with powder metallurgy (P/M). It is shown in this study
that the P/M is a method which enables the preparation
of the qualitatively new Al-based alloys containing
transition metals, such as Fe, in the concentrations far
exceeding those in common casting and wrought alloys.1
Therefore, aluminium scraps containing high amounts of
iron can be processed with this technique. A rapid
solidification of a melt, which generally occurs during
the powder preparation, refines the structure, reduces the
volume fraction of the intermetallic phases and forms
new metastable crystalline, quasi-crystalline and
amorphous phases.1 All of these structural features are
beneficial for achieving desirable combinations of
strength and ductility. Powder compaction is generally
performed by pressing and sintering or by hot extrusion.2
2 EXPERIMENT
The P/M alloys with the nominal chemical compo-
sitions (in mass fractions, w/%) of Al-14Si-8Fe and
Al-22Si-8Fe-1Cr are studied in this work. The P/M
alloys are compared with the commercial, Al-Si-based
casting alloy with the nominal composition of
Al-12Si-1Cu-1Mg-1Ni. This alloy is used for the pistons
of combustion engines and is generally considered as
thermally very stable. The P/M alloys were prepared by
melting pure elements and master alloys in a vacuum
induction furnace under an argon atmosphere. After a
sufficient homogenization of the melt was achieved, the
melt was poured into a cast-iron metal mould to prepare
an ingot of 20 mm in diameter and 150 mm in length.
Solidified powder was rapidly prepared with centrifugal
atomization, which is schematically shown in Figure 2.
The alloy ingot was remelted under argon and ejected
onto a rapidly rotating graphite wheel (the rotation speed
of 30 000 r/min) to produce flake-like particles (Figure
3). The dimensions of the powder particles ranged from
0.1 mm to 2 mm with a thickness of approximately 50
μm. The centrifugal atomization technique allows a
production of large quantities of powders at low cost.
The powder was then placed into a tungsten-carbide
mould and was compacted by uni-axial pressing at an
ultra-high pressure of 6 GPa to prepare cylindrical
samples of 15 mm in diameter and 5 mm in height. The
pressing temperature and time were 450 °C and 1 h,
respectively. The casting alloy was provided by an
industrial supplier in the form of an ingot with a
thickness of 20 mm and a length of 100 mm. The alloy
was heat treated by the T6 regime consisting of solution
annealing at 510 °C/5 h, water quenching and artificial
aging at 230 °C/6 h.3
For all of the investigated materials, compressive
mechanical properties were measured with an Instron
5882 machine at a deformation rate of 1 mm/min.
Vickers hardness tests were performed at room
temperature with a 5-kg load. To examine the thermal
stability of the alloys, compressive tests were also
performed after annealing the samples for 100 h at
300–400 °C, and the development of the Vickers
hardness at room temperature was observed to follow the
structural changes induced by the long annealing time.
The structures of the alloys were investigated using light
microscopy (LM), scanning electron microscopy (SEM,
Tescan Vega 3) and energy dispersion spectrometry
(EDS, Oxford Instruments Inca 350). The phase compo-
sition was determined with X-ray diffraction (XRD, X
Pert Pro).
D. VOJTÌCH, F. PRÙ[A: APPLICATION OF POWDER METALLURGY IN THE PROCESSING OF ALUMINIUM SCRAPS ...
340 Materiali in tehnologije / Materials and technology 46 (2012) 4, 339–343
Figure 3: Flake-like powder particles prepared with the centrifugal
atomization
Slika 3: Pra{ni delci v obliki plo{~ic, izdelani s centrifugalno atomi-
zacijo
Figure 2: Schematic drawing of the centrifugal atomization used in
this study to prepare the P/M alloys
Slika 2: Shemati~ni prikaz centrifugalne atomizacije, ki je bila upo-
rabljena v tej {tudiji za pripravo P/M zlitin
3 RESULTS AND DISCUSSION
The structures of the investigated alloys are shown in
Figure 4. The P/M Al-14Si-8Fe alloy (Figure 4a) is
dominated by the α(Al) (dark) and 	-Al5FeSi (light)
phases. XRD (not shown) also proved the presence of a
small volume fraction of eutectic silicon that is hardly
seen in an SEM micrograph in the secondary-electron
regime. The 	-Al5FeSi phase appears as very fine
needles with an average thickness of 2 μm and the length
of 8 μm. The P/M Al-22Si-8Fe-1Cr alloy is shown in
Figure 4b. This alloy has the same phase composition as
the previous one, i.e., it contains the 	(Al) (dark),
	-Al5FeSi (light) and Si (not seen) phases. The EDS
analysis reveals that chromium does not create any new
phases. Instead, it is incorporated in the 	-Al5FeSi phase
where it substitutes iron atoms. In addition, chromium
significantly changes the morphology of the
	-Al5FeSi(Cr) phase from a needle-like to an almost
spherical structure. The observed phase composition of
the P/M alloys corresponds to the well known ternary
equilibrium diagram of the Al-Si-Fe system4 that is
illustrated in Figure 5. The structure of the casting
Al-12Si-1Cu-1Mg-1Ni (T6) alloy in Figure 4c consists
of primary (Al) dendrites (light), (Al)+Si eutectic and
intermetallic phases containing mainly Ni and Al.
During the T6 heat treatment of the alloy, plate-like
Al2CuMgNi precipitates were formed which are seen in
the TEM micrograph in Figure 4c.
By comparing Figures 1 and 4a, one can see that the
centrifugal atomization of the P/M alloys causes a
significant structural refining. This refining is a direct
consequence of the rapid solidification occurring during
the centrifugal atomization of the melt. In this process
(Figure 2), the melt stream falls onto a fast rotating
wheel, where it is broken into small droplets by the
centrifugal force. These droplets are ejected towards the
cooled wall of the atomizer, where they rapidly solidify
and obtain the typical flake-like shape. It will be shown
later that the structural refining has a positive impact on
the resulting mechanical properties. It is also important
to note that the P/M alloys exhibit almost no porosity
and have very good contacts between the particles.
Apparently, sufficient diffusion bonding between the
D. VOJTÌCH, F. PRÙ[A: APPLICATION OF POWDER METALLURGY IN THE PROCESSING OF ALUMINIUM SCRAPS ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 339–343 341
Figure 6: Compressive stress-strain diagrams of the investigated
alloys
Slika 6: Diagram odvisnosti med deformacijo in tla~no trdnostjo
preiskovanih zlitin
Figure 4: Microstructures of the investigated alloys: a) the PM
Al-14Si-8Fe alloy (SEM), b) the PM Al-22Si-8Fe-1Cr alloy (SEM), c)
the casting Al-12Si-1Cu-1Mg-1Ni (T6) alloy (LM, TEM)
Slika 4: Mikrostrukture preiskovanih zlitin: a) P/M Al-14Si-8Fe
(SEM), b) P/M Al-22Si-8Fe-1Cr (SEM), c) litina Al-12Si-1Cu-
1Mg-1Ni (T6) (LM, TEM)
Figure 5: Aluminium corner of the Al-Si-Fe phase diagram4 (the
chemical compositions of the investigated P/M alloys are marked with
black cubes)
Slika 5: Al kot v ternarnem faznem diagramu Al-Si-Fe (kemijske
sestave preiskovanih litin so ozna~ene s ~rnimi kockami)4
particles was induced by the high pressure of 6 GPa used
in this experiment. A similar finding was also reported
by Cieslak et al.5
Figure 6 compares the compressive stress-strain
diagrams of the investigated P/M and casting alloys. It is
observed that the P/M Al-14Si-8Fe and Al-22Si-8Fe-1Cr
alloys exhibit a significantly higher compressive strength
(670 MPa and 720 MPa, respectively) than the
Al-12Si-1Cu-1Mg-1Ni alloy (600 MPa). Moreover, the
P/M alloys also exhibit a higher hardness than their
casting counterparts. Vickers hardness values for the
Al-14Si-8Fe, Al-22Si-8Fe-1Cr and Al-12Si-1Cu-1Mg-
1Ni alloys are (170 ± 10) HV5, (185 ± 16) HV5 and (130
± 8) HV5, respectively. Higher compressive strength and
hardness of the P/M alloys are caused by their observed
structural refining due to the rapid solidification (Figure
4). An additional contribution to both the hardness and
compressive strength is the high volume fraction of the
hard phases, namely 	-Al5FeSi, in the structure of the
P/M alloys (Figure 4). In contrast, the casting alloy
contains large (Al) grains, whose contribution to the
Hall-Petch strengthening is negligible (Figure 4c). The
most pronounced strengthening mechanism operating in
this alloy is the precipitation strengthening caused by the
semi-coherent Al2CuMgNi precipitates (Figure 4c). The
compressive strengths of the P/M alloys investigated in
this study (670 MPa and 720 MPa) are higher than the
recently reported compressive strengths (
500 MPa) of
similar alloys (Al-20Si-5Fe-2(Cu, Ni, Cr)) 6 compacted
at 400 °C/1 h, but at a lower pressure of 300 MPa. This
suggests that the ultra-high-pressure compaction at 6
GPa positively influences the resulting mechanical
properties.
Another important finding is that both P/M alloys
exhibit a very good plasticity, despite the high iron
content (Figure 6). The alloys with the same chemical
composition prepared with the classic casting metallurgy
would be extremely brittle due to the presence of the
large 	-Al5FeSi plates (Figure 1), acting as primary
stress concentrators during the compressive loading. The
plasticity of the P/M alloys can be attributed to the
preparation procedure, including the rapid solidification
D. VOJTÌCH, F. PRÙ[A: APPLICATION OF POWDER METALLURGY IN THE PROCESSING OF ALUMINIUM SCRAPS ...
342 Materiali in tehnologije / Materials and technology 46 (2012) 4, 339–343
Figure 9: Arrhenius plots of diffusion coefficients Ds, of various
metals in Al.7 Transition metals like Fe have significantly lower
diffusion coefficients than the commonly used alloying elements like
Cu, Si or Mg.
Slika 9: Arrheniusov diagram difuzijskih koeficientov Ds razli~nih
kovin v Al.7 Prehodne kovine, kot je Fe, imajo manj{e difuzijske
koeficiente kot pogosti zlitinski elementi v Al-zlitinah (Cu, Si ali Mg).
Figure 7: Room-temperature Vickers hardness as a function of the
annealing time at 300 °C and 400 °C
Slika 7: Vickersova trdota, izmerjena pri sobni temperaturi v odvis-
nosti od ~asa `arjenja pri 300 °C in 400 °C
Figure 8: Compressive strength of the alloys as a function of
annealing regime
Slika 8: Tla~na trdnost izdelanih zlitin v odvisnosti od re`ima `arjenja
of the melt. Due to this procedure, the 	-Al5FeSi parti-
cles became very fine, short or almost spherical (Figures
4a and 4b). As the stress concentration is proportional to
the dimensions of the needles, the refining caused by
rapid solidification is able to reduce the local stresses
and allow the plastic (Al) phase to be deformed before
the formation of fracture micro-cracks.
Figures 7 and 8 illustrate the thermal stability of the
mechanical properties of the investigated alloys. They
show the development of the room-temperature Vickers
hardness and compressive strength of the alloys during
the long-term annealing at 300 °C and 400 °C. Both the
hardness and the compressive strength of the P/M alloys
decrease during the annealing at 300–400 °C, but these
decreases occur much more slowly than in the case of
the casting Al-12Si-1Cu-1Mg-1Ni alloy. The P/M alloys
retain a relatively high hardness and compressive
strength of above 120 HV5 and 450 MPa, respectively,
even after the 100 h annealing at 400 °C. One reason for
a better thermal stability of the P/M alloys is a lower
diffusivity of iron in solid aluminium in comparison with
silicon, copper, magnesium and nickel (Figure 9).7 The
structural coarsening of the 	-Al5FeSi particles is
therefore very slow in the P/M alloys. In contrast, the
growth and transformations of precipitates in the casting
alloy (Figure 4c) occurs fast and, therefore, this alloy
softens rapidly during the annealing.
4 CONCLUSIONS
It is demonstrated in this work that the Al-based
alloys with a high iron content of 8 % can be success-
fully prepared with the powder-metallurgy technology
that consists of a combination of centrifugal atomization
and high-pressure compaction. The resulting alloys have
significantly refined structures that have a positive
impact on strength, hardness and, particularly, plasticity.
The alloys also exhibit excellent thermal stability in
comparison with the casting Al-12Si-1Cu-1Mg-1Ni
alloy commonly used for the production of the com-
bustion-engine pistons. The classic casting processes
cannot be used for the preparation of the high-iron alloys
because iron would form large, needle-like 	-Al5FeSi
particles, resulting in an extreme brittleness of the
as-cast alloys.
Acknowledgements
The authors wish to thank the Czech Academy of
Sciences (project no. KAN 300100801) for its financial
support of this research.
5 REFERENCES
1 D. Vojtìch, J. Verner, J. [erák, F. [iman~ík, M. Balog, J. Nagy, Mat.
Sci. Eng., A458 (2007), 371
2 L. Shaw, H. Luo, J. Villegas, D. Miracle, Scripta Mater., 50 (2004),
921
3 S. Michna, Aluminium Materials and Technologies from A to Z.,
Adin, Presov 2007, 375
4 L. F. Mondolfo, Aluminum Alloys: Structure and Properties, Butter
Worths, London 1976, 535
5 G. Cieslak, J. Latuch, Rev. Adv. Mater. Sci., 18 (2008), 425
6 M. Rajabi, M. Vahidi, A. Simchi, P. Davami, Mater. Charact., 60
(2009), 1370
7 Y. Du, Y. A. Chang, B. Huang, W. Gong, Z. Jin, H. Xu, Z. Yuan, Y.
Liu, Y. He, F. Y. Xie, Mat. Sci. Eng., A363 (2003), 140
D. VOJTÌCH, F. PRÙ[A: APPLICATION OF POWDER METALLURGY IN THE PROCESSING OF ALUMINIUM SCRAPS ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 339–343 343

P. BORKOVI] et al.: FATIGUE-LIFE BEHAVIOUR AND A LIFETIME ASSESSMENT ...
FATIGUE-LIFE BEHAVIOUR AND A LIFETIME
ASSESSMENT OF A DOUBLE-LEAF SPRING USING
FEM-BASED SOFTWARE
UTRUJANJE IN OCENA DOBE TRAJANJA DVOLISTNATIH
VZMETI Z UPORABO MKE-ORODJA
Predrag Borkovi}, Borivoj [u{tar{i~, Milan Male{evi}, Borut @u`ek, Bojan Podgornik,
Vojteh Leskov{ek
Institute of Metals and Technology, Lepi pot 11, SI-10000 Ljubljana, Slovenia
predrag.borkovic@imt.si
Prejem rokopisa – received: 2011-10-20; sprejem za objavo – accepted for publication: 2012-04-05
A lifetime assessment of any component by performing classical fatigue tests on real geometry is a very expensive and
time-consuming task. For this reason, many efforts have been made to produce a computer software which will be able to
replace this conventional fatigue testing. Although there is some dissimilarity between these two testing methods, we are free to
use such a computer tool to shorten the testing time. The Finite Element Method (FEM) is one of the most frequently used
methods for component designing, as well as for fatigue assessment. This assessment includes the stresses, the loads and the
strength behavior of the material component. As the input values for the material properties, S/N curves were defined on the
basis of the test results. The tensile and fatigue tests on the notched, as well as on the unnotched specimens under
compression-tension loading were carried out in order to obtain the static and dynamic (S/N curves) properties of the
investigated spring steel (51CrV4). Two groups of specimens were prepared: the longitudinal and perpendicular specimens
relative to the rolling direction, both with two different heat-treatments (tempering temperatures).
The results of the performed investigation showed a clear difference between the fatigue strengths of the specimens with two
outmost segregation orientations and two different tempering temperatures. In addition, other influences, such as the residual
stresses due to shoot peening and machining, surface roughness and metallurgical variables, can affect the fatigue strength of a
material. All these effects can be taken into account by using the FEM-based simulations.
Keywords: fatigue life, leaf spring, FEM simulation, spring steel
Ocenjevanje dobe trajanja komponent pri izvajanju klasi~nega preizkusa utrujanja realne geometrije je zelo drag in dolgotrajen
postopek. Zaradi tega je bil vlo`en velik napor, da bi pridobili ra~unalni{ko podprto orodje, ki bi bilo sposobno zmanj{ati {tevilo
konvencionalnih preizkusov utrujanja. ^eprav realno lahko pri~akujemo, da obstajajo nekatere razlike med rezultati
preizku{anja na realnih preizku{ancih in tistimi, dobljenimi z ra~unalni{ko podprtim orodjem, je prednost slednjega nedvomna,
saj mo~no skraj{a ~as do pridobitve dokaj uporabnih podatkov. Metode kon~nih elementov (MKE) so med najbolj uporabnimi
za na~rtovanje komponent tudi pri oceni vpliva utrujanja nanje za neko izbrano realno geometrijo. Ta ocena je sestavljena iz
trdnostnih, obremenitvenih in napetostnih zna~ilnosti materiala oz. komponent. Pri vzorcih iz izbranega materiala (vzmetno
jeklo 51CrV4) smo dolo~ili njegovo trdoto ter stati~ne (natezno trdnost) in dinami~ne (S/N-krivulje) lastnosti. Preizkusi
utrujanja oz. dolo~evanje S/N-krivulj je potekalo na zarezanih in nezarezanih vzorcih v natezno-tla~nem na~inu obremenitve.
Pripravljeni sta bili dve skupini vzorcev: v pre~ni in vzdol`ni smeri glede na smer valjanja (smer izcejanja), obe skupini pa sta
bili {e dodatno razli~no toplotno obdelani (pri dveh razli~nih temperaturah popu{~anja).
Rezultati preiskav ka`ejo jasno razliko med dinami~no trajno trdnostjo vzorcev za izbrani orientaciji izcej in izbrani temperaturi
popu{~anja. Poleg tega moramo upo{tevati, da tudi drugi vplivi (metalurgija izdelave jekla, zaostale napetosti zaradi peskanja in
mehanske obdelave, hrapavost povr{ine itd.) lahko vplivajo na dinami~no trajno trdnost materiala. Vsi navedeni vplivi se lahko
upo{tevajo pri uporabi MKE-orodja.
Klju~ne besede: doba trajanja, listnate vzmeti, MKE-simulacije, vzmetno jeklo
1 INTRODUCTION
The basic function of a leaf spring is to store
mechanical energy when it initially elastically deforms
and to recoup this energy after being released. Since a
leaf spring is subjected to dynamic loading it is possible
for a failure to occur even at a stress level considerably
lower than the tensile or the yield strengths determined
by the quasi static loading. In addition, the fatigue
strength of a leaf spring may be influenced by many
factors such as: surface roughness, surface treatment,
component size, residual stresses, corrosion, tempera-
ture, metallurgical structure, etc. In order to provide
stability and safety to a vehicle, all these influences
should to be taken into account in the manufacturing and
designing processes of spring steel, as well as of spring
semifinal products (plates or profiles). Unfortunately, all
of these influencing factors often cannot be accurately
taken into consideration without an exact experimental
investigation.
Particular care has to be taken during a spring-
manufacturing process (profiling, eye making, punching,
etc) that, together with the final heat-treatment, has a
crucial effect on the fatigue behavior of a spring. Among
all the mentioned factors that influence the dynamic
property of a spring, in this paper the influences of the
intensities of the alloying elements, as well as of
different tempering temperatures, on the fatigue behavior
were examined.
Materiali in tehnologije / Materials and technology 46 (2012) 4, 345–349 345
UDK 620.178.3:519.61/.64 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)345(2012)
2 EXPERIMENTAL WORK
The experimental work consisted of the com-
pression-tension fatigue tests carried out on the notched
and smooth (unnotched) specimens on a servo-hydraulic
testing rig ±250 kN INSTRON 8802 at the frequency of
30 Hz. Some additional tests were also performed on a
high-frequency pulsator (Rumul, Switzerland) that is
known for providing a fast and reliable method of
steel-quality assessment1,2. The tested spring steel
51CrV4 was produced with a modified deoxidation
technology at the steel plant [tore Steel, Slovenia. The
chemical composition of the selected spring steel is
given in Table 1.
To determine the dynamic properties of the spring
steel two groups of specimens were prepared and tested.
As it is illustrated in Figure 1, one group of specimens is
oriented along the rolling direction, while another group
is perpendicular to the rolling direction.
Both the perpendicularly ( = 90°) and longitudinally
( = 0°) oriented specimens relative to the rolling direc-
tion are presented in Figures 2 and 3.
For the purpose of both static-tensile test and
dynamic-loading test, all specimens are cut off from the
base spring-steel material in the as-delivered condition (a
flat profile with the dimensions of 90 mm × 28 mm).
After cutting and machining, the specimens were heat-
treated, quenched in nitrogen at 5 bar of overpressure
and then tempered at two different temperatures. Both
perpendicular and longitudinal specimens were divided
into these two groups of tempering temperatures: 425 °C
(HT1) and 475 °C (HT2).
At the end of the experiment, an FEM simulation of
the dynamic loading (fatigue) of a double-leaf spring
with a selected geometry was carried out with the use of
the experimental testing results. Among the many
accessible computer codes based on the stress or
strain-life approaches3–7 and on a cumulative-damage
analysis8 relating to the fatigue-life prediction of the
double-leaf spring, ANSYS computer software was used.
The presented FEM simulation was based on the
high-cycle fatigue approach and the S/N curves for the
investigated spring steel 51CrV4 obtained with the
compression-tension tests of differently oriented speci-
mens at both tempering temperatures. The used com-
puter tool was first successfully tested by performing an
FEM simulation on the smooth and notched cylindrical
specimens during the mechanical testing.
3 RESULTS AND DISCUSION
3.1 Static test
Before the compression-tension fatigue testing, the
static-tensile tests were also carried out using a 500-kN
Instron 1255 test rig with an extensometer. As illustrated
in Table 2 it is obvious that both tensile and yield
strengths increase with a decrease in the tempering
P. BORKOVI] et al.: FATIGUE-LIFE BEHAVIOUR AND A LIFETIME ASSESSMENT ...
346 Materiali in tehnologije / Materials and technology 46 (2012) 4, 345–349
Table 1: Chemical composition of the spring steel 51CrV4 in mass fraction (w/%)
Tabela 1: Kemijska sestava vzmetnega jekla 51CrV4 v masnih dele`ih (w/%)
Chemical element C Si Mn P S Cr Mo Ni V
Composition 0.52 0.35 0.96 0.011 0.004 0.94 0.05 0.13 0.12
Figure 1: Orientation of the specimens with respect to the rolling
direction
Slika 1: Orientacija preizku{ancev glede na smer valjanja
Figure 3: Standard notched cylindrical specimen; Kt = 2 (calculated
with ANSYS)
Slika 3: Standardni cilindri~ni preizku{anec z zarezo; Kt = 2 (izra-
~unan v ANSYS-u)
Figure 2: Standard smooth cylindrical specimen; Kt = 1 (calculated
with ANSYS)
Slika 2: Standardni cilindri~ni preizku{anec brez zareze; Kt = 1
(izra~unan v ANSYS-u)
temperature. The fracture elongation also varies with the
changes in the tempering temperature. When the
tempering temperature decreases the fracture elongation
decreases as well. Ductility is larger in the longitudinal
direction than in the perpendicular one.
3.2 Hardness measurement
A slight difference between the hardness measure-
ments at 425 °C and at 475 °C of the tempering tempe-
rature was also found (Table 3). Before the heat treat-
ment, the base material had the Rockwell hardness of
about 30 HRC, whereas after the heat treatment the
average hardness increased to about 45 HRC at 425 °C
and to 43 HRC at the tempering temperature of 475°C.
Nearly the same results were obtained for both the
perpendicularly and parallelly oriented specimens.
Table 3: Hardness in dependence of the heat-treatment condition and
the segregation orientation
Tabela 3: Trdota v odvisnosti od pogojev toplotne obdelave in
orientacije izcej
Longitudinal ( = 0°) Perpendicular ( = 90°)
425 °C 475 °C 425 °C 475 °C
45.0 HRC 43.4 HRC 45.8 HRC 43.0 HRC
3.3 Dynamic test
The influences of the heat treatment and the segre-
gation orientation under the compression-tension dyna-
mic loading were investigated. These experiments were
performed at the Institute of Metals and Technology
(IMT), Ljubljana, Slovenia, using a compression-tension
testing machine Instron 8802 of ±250 kN with the
operating frequency of 30 Hz.
The S/N curves obtained with the compression-
tension fatigue tests on the notched, as well as on the
smooth specimens are presented in the diagrams in
Figures 4 and 5, respectively. The specimens shown in
Figures 2 and 3 were mechanically polished (fine
metallographic grinding; final paper 800) after the heat
treatment in order to achieve the required surface
roughness.
Figures 4 and 5 show that the fatigue strength of the
specimens tempered at 425 °C is slightly higher than that
of the specimens tempered at 475 °C. It is also clear that
the fatigue strength of the longitudinally oriented
specimens is significantly higher than the fatigue
strength of the perpendicularly oriented specimens. We
still need to finish the fatigue testing on the smooth,
perpendicularly oriented specimens tempered at 475 °C
to get more information about the influence of the notch
on the fatigue strength. A tabular review of the fatigue-
strength dependence on the segregation orientation and
the heat-treatment conditions, for both notched and
smooth specimens, is given in Table 4.
The 5.2 slope of the S/N curve for the perpendicu-
larly oriented HT2 specimens is surprisingly lower than
the 5.9 slope of the HT1-S/N curve. For more accurate
findings, this fatigue test has to be repeated.
P. BORKOVI] et al.: FATIGUE-LIFE BEHAVIOUR AND A LIFETIME ASSESSMENT ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 345–349 347
Table 2: Tensile-test results
Tabela 2: Rezultati nateznega preizkusa
Orientation Temperingtemperature
Yield strength
(MPa)
Tensile strength
(MPa)
Fracture elongation
(%)
Fracture contraction
(%)
Perpendicular
( = 90°)
475 °C/1 h 1373 1448 7.04 24.6
425 °C/1 h 1502 1591 5.16 15.8
Longitudinal
( = 0°)
475 °C/1 h 1366 1442 10.6 41
425 °C/1 h 1502 1606 9.9 42
Figure 5: S/N curves for the smooth, longitudinally orientated speci-
mens
Slika 5: S/N-krivulje nezarezanih preizku{ancev za obe orientaciji
izcej
Figure 4: S/N curves for the notched specimens of both segregation
orientations
Slika 4: S/N-krivulje zarezanih preizku{ancev za obe orientaciji izcej
Table 4: Fatigue limits of the tested specimens
Tabela 4: Trajna dinami~na trdnost preizku{ancev
Segregation
orientation
Tempering
temperature
Notched
specimens
Smooth
specimens
Perpendicular
( = 90°)
HT1 (425 °C) 184 MPa 444 MPa
HT2 (475 °C) 171 MPa 386 MPa
Longitudinal
( = 0°)
HT1 (425 °C) 225 MPa 479 MPa
HT2 (475 °C) 214 MPa 464 MPa
From the fatigue-limit values for the smooth and
notched specimens it is evident that the fatigue limits of
the notched specimens are nearly reduced by the
theoretical stress-concentration factor (Kt = 2 for the
notched specimens).
3.4 Double-leaf-spring simulation using FEM-based
software
Fatigue simulations of the double-leaf spring with the
selected geometry were performed using three standard
loading conditions that the spring manufacturers
normally use for the structural testing of leaf springs
(Figure 7). The fatigue simulations were also based on
the results of the above described mechanical testing.
First, the simulations on the specimens were performed
in order to evaluate the usability of ANSYS software for
the double-leaf spring. All specimens were first
generated in Solid Works 3D computer programme.
After exporting, the specimen models were further
prepared (meshed, constrained, loaded, etc) and finally
simulated with ANSYS computer programme using the
fatigue module. No remarkable difference was found in
the number of cycles between the mechanically tested
specimens and the computer-simulated specimens using
ANSYS software with the corresponding S/N curves.
Afterwards the fatigue simulations using the ANSYS
computer tool were carried out on the double-leaf spring.
One model of the double-leaf spring is presented in
Figure 6.
The ANSYS material database was created using a
new spring-steel material with the required static
properties, forming six S/N curves obtained with the
dynamical testing. The fatigue-life results for the
double-leaf spring obtained with the ANSYS simulation
are gathered in Figure 7. The results are based on the
variation of different parameters, i.e., the segregation
orientation, the heat treatment and the notch effect. The
double-leaf-spring model was loaded in the case of a free
displacement in one direction (z) and constrained in the
other two directions (x and y), Figure 6a.
From the fatigue simulation it is clear that the longest
fatigue life of the double-leaf spring was obtained by
using the dynamic properties of the smooth,
longitudinally oriented specimens, tempered at 425 °C.
For the lowest loading condition it can be seen that all
three S/N curves relating to the smooth specimens give a
fatigue lifetime of more than 5 million cycles, which is
already in the range of the fatigue limit. The S/N curves
obtained with the fatigue testing of the notched
specimens were corrected with the stress-concentration
factor to avoid the influence of the notch.
The expected life of the double-leaf spring is above
1.5 · 105 cycles. It is evident that this expectation is ful-
filled only at the lowest loading condition. Accordingly,
the suspension system of a truck has to be designed in
such a way that this dynamic loading condition is not
exceeded. One can also notice that the highest prescribed
testing conditions are already in the low-cycle fatigue
region (the regime of loading very close to, or even
above, the yield point of the selected spring steel; Table
2). From this point of view the selected structural testing
condition seems problematic and inadequate.
4 CONCLUSION
The investigations have shown scattered results of the
fatigue testing performed at a certain stress level. This
dispersion of the results among the repeated specimens is
P. BORKOVI] et al.: FATIGUE-LIFE BEHAVIOUR AND A LIFETIME ASSESSMENT ...
348 Materiali in tehnologije / Materials and technology 46 (2012) 4, 345–349
Figure 7: Simulated fatigue lifetime of the tested double-leaf spring
Slika 7: Simulacijska doba trajanja preizku{anih dvolistnatih vzmeti
Figure 6: a) Basic dimensions of the double-leaf spring generated
with Solid Works and b) an ANSYS fatigue simulation of the
double-leaf spring
Slika 6: a) Osnovne dimenzije dvolistnate vzmeti, modelirane v Solid
Worksu in b) ANSYS-simulacije utrujanja dvolistnate vzmeti
mainly influenced by different nonmetallic inclusions9 in
the spring steel (MnS, CaS, Al2O3). Insufficient polishing
quality after the heat treatment of the specimen surface
can also lead to the number of cycles being lower than
expected, causing a lower fatigue strength of the
investigated steel. In general, in the case of the selected
spring steel 51CrV4, the fatigue strength of the
perpendicularly oriented specimens decreases by about
25 % more than the strength of the longitudinally orien-
ted specimens, accompanied with the lower tensile and
yield strengths. Considering different stress-concen-
tration factors relating to the notched and smooth
specimens, it is shown that the fatigue strength of the
notched specimens is effectively lower, as indicated by
the stress-concentration factor.
Generally speaking, when dealing with a fatigue-
module simulation, one has to take into consideration not
only the metallurgical and manufacturing effects, but
also the other effects related to the used software, such as
the type and the size of the mesh element (tetrahedral,
hexahedral etc), the input file (IGES, STEP, SLDPRT
etc), the boundary conditions and the constraint type.
5 REFERENCES
1 B. [u{tar{i~, B. Sen~i~, B. Arzen{ek, P. Jodin, The notch effect on
the fatigue strength of 51CrV4Mo spring steel, Mater. Tehnol., 41
(2007) 1, 29–34
2 B. [u{tar{i~, B. Sen~i~, V. Leskov{ek, Fatigue strength of spring
steels and life-time prediction of leaf springs, Assessment of
reliability of materials and structures (RELMAS’2008), St.
Petersburg, Russia, June 17–20, 2008; problems and solutions;
international conference, Volume 1, St. Petersburg, Polytechnic
Publishing House, 2008, 361–366
3 S. Tavakkoli, F. Aslani et al., Analytical Prediction of Leaf Spring
Bushing Loads Using MSC/NASTRAN and MDI/ADAMS, http://
www.mscsoftware.com/support/library/conf/wuc96-/11b_asla.pdf
4 M. S. Kumar, S. Vijayarangan, Static analysis and fatigue life
prediction of steel and composite leaf spring for light passenger
vehicles, Journal of Scientific and Industrial research, 66 (2007) 2,
128–134
5 F. N. A. Refngah, S. Abdullah, A. Jalar, L. B. Chua, Fatigue life
evaluation of two types of steel leaf springs, International Journal of
Mechanical and Materials Engineering (IJMME), 4 (2009) 2,
136–140
6 N. Philipson, Leaf spring modeling http://www.modelica.org/
events/modelica2006/Proceedings/sessions/Session2d1.pdf
7 G. S. S. Shankar, S. Vijayarangan, Mono Composite Leaf Spring for
Light Weight Vehicle – Design, End Joint Analysis and Testing,
Materials Science (Med`iagotyra), 12 (2006) 3, 220–225
8 W. Eichlseder, Enhanced Fatigue Analysis – Incorporating Down-
stream Manufacturing Processes, Mater. Tehnol., 44 (2010) 4,
185–192
9 M. Kova~i~, S. Sen~i~, Critical inclusion size in spring steel and
genetic programming, RMZ – Materials and Geoenvironment,
(2010), 17–23
P. BORKOVI] et al.: FATIGUE-LIFE BEHAVIOUR AND A LIFETIME ASSESSMENT ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 345–349 349

P. GSELMAN et al.: CHARACTERIZATION OF DEFECTS IN PVD TiAlN HARD COATINGS
CHARACTERIZATION OF DEFECTS IN PVD TiAlN
HARD COATINGS
KARAKTERIZACIJA DEFEKTOV PVD TiAlN TRDIH PREVLEK
Peter Gselman1,2, Tonica Bon~ina2, Franc Zupani~2, Peter Panjan1, Darja Kek Merl1,
Miha ^ekada1
1Jo`ef Stefan Institute, Jamova cesta 39, SI-1000 Ljubljana, Slovenia
2University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, SI-2000 Maribor, Slovenia
peter.gselman@ijs.si
Prejem rokopisa – received: 2011-10-23; sprejem za objavo – accepted for publication: 2012-03-07
PVD hard coatings are continuously gaining their importance in different fields of applications. In industrial use, they are often
exposed to corrosive environments. Hard coatings possess inherently good corrosion resistance, but the substrate-hard coating
systems may suffer from a severe corrosion attack due to the defects (craters, pin holes) in the coatings. On the sites, where
defects extend through the coating, pitting corrosion can take place. These sites are drawbacks in the applications of hard
coating.
A PVD TiAlN hard coating was prepared on cold-work, tool-steel (AISI D2) substrates by sputtering using unbalanced
magnetron sources. The growth defects incorporated into the coating were studied after the deposition and corrosion
experiments. We used two methods: (1) scanning electron microscopy (SEM) for general overview of the coating topography
and 2D-characterization of defects, and (2) scanning electron microscopy with a focused ion beam (SEM-FIB) for making serial
cross-sections through the selected defects in order to provide images for a 3D-reconstruction of defects. In this work we tried to
investigate the formation of a defect at a specific location and find out whether the selected defect causes pitting corrosion.
Keywords: PVD hard coating, defect, 3D-reconstruction, FIB, SEM
PVD trde prevleke pridobivajo pomen na razli~nih podro~jih uporabe. Pri razli~nih aplikacijah so pogosto izpostavljene
korozijskemu okolju. Trde prevleke so kemijsko inertne in same po sebi korozijsko obstojne. Med njihovo pripravo pa nastane v
njih veliko defektov (kraterji, "pinholi"). Nekateri od njih segajo do podlage. Na teh mestih pride do jami~aste korozije, kar pa
je z vidika uporabe ne`eleno.
TiAlN-prevleko smo nanesli na podlago orodnega jekla D2 po postopku napr{evanja z neuravnote`enim magnetronskim
izvirom. Defekte, ki so nastali med pripravo plasti, smo preu~ili po korozijskem preizkusu. Za analizo smo uporabili dve
metodi. Vrsti~ni elektronski mikroskop (SEM), s katerim smo pridobili splo{ni pregled prevleke, kot tudi 2D-karakterizacijo
defektov. Kot drugo metodo smo uporabili fokusiran ionski curek, ki je integriran v klasi~ni SEM, s katerim smo izvedli serijo
rezov izbranih defektov za njihovo 3D-karakterizacijo. V delu smo posku{ali na specifi~nih mestih ugotoviti mehanizem
nastanka defektov in ugotovljali, ali povzro~ajo jami~asto korozijo.
Klju~ne besede: trde PVD-prevleke, defekti, 3D-rekonstrukcija, FIB, SEM
1 INTRODUCTION
In hard coatings, macro- and microdefects typically
appear during the deposition1–5. The most common
macrodefects are: (a) large and shallow craters with
diameters greater than 5 μm, (b) nodular defects (sphe-
rical droplets, conical features) with diameters in the
range of 1–5 μm and (c) disk-like holes arising from
droplet wrenching and pin holes1,2. The defects smaller
than 1 μm can be regarded as microdefects. Their main
representatives are: (I.) small micro-sized holes
(so-called pin holes) extended through the entire
thickness of a coating and (II.) small islands generated
by the built-in particles with a diameter of slightly less
than 1 μm. These surface imperfections in the coatings
can cause local stresses, higher friction, sticking of the
material, local loss of adhesion and pitting corrosion. All
these facts are drawbacks in the application of hard
coatings. Therefore, it is very important to reduce their
concentration. However, the question is how to deposit
denser and less defective coatings.
In this paper we tried to investigate the formation of a
defect at a specific location and find out whether such a
defect can cause pitting corrosion. For scientific and
technological reasons, this topic remains very important
also for further investigations and our main goal was to
make a contribution to the researched area by using
microstructural characterization and 3D-reconstruction
of defects.
2 EXPERIMENTAL WORK
The industrial unbalanced sputtering system CC800/7
(CemeCon) was used for the deposition of the TiAlN
hard coating. The sputtering system was equipped with
four unbalanced magnetron sources that were positioned
pairwise at the two sides of the vacuum chamber. A
cold-work, high-chromium, high-carbon-type, tool-steel
D2 (OCR12VM) was used as the substrate material. The
substrates were polished to Ra = 9–15 nm, ultrasonically
cleaned in detergents, rinsed in deionized water and
dried in hot air. The final cleaning was performed by ion
etching. The RF bias with the maximal power of 2000 W
Materiali in tehnologije / Materials and technology 46 (2012) 4, 351–354 351
UDK 621.793:669.058.67:66.019 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)351(2012)
was applied in CC800/7 for 85 min. During the
deposition, the bias voltage and the substrate temperature
were 125 V and 450 °C. The thicknesses of the coatings
were around 4 μm and a 2-fold rotation was used during
the sputtering.
The growth defects integrated into the hard coating
were studied after the corrosion test. The corrosion
experiment was performed in a 0.1 M chloride solution
at pH = 6.8 using a PARSTAT 2263 device for electro-
chemical impedance spectroscopy. The immersion time
was 24 h. The defect morphology and the distribution of
defects were studied with field emission scanning
electron microscopy (SIRION NC400, FEI).
The focused ion beam (SEM-FIB) installed in a
conventional scanning electron microscope (QUANTA
200 3D, FEI) was used to prepare different cross-sec-
tions through the defect. The primary opening (required
for further observations of cross-sections) with the
dimensions of 20 μm × 10 μm × 10 μm was ion milled
with a 20-nA beam current, while the acceleration
voltage of ions was 30 kV. Then the cross-sectioning was
executed with the medium-high beam current (5 nA).
The next step was polishing, which was divided into two
stages. First, there was rough polishing with a 3-nA
beam current and an exposure time of around 3 min,
while the second one was fine polishing with a 1-nA
beam current for another 3 min. At the end of the
polishing, the specimen was tilted and an image of the
cross-section was taken using low ion-beam current (10
pA). After the image acquisition the next slice of sample,
with the thickness of about 1 μm, was removed by ion
milling. The image of the new cross-section was
acquired again by ions. This method was repeated until
the entire area of the defect was examined.
The SEM-FIB cross-section images were put
together into a 3D-image of the defect with the Avizo
Fire 6.3 software. First, the images were imported into
the software, where the voxel size and the slice thickness
were defined. Then the images were arranged and
aligned in such a way that a pixel of any picture
represents the same site of interest on all the images.
After that the voxel was cropped to a volume which
contained only the defect, the coating and the part of the
substrate, where pitting corrosion took place. All the
images were then labeled. During the next step, several
algorithms (resampling, surface generation and smooth-
ing) were used for the 3D-image processing. Finally, a
3D-reconstruction was performed utilizing the same
software6.
3 RESULTS AND DISCUSSION
A corrosion test of TiAlN on a D2 substrate, using
electrochemical spectroscopy, was carried out in a
chloride medium. After that a characterization of defects
was performed.
As a very useful tool for cutting and examining
internal structure of defects, we applied focused-ion-
beam milling. We prepared cross-sections through the
growth defects by ion-beam milling (Figures 1 and 2).
After several cuts of different defects we found out that
with a non-destructive top-view examination (SEM, LM)
it is almost impossible to determine which defect can
cause pitting corrosion, because we cannot observe if
there is any direct diffusion path (pin holes, crevices,
columnar grain boundaries) through the coating that
enables a corrosion attack on the substrate4,7.
To reveal the problem we analyzed two typical
defects. Figure 1 shows a crater that was formed in the
TiAlN coating on the D2 tool steel, while Figure 2
displays a spherical droplet on the same substrate. From
these two images it is difficult to conclude that corrosion
P. GSELMAN et al.: CHARACTERIZATION OF DEFECTS IN PVD TiAlN HARD COATINGS
352 Materiali in tehnologije / Materials and technology 46 (2012) 4, 351–354
Figure 2: SEM-micrograph and an FIB cross-section image of a
spherical droplet on a sample TiAlN/D2 tool steel
Slika 2: SEM- in FIB-slika prereza sferi~ne kapljice na vzorcu
TiAlN/D2 orodnega jekla
Figure 1: SEM image and FIB cross-sectional micrograph of the same
growth defect (crater) on a sample TiAlN/D2 tool steel
Slika 1: SEM- in FIB-slika prereza defekta (kraterja) na vzorcu
TiAlN/D2 orodnega jekla
took place, but when we continue with a series of slices,
we can realize that the crater extends through the whole
coating and originates in a small hole on the surface of
the D2 tool steel (Figure 1). It is known that PVD pro-
cesses have a poor ability to cover small pits or holes
with a low aspect ratio (where depth is similar to their
width) due to the shadowing effect.
Spherical droplets (Figure 2) are created as a result
of arcing. The arc tendency during the deposition or
etching cycles is increased with impurities on substrates,
shields, fixtures, turntables and targets. Arcs are the
origin of microdroplet formation. Some of these micro-
droplets can be incorporated into the coating during its
growth or deposit on the surface of the substrate.
In our case a microdroplet was deposited on the
surface of the substrate material during etching or at the
beginning of the deposition cycle (Figure 2). This
conclusion can be confirmed by the fact that the droplet
is in contact with the uncoated surface and that the
substrate surface under the droplet is untreated by ion
etching due to the shadowing effect (the step on the
surface).
Figure 2 shows that the region under the spherical
droplet is not completely filled with coating due to the
shadowing effect. Unexpectedly, the corrosion did not
occur under it. After a detailed FIB analysis we
recognized that pitting corrosion did not appear because
there were no pin holes or porosities that could allow the
solution to penetrate into the substrate. However, it is
necessary to point out that corrosion was found under the
other examined spherical droplets due to the existing pin
holes.
The 3D image clearly shows that under the crater a
pin hole extends through the entire thickness of the
coating (Figure 3). It is known that corrosion takes place
on the spots, where the solution can directly reach the
substrate. Black pits on the image represent the corroded
volume of the substrate. The calculated volume of the
corroded area is surprisingly large (around 23 μm3),
while the volume of the crater is approximately two
times smaller (12 μm3). The corrosion products had a
depth of 5 μm in the base material. On the spots where
the corrosion has occurred, the substrate cannot provide
sufficient support to the coating, so it can be damaged
easily.
A 3D rendering of a spherical droplet without a
nucleus (microdroplet) is shown in Figure 4.
4 CONCLUSION
Growth defects (large and shallow craters or cone
structures due to the inclusion in the middle of a dense
coating layer) do not always present a break point for the
corrosion resistance. We found that small pits on the
substrate surface and growth defects lead to the
formation of pin holes. It is well known that pin holes
cause pitting corrosion.
Selected pits were analyzed by FIB in combination
with SEM. In this way we obtained an insight into the
pitting-corrosion process.
We can conclude that 3D imaging and analyzing
provide new insights in understanding the growth-defect
formation as well as its influence on the corrosion
processes, which will help us solve challenging problems
regarding the extension of the coating/substrate-system
life time.
Acknowledgement
This work was supported by the Slovenian Research
Agency, ARRS (project L2-2100). Special thanks go to
the company VSG for the free Avizo Fire 6.3 software.
5 REFERENCES
1 P. Panjan, D. Kek-Merl, F. Zupani~, M. ^ekada, M. Panjan, Surf.
Coat. Technol., 202 (2008) 11, 143
2 M. ^ekada, P. Panjan, D. Kek-Merl, M. Panjan, G. Kapun, Vacuum,
82 (2008) 2, 252
3 D. B. Lewis, S. J. Creasey, C. Wüstefeld, A. P. Ehiasarian, P. Eh.
Hovespian, Thin Solid Films, 503 (2006), 143
P. GSELMAN et al.: CHARACTERIZATION OF DEFECTS IN PVD TiAlN HARD COATINGS
Materiali in tehnologije / Materials and technology 46 (2012) 4, 351–354 353
Figure 4: 3D-transparent image of a spherical droplet
Slika 4: 3D-transparentna slika sferi~ne kapljice
Figure 3: 3D reconstruction of a corrosion volume under a TiAlN
coating on a site where a crater with pin holes was present
Slika 3: 3D-rekonstrukcija korodiranega obmo~ja pod TiAlN-prevleko
na mestu, kjer se je nahajal krater s "pinholi"
4 H. A. Jehn, Surf. Coat. Technol., 125 (2000), 212
5 J. Vetter, M. Stuber, S. Ulrich, Surf. Coat. Technol., 168 (2003), 169
6 J. Ohser, K. Schladitz, 3D Images of Materials Structures –
Processing and Analysis, Wiley, 2009
7 S. H. Tsai, J. G. Duh, Journal of The Electrochemical Society, 157
(2010) 5, 89
P. GSELMAN et al.: CHARACTERIZATION OF DEFECTS IN PVD TiAlN HARD COATINGS
354 Materiali in tehnologije / Materials and technology 46 (2012) 4, 351–354
M. SUBAN et al.: DIGITAL IMAGING ANALYSIS OF MICROSTRUCTURES AS A TOOL ...
DIGITAL IMAGING ANALYSIS OF MICROSTRUCTURES
AS A TOOL TO IDENTIFY LOCAL PLASTIC
DEFORMATION
DIGITALNA ANALIZA POSNETKOV MIKROSTRUKTUR KOT
ORODJE ZA UGOTAVLJANJE LOKALNE PLASTI^NE
DEFORMACIJE
Marjan Suban, Robert Cvelbar, Borut Bundara
Institute of metal constructions, Mencingerjeva 7, 1001 Ljubljana, Slovenia
marjan.suban@imk.si
Prejem rokopisa – received: 2011-10-25; sprejem za objavo – accepted for publication: 2012-03-22
This paper presents a methodology to detect plastic deformation at the micro level. The analysis is based on the statistical data
describing the morphological and crystallographic textures of a sample microstructure. This data was obtained from optical
microscopy using a digital imaging analysis. The important parameters necessary to describe the microstructure were identified
as the grain-size and grain-orientation distributions. A change in the weighted product of these two parameters, the grain size as
the area of grains and the grain orientation as the moment of inertia of grains, can represent a measure to identify plastic
deformation on a small area. A demonstration of its applicability was performed on a real object as part of a
ruptured-pipe-failure analysis in a thermal-power-plant boiler. The presented analysis leads to a fast identification of the local
plastic deformation and, in the case of periodical analysis of the same sample, it can even be used as a measure to identify
creeping.
Keywords: digital imaging analysis, plastic deformation, grain orientation, local-deformation analysis
V prispevku je predstavljena metodologija za odkrivanje plasti~ne deformacije na mikroravni. Metoda temelji na podlagi
statisti~nih podatkov, ki opisujejo morfolo{ke in kristalografske teksture mikrostrukture. Vir podatkov je digitalna analiza
posnetkov, pridobljenih iz opti~ne mikroskopije. Kot pomembna parametra, potrebna za opis mikrostrukture, sta opredeljena
distribucija velikosti zrn in usmerjenosti zrn. Sprememba tehtanega zmno`ka teh dveh parametrov: velikosti zrna kot povr{ine
zrna in usmerjenosti zrna kot vztrajnostni moment zrna, je lahko merilo za identifikacijo plasti~nih deformacij na majhnem
podro~ju. Za prikaz uporabnosti metode je bila izvedena analiza na realnem objektu kot del analize po{kodb po~ene cevi v kotlu
termoelektrarne. Predstavljena analiza omogo~a hitro identifikacijo lokalne plasti~ne deformacije, pri ~asovno zaporedni analizi
istega vzorca pa se lahko celo uporablja za identifikacijo lezenja.
Klju~ne besede: digitalna analiza posnetkov, plasti~na deformacija, orientacija kristalnega zrna, analiza lokalne deformacije
1 INTRODUCTION
The quantification of a material microstructure is
important for an evaluation of a material and some of its
properties. Many properties of the material are strongly
influenced by the size and shape of the grains. Obser-
vations of the changes in the geometric features of the
grains may be quantified by an image analysis. Micro-
scopy is a powerful non-invasive tool for studying a
material microstructure, especially if complemented with
an image analysis.
The paper presents the possibilities of evaluating
plastic deformation based on microstructural grain
layouts. Contrary to the other researchers’ work, our
methodology of identifying the grain shape and
evaluating their orientation is based on the moments of
inertia of individual grains. When a load is applied to a
material, it will cause the material to change its shape.
The observed effects of the deformation process on the
material’s microstructure using optical microscopy1 can
be evidently seen on Figure 1. The changes in the
microstructure can be easily described for the grain level
by using three parameters: the change in the grain area,
the shape (elongation) and the orientation. The new pro-
posed methodology for evaluating these three parameters
is presented in the following sections.
1.1 Digital imaging
There are many imaging techniques available for
viewing a microstructure in 2D and thus useful for
collecting images of each cross section. These imaging
methods include the techniques like secondary electron
(SE), back-scattered electron (BSE), ion-induced secon-
dary electron (ISE) and electron back-scatter diffraction2
(EBSD) (Figure 2). While each method generates an
image in 2D, important issues are to be discussed with
regard to some of the techniques3.
Many laboratories are unfortunately limited to the
use of the optical microscopy as an essential and the only
available method for evaluating microstructures. If grain
boundaries are very clearly defined, the computer
programs can be written in such a way that they allow
higher-dimensional measurements (the area and shape
measurements) recorded on digital images, but often
many grain boundaries are hard to distinguish. The
Materiali in tehnologije / Materials and technology 46 (2012) 4, 355–359 355
UDK 539.374:620.186 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)355(2012)
current state of the optical recognition software offers
many tools to improve image resolutions, so that grain
boundaries can be easily recognized.
1.2 Quantitative characterization
The quantitative characterization is an obvious tool
for the researchers to use when attempting to establish a
relationship between a microstructure and its properties.
As a result of many studies numerous different methods
and rules have been developed. In reality the application
of the characterization greatly dictates the nature of the
measurement technique chosen. Often the measurements
are performed only in 2D using optical microscopy and
digital imaging. In order to maximize the value of the
data, microstructural quantification should be designed
effectively. Almost an infinite number of parameters and
correlations can be used to describe a microstructure, but
just some of these parameters are actually important.
The measurement of the grain size is the most used
method of all the techniques for quantifying the micro-
structural features. It is known that the grain size of a
polycrystalline material is extremely important for
determining its properties. Various properties exhibit a
correlation with the grain size, such as: yield stress,
ductility, and hardness4. Generally, the grain size is
measured as an average scalar value, such as the inter-
cept length, the grain area or the grain volume5.
1.3 Grain shape and the principal-axis orientation
The shape of grains is likely to be of major signifi-
cance in a number of applications, but irregular geome-
tries of the grains in a polycrystalline microstructure
make the grain shape a difficult parameter to quantify.
The difficulty lies in the need for the data to express the
true grain shape. Usually a number of simplifying
assumptions are made.
Like the calculation of the grain size, the determi-
nation of the grain shape has been greatly aided by the
EBSD maps. The boundaries of each grain can be clearly
found and measured by identifying the local changes in
the orientation. A problem may occur when two neigh-
boring grains have the same orientation. The image
analyzer can consider these two grains as one grain and
that can lead to an error in determining the grain size and
its shape. The difference between quantifying a grain
shape and a grain size is related to the inability to clearly
describe the shape. Determining a grain area is a rela-
tively straightforward measurement yielding a scalar
value, while a shape really needs to be described by the
local curvatures, which is more complicated and requires
higher-order mathematical descriptors. It is most
common in everyday work to use simple objects to
describe a shape, such as an ellipse. In this case, a
M. SUBAN et al.: DIGITAL IMAGING ANALYSIS OF MICROSTRUCTURES AS A TOOL ...
356 Materiali in tehnologije / Materials and technology 46 (2012) 4, 355–359
Figure 2: Analysis of the microstructure with: a) optical microscopy, b) secondary electron microscopy image, c) EBSD3
Slika 2: Analiza mikrostrukture: a) opti~na mikroskopija, b) SEM, c) EBSD3
Figure 1: Microstructures of the C-Mn steel: a) the initial state; b) the
microstructure after a deformation1
Slika 1: Mikrostruktura C-Mn-jekla: a) za~etno stanje; b) mikrostruk-
tura po deformaciji1
quantification of the grain shape is done with the calcu-
lation of the length of an equivalent ellipse major Emax
and minor axis Emin, where the area of a grain is
equivalent to the area of the approximated ellipse. The
ratio of the minor axis of the equivalent ellipse to its
major axis represents the measure for grain roundness.
The roundness is a measurement of the length – width
relationship, with a value in the range from 0 to 1. A
perfect round grain has a roundness of 1, while a very
narrow, elongated grain has a roundness of about 0.
In addition, many shape descriptors involve com-
bining groups of size parameters to generate dimension-
less values like length/width (aspect ratio), area/convex
area (solidity) and length/fiber length (curl)6, which can
sometimes be ambiguous. Some scalar parameters, such
as the moments of inertia, can do an adequate, or a
better, job of expressing the grain shape and orientation.
The principal-axes orientation quantifies the orien-
tation of the grains’ principal axes relative to the global
coordinate system of the cross section. It is important to
distinguish between the principal-axes orientation, which
is the orientation referring to the principal axes and the
crystal-lattice (crystallographic) orientation, which is
usually determined with EBSD. In the case of an ellipse,
the orientation is presented as angle  measured counter-
clockwise from the horizontal axis x to the axis of lowest
moment of inertia x1, as shown on Figure 3.
2 METHODOLOGY
Generally, the microstructure images suffer from the
defects of improper illumination, artifacts and noise that
are developed at the time of the sample preparation. The
first stage is important for attaining higher grain-segmen-
tation accuracy. An RGB microstructure image was
processed using the tools of digital imaging with the
final goal to achieve an image without any noise,
showing a set of individual grains and excluding the
border grains. This is also the most important step for
analyzing the differentiation between individual grains in
the microstructure. If the grains are not clearly
distinguished then the next steps of the methodology are
purposeless. An application called the Particle Analysis,
or the Blob Analysis, in the National Instruments IMAQ
Vision software was used for digital imaging and
calculating. The blob analysis is the process of detecting
and analyzing distinct two-dimensional shapes within a
region of the image. It can provide information about the
number of grains, location, shape, area, perimeter, and
the orientation of grains.
One of main results provided by the analyzing
software is the moment of inertia, or the second moment
of area, which is a property of a cross section and can be
used to predict the resistance of the cross-section areas to
the bending and deflection around an axis lying in the
cross-sectional plane. The moments of inertia for any
cross section defined as a simple polygon on the x – y
plane can be computed in a generic way by summing the
contributions from each segment of the polygon. The
equations marked as 1, 2 and 3 can be used to calculate
the moments of inertia, where the parameters xi and yi
represent the distance from the origin to the elementary
triangle center point and ai represents the area of this
triangle.
( )I y y y y ax i i i i i
i
n
= + ++ +
=
−
∑1
12
2
1 1
2
1
1
(1)
( )I x x x x ay i i i i i
i
n
= + ++ +
=
−
∑1
12
2
1 1
2
1
1
(2)
( )I x y x y x y x y axy i i i i i i i i i
i
n
= + + ++ + + +
=
−
∑1
24
2 21 1 1 1
1
1
(3)
From these values that represent the moment of
inertia for the x – y axis, the moments of inertia for the
principal axis x1 – y1 can be calculated (for the axis
presentation see Figure 3).
I
I I I I
Ix y
x y x y
xy1 1
2
2
2 2,
=
+
±
−⎛
⎝
⎜
⎞
⎠
⎟ + (4)
In this case Ix1 represents the minimum eigenvalue of
the moment of inertia, while Iy1 is the maximum value.
The roundness of a grain can then be calculated using the
ratio of these two values and it ranges from 0 (an
extremely elongated grain) to 1 (a round grain).
R
I
I
x
y
= 1
1
(5)
The angle of the rotation of the coordinate system
around the grain center of mass  can be calculated
using Eq. 6. It is used as a parameter to quantify the
orientation of an individual grain. A correction of angle
 is needed to transform the angle range from 0–180° to
M. SUBAN et al.: DIGITAL IMAGING ANALYSIS OF MICROSTRUCTURES AS A TOOL ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 355–359 357
Figure 3: Schematic representation of the grains approximation with
the equivalent ellipse
Slika 3: Shematski prikaz aproksimacije kristalnega zrna z njemu
enakovredno elipso
0–90° (e.g., the grain orientation for the angle of 1° is
almost equal to the one for the angle of 179°).
 =
−
1
2
2
atan
I
I I
xy
x y
(6)
3 EXPERIMENTAL RESULTS AND DISCUSSION
For the experimentation we have used many micro-
structure images of low-carbon and austenitic stainless
steel at various resolutions (i.e., magnifications) and with
different grain shapes. These images were obtained with
optical microscopy. An example of a microstructure
image (500-times) of the austenitic stainless steel and the
steps of digital imaging are shown in Figure 4.
On the left-hand side of Figure 5 the results for the
examined sample from Figure 4 are shown. From the
graph it can be concluded that the roundness of the
grains in the microstructure is around 0.46 and that the
median of the grain-orientation angle is approximately
15°, while the average value is 21°. In short, it can be
said, that the grains are elongated in the direction close
to the horizontal axis.
The results presented on the right-hand side of
Figure 5 show a comparison between the roundness
calculated by using the ratio of the equivalent-ellipse
axis and the roundness calculated by using the ratio of
the moment of inertia for the principal axis. It is
observed that the latter is much more sensitive to the
grain shape then the former. The range of the calculated
results is also much wider, which is a more authentic
description of the real state.
4 CASE STUDY
For the practical demonstration, two samples taken
from the ruptured steam pipe, marked as No. 2 and No. 4
(200-times), were analyzed using the above-described
methodology. The average value of the roundness cal-
culated by using Eq. 5 for sample No. 4 is 0.4, while it is
0.36 for sample No. 2. As it can be seen from Figure 6,
the grains are more elongated in the case of sample No. 4
than in the other case. A similar difference can also be
found when observing the grain-orientation angles. The
grain-orientation angle distribution for sample No. 4 is
M. SUBAN et al.: DIGITAL IMAGING ANALYSIS OF MICROSTRUCTURES AS A TOOL ...
358 Materiali in tehnologije / Materials and technology 46 (2012) 4, 355–359
Figure 4: Schematic representation of the steps of digital imaging and
the calculated result for one grain
Slika 4: Shematski prikaz faz digitalne obdelave posnetkov in rezultat
izra~una za eno kristalno zrno
Figure 6: a) Positions of samples No. 4 (left) and 2 (right) of the
ruptured steam pipe, b) the corresponding microstructures and c) the
results of the analysis
Slika 6: a) Polo`aj vzorcev {t. 4 (levo) in 2 (desno) na po~eni cevi par-
nega kotla, b) pripadajo~i mikrostrukturi ter c) rezultati analize
Figure 5: a) Correlation between the roundness and the orientation
angle  and b) the correlation between different calculations of the
roundness
Slika 5: a) Korelacija med okroglostjo in kotom usmerjenosti  in
b) korelacija med razli~nima metodama izra~una okroglosti
almost flat and its median is 45°, which means that these
grains have no orientation tendency. In the second case
the median of the grain-orientation angle is 64°, while
the average value is 59°. This can also be observed from
the grains in the microstructure on the right-hand side of
Figure 6.
5 CONCLUSIONS
In this paper, a semi-automatic, digital, microstruc-
ture-image-analysis system for quantifying a micro-
structure is developed. As an input in the analysis
process, the image obtained with optical microscopy was
used. In the calculation part of the investigation, the
moment of inertia was introduced as the major parameter
to quantify the grain shape and the orientation. The
results obtained with the above method are reproducible
and repeatable, but can be misleading if the differen-
tiation between individual grains in the microstructure is
not clear. We can demonstrate the accuracy and
usefulness of this method by comparing the computed
values with the results of the other methods. The
presented case study of the sample of a ruptured steam
pipe demonstrates its usefulness for the laboratory and
industrial applications in the field of material investi-
gations.
6 REFERENCES
1 R. Song, D. Ponge, D. Raabe, Texture evolution of an ultrafine
grained C-Mn steel and its evolution during warm deformation and
annealing, Acta Materialia, 53 (2005) 18, 4881–4892
2 S. Xia, B. Zhou, W. Chen, Grain Cluster Microstructure and Grain
Boundary Character Distribution in Alloy 690, Metallurgical and
materials transactions A, 40 (2007) 12, 3016–3030
3 M. Groeber, Development of an automated characterization-repre-
sentation framework for the modeling of polycrystalline materials in
3D, Dissertation, Ohio State University, 2007
4 R. W. Armstrong, I. Codd, R. M. Douthwaite, N. J. Petch, Plastic
Deformation of Polycrystalline Aggregates, Philosophical Magazine,
7 (1962), 45–58
5 L. Ciupinski, B. Ralph, K. J. Kurzydlowski, Methods for the charac-
terization of grain size, Materials Characterization, 38 (1997) 3,
177–185
6 S. Ghosh, D. Dimiduk (Eds.), Computational Methods for Micro-
structure-Property Relationships, 1st Edition, Springer New York,
2011
M. SUBAN et al.: DIGITAL IMAGING ANALYSIS OF MICROSTRUCTURES AS A TOOL ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 355–359 359

L. KRAJNC et al.: THERMODYNAMIC ANALYSIS OF THE FORMATION OF NON-METALLIC INCLUSIONS ...
THERMODYNAMIC ANALYSIS OF THE FORMATION
OF NON-METALLIC INCLUSIONS DURING THE
PRODUCTION OF C45 STEEL
TERMODINAMI^NA ANALIZA NASTANKA NEKOVINSKIH
VKLJU^KOV PRI IZDELAVI JEKLA C45
Luka Krajnc1, Grega Klan~nik2, Primo` Mrvar2, Jo`ef Medved2
1[tore Steel, d. o. o., @elezarska 3, 3220 [tore, Slovenia
2University of Ljubljana, Faculty of Natural Science and Engineering, Department for Materials and Metallurgy,
A{ker~eva 12, 1000 Ljubljana, Slovenia
luka.krajnc@store-steel.si
Prejem rokopisa – received: 2011-11-11; sprejem za objavo – accepted for publication: 2012-03-01
C45 steel belongs to the group of carbon steels and is used in the normalized and tempered states for heavily stressed parts in
the automobile industry. Nowadays, a major problem in steelmaking is non-metallic inclusions, which can form during various
steps of the steelmaking process and are detrimental to the mechanical properties of the steel.
In scope of this work we have tried to determine during which steps of the steelmaking process the non-metallic inclusions are
formed. Samples were taken from three different steps of the steelmaking process, i.e., from the electric arc furnace, from the
ladle furnace and from the tundish. The samples were then chemically analysed. The results were used for a thermodynamic
simulation with Thermo-Calc. Other samples were prepared from another simultaneously taken probe and were partly subjected
to differential scanning calorimetry (DSC) and partly metallographically prepared. Photographs were taken with a light
microscope, from which the phase composition was calculated. The non-metallic inclusions in our samples were analysed with
EDS – photographs were taken and point and mapping analyses were made on them.
We found that in our C45 steel sample, spinel (MgO · Al2O3) and aluminate (Al2O3) inclusions can be found at the end of the
ladle-furnace treatment. In the tundish sample only small aluminate and spinel inclusions and a lot of MnS inclusions were
found.
Keywords: non-metallic inclusions, thermodynamics, C45 steel
Jeklo C45 je ogljikovo jeklo, ki se v normaliziranem in pobolj{anem stanju uporablja za obremenjene dele v avtomobilski
industriji. Slab{e mehanske lastnosti in kraj{a uporabna doba pa so predvsem povezane z nekovinskimi vklju~ki. Le-ti lahko
nastanejo pri razli~nih stopnjah procesa izdelave jekla.
V okviru tega dela smo posku{ali ugotoviti, pri katerih stopnjah procesa izdelave jekla nastanejo nekovinski vklju~ki. V ta
namen smo vzeli vzorce pri treh stopnjah izdelave jekla: na elektrooblo~ni pe~i, na ponov~ni pe~i in iz vmesne ponovce na
napravi za kontinuirno litje jekla. Naredili smo kemi~no analizo vzorcev, ki smo jo uporabili za termodinami~ni ravnote`ni
izra~un s programskim orodjem Thermo-Calc. Vzorce smo razrezali in naredili diferen~no vrsti~no kalorimetrijo (DSC),
metalografsko pripravljene vzorce pa smo slikali s svetlobnim mikroskopom. Za kvalitativno metalografsko analizo smo
uporabili program analySIS 5.0 in z njim ugotovili dele` mikrostrukturnih sestavin. Vzorce smo pregledali in slikali {e z
vrsti~nim elektronskim mikroskopom, naredili smo tudi to~kovno analizo vklju~kov in dolo~ili porazdelitev elementov glede na
povr{ino.
Ugotovili smo, da so se v na{em vzorcu jekla C45 pojavljali pri koncu obdelave na ponov~ni pe~i {pinelni (MgO · Al2O3) in
aluminatni (Al2O3) vklju~ki, ki so ne`eleni, trdi in krhki vklju~ki. Pri litju jekla na napravi za kontinuirno litje so se v na{em
vzorcu v jeklu pojavljali le majhni {pinelni in aluminatni vklju~ki, videli pa smo mnogo MnS-vklju~kov.
Klju~ne besede: nekovinski vklju~ki, termodinamika, jeklo C45
1 INTRODUCTION
With the advances in technology, with the
introduction of new secondary-steelmaking processes
and with the increasingly higher quality demands from
buyers, there has been, especially in recent years, a
greater demand for better mechanical properties from
steel. This can only be guaranteed, however, if we can
control the quantity, size and distribution of the
non-metallic inclusions in the steel.
The non-metallic inclusions in steel are formed from
metal elements, such as iron, manganese, silicon,
aluminium and calcium and non-metal elements such as
oxygen, sulphur, nitrogen and phosphor. The
non-metallic inclusions are divided in two groups, i.e.,
external and internal. The external inclusions are formed
because the melt is in contact with slag and refractory
and some parts of these two can be caught in the melt.
These inclusions are large, with an irregular shape and
they are detrimental to the mechanical properties of the
steel. Internal inclusions, on the other hand, are formed
with chemical reactions between the melt, slag and
refractory. These are, in general, smaller. The biggest
problem comes from hard and brittle oxide inclusions1,2.
In the scope of this work we have focused our
attention on the inclusions that are formed during the
de-oxidation process with aluminium, on the
modification of alumina inclusions and on the formation
of spinel-type inclusions.
Steel can, after oxygen rafination, have as much as
0.1 % mass fractions of oxygen; therefore, it is essential
Materiali in tehnologije / Materials and technology 46 (2012) 4, 361–368 361
UDK 669.18 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)361(2012)
to have the steel deoxidized3. To achieve this we add
different materials with a high affinity for oxygen, such
as aluminium, silicon, etc. The process of de-oxidation is
based on the following chemical reaction4:
x [Me] + y [O] = (MexOy)
slag (1)
where [Me] is the dissolved metal in the liquid steel, [O]
is the dissolved oxygen in the liquid steel and (MexOy)
is the product of the reaction, an inclusion which is
reduced to slag.
Aluminium forms with oxygen alumina (Al2O3)
inclusions, which are large, hard and brittle and are
detrimental to the mechanical properties of the steel.
Furthermore, they have a negative effect on the casta-
bility of the steel, because they clog the nozzles from the
tundish in the continuous casting machine. With
aluminium de-oxidized steel the addition of calcium
modifies the alumina inclusions. The results of this
reaction are calcium alumina inclusions (CaO · Al2O3),
which have a globular shape and their melting point is
lower than the melting point of the steel5. They are
normally larger than non-modified inclusions, however,
and it is therefore essential that they are removed into the
slag. Studies have shown that when we add too much
calcium and the concentration of sulphur is high enough,
inclusions of calcium sulphide (CaS) are formed, which
also has a negative influence on the steel’s castability.
Thus it is important to add the correct amount of calcium
to the steel, to be within the calcium concentration range,
when most of the alumina inclusions are modified and
the calcium content is not high enough for the CaS
inclusions to form6.
1.1 The formation of spinel phase
After de-oxidation and the formation of alumina
inclusions, aluminium has a high activity as an element,
because of the higher concentration in the steel and can
reduce magnesium oxide from magnesite refractory and
slag7. Magnesium forms with oxygen magnesite (MgO)
inclusions. The increase in the concentration of MgO
and Al2O3 increases the driving force for the formation
of spinel-type inclusions in the calcium alumosilicate
inclusion systems. Spinel-type crystals can grow almost
as large as crystals of base calcium silicate.
This theory, which was suggested by Park et al.8,
shows the formation of spinel-type inclusions in large
CaO-SiO2-MgO-Al2O3-type inclusions with a low
concentration of oxygen during the transportation of the
melt from the ladle furnace to the continuous casting
machine.
Another theory about the formation of spinel-type
inclusions, which was suggested by Jiang et al.9, shows
us a different mechanism. As a result of de-oxidation
alumina inclusions are formed. During the reaction
between the melt, slag and refractory the magnesium is
reduced from the slag or refractory:
2[Al] + 3(MgO)slag or refractory = 3[Mg] + (Al2O3) (2)
Before the addition of calcium for modifying the
alumina inclusion, the activity of magnesium is higher
than that of calcium, and therefore MgO · Al2O3 and
MgO inclusions are formed:
3[Mg] + 4(Al2O3) = 3(MgO · Al2O3)
inclusion + 2[Al] (3)
[Mg] + [O] = (MgO)inclusion (4)
After the treatment in the ladle furnace and after the
addition of calcium, its activity is increased and the
CaO·MgO · Al2O3 inclusions are formed:
4[Ca] + 5(MgO · Al2O3)
inclusion =
= 4(CaO · MgO · Al2O3)
inclusion + [Mg] (5)
The purpose of this work was to find out which
inclusions are formed in C45 steel and at what point in
the steelmaking process they are formed. The study was
made on C45 steel as a typical representative of the steel
used for heavily stressed parts in the automobile
industry.
2 EXPERIMENTAL
Samples of carbon C45 steel were taken at the
company [tore Steel, d. o. o. Six samples were taken in
all: the first one from the electric arc furnace, the next
four from the ladle furnace and the final one from the
tundish on the continuous casting machine. For the
sampling we used the Sample-on-Line method or the
Lollipop-Sampling method. The samples were taken at
the same time as the samples for the company and they
were marked from 1 to 6.
In [tore Steel, d. o. o., the chemical analysis was
made on samples that were taken in parallel with our
samples. Optical emission spectroscopy was used with
the instrument, made by Spectro, LAVMC12A. The
samples taken for further analysis were first cut, and then
the smaller part was prepared for simultaneous thermal
analysis (STA), while the larger part was used for the
metallographic analysis, as shown in Figure 1. The
chemical analysis of the slag was performed using a
standard chemical analysis.
L. KRAJNC et al.: THERMODYNAMIC ANALYSIS OF THE FORMATION OF NON-METALLIC INCLUSIONS ...
362 Materiali in tehnologije / Materials and technology 46 (2012) 4, 361–368
Figure 1: Cutting and preparing the samples
Slika 1: Razrez in priprava vzorcev
The characteristic temperatures were determined with
a simultaneous thermal analysis (STA), which was made
on a Jupiter 449c from NETSCH. For the reference an
empty corundum crucible was used. The measurements
were made using a dynamic argon flow.
The thermodynamic predictions were performed with
Thermo-Calc TCW5 for the prediction of the solidi-
fication. The calculations were made using the TCFE3,
SSUB3 and SLAG1 databases.
The metallographic analysis was made using a scan-
ning electron microscope JEOL JSM – 5610, equipped
with energy-dispersive spectroscopy (EDS). The light
microscopy was made using OLYMPUS SZ61 and
OLYMPUS BX61 microscopes. Finally, the micrographs
were processed with the Analysis 5.0 computer program.
3 RESULTS AND DISCUSION
3.1 Chemical composition
The chemical composition of the C45 steel samples
is given in Table 1. The chemical composition of the
slag sample, which was taken at the end of ladle furnace
treatment, is given in Table 2.
3.2 Thermodynamic calculations
The isopletic phase diagrams were predicted with
Thermo-Calc. The chemical composition was taken for
the calculations as shown in Table 1. The following
elements, C, Si, Mn, Cr, Al, Mo, Ni and Fe, were taken
into account; however, oxygen was not taken into
account. In Figure 2 an example of a calculated isopletic
phase diagram for iron – carbon is shown, for sample 6,
taken from the tundish on the continuous casting
machine, with the correct carbon content marked.
In Figure 3 the weight fraction of phases in depen-
dence of the temperature is shown for sample 6, taken
from the tundish on the continuous casting machine.
In Figure 3, in the sample taken from the tundish on
the continuous casting machine, it is clear that initially a
small fraction of -ferrite solidifies at 1490 °C, then the
-ferrite begins its transformation into austenite at 1489
°C, and the rest of the melt begins to solidify at 1421 °C.
At 761 °C the austenite begins its transformation into
ferrite and later during the eutectoid phase transition into
pearlite (Fe + Fe3C). At lower temperatures the carbides
precipitate from the austenitic matrix.
L. KRAJNC et al.: THERMODYNAMIC ANALYSIS OF THE FORMATION OF NON-METALLIC INCLUSIONS ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 361–368 363
Figure 2: Isopletic phase diagram for iron–carbon for sample 6, with
our carbon content marked
Slika 2: Navpi~ni prerez faznega diagrama `elezo-ogljik za vzorec 6,
z ozna~eno vsebnostjo ogljika
Table 1: Chemical composition of the C45 steel samples in mass fractions, w/%
Tabela 1: Kemijska sestava vzorcev jekla C45, w/%
Chemical element C Si Mn P S Cr Mo
Sample 1 0.13 0.03 0.13 0.011 0.036 0.11 0.04
Sample 2 0.40 0.26 0.75 0.014 0.022 0.22 0.04
Sample 3 0.47 0.28 0.75 0.014 0.017 0.24 0.04
Sample 4 0.47 0.28 0.76 0.015 0.03 0.24 0.04
Sample 5 0.47 0.27 0.75 0.014 0.027 0.24 0.04
Sample 6 0.48 0.27 0.75 0.014 0.021 0.25 0.04
Chemical element Ni Al Cu Sn Ca N O Fe
Sample 1 0.10 0.257 0.21 0.012 0.0004 0.007 0.0068 98.92
Sample 2 0.13 0.005 0.2 0.012 0.0004 0.008 0.0056 97.93
Sample 3 0.16 0.005 0.2 0.012 0.0005 0.008 0.0042 97.80
Sample 4 0.16 0.005 0.2 0.013 0.0003 0.007 0.0056 97.77
Sample 5 0.16 0.031 0.2 0.013 0.0003 0.008 0.0036 97.77
Sample 6 0.16 0.025 0.2 0.013 0.0001 0.008 0.0034 97.77
Table 2: Chemical composition of the slag sample
Tabela 2: Kemijska sestava vzorca `lindre
Chemical element Si Al Fe Mn Ca Mg O
Content (w/%) 6.42 10.44 0.64 0.19 43.25 2.97 36.09
With Thermo-Calc it was calculated at which con-
tents of magnesium and oxygen certain inclusions at
1550 °C are formed. Here, X represents the existence of in-
clusions inside the melt. The results are given in Table 3.
We can see that with small contents of oxygen,
magnesite (MgO) inclusions are formed; with an
increase of oxygen content, first the MgO inclusion and
then also the spinel-type inclusions (MgO·Al2O3) are
formed. With the highest oxygen content the spinel-type
inclusions and alumina inclusions are formed. The
content of magnesium does not have an effect on the
sequence of inclusion formation; it does, however, have
an effect on when the inclusions are formed. At a magne-
sium content (in mass fractions, w) of 1 · 10–3 % the first
inclusions are formed at an oxygen content of 1 · 10–5 %.
At a magnesium content of 1 · 10–4 % the first inclusions
are already formed at an oxygen content of 1 · 10–6 %.
With Thermo-Calc we have also analysed the slag,
using the chemical composition presented in table 2. We
found that at the temperatures at which the steel is in the
molten state there are four different oxides in the molten
or solidified state. 3CaO · SiO2, MgO, CaO · Al2O3 and
MgO · Al2O3 can be reduced to steel or they can first
react with phases in the slag, steel, refractory or atmo-
sphere and then be reduced to steel.
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364 Materiali in tehnologije / Materials and technology 46 (2012) 4, 361–368
Figure 3: Weight fraction of phases in dependence of the temperature
for sample 6
Slika 3: Masni dele` faz v odvisnosti od temperature za vzorec 6
Table 3: Calculation of the possibility of inclusion formation at different contents of oxygen and magnesium
Tabela 3: Izra~un mo`nosti nastanka vklju~ka pri razli~nih vsebnostih kisika in magnezija
Type of
inclusion Type of inclusion
Sample 1 Content ofoxygen (%) Al(O,C) Sample 2
Content of
oxygen (%) Spinel MgO Al2O3 mulite
Content of Mg
is 0.001(%)
0.000001
Content of Mg
is 0.001(%)
0.000001
0.00001 x 0.00001 x
0.0001 x 0.0001 x
0.001 x 0.001 x x
0.0068 x 0.0056 x x
Content of Mg
is 0.0001(%)
0.000001
Content of Mg
is 0.0001(%)
0.000001 x
0.00001 x 0.00001 x
0.0001 x 0.0001 x x
0.001 x 0.001 x x
0.0068 x 0.0056 x x
Type of inclusion Type of inclusion
Sample 5 Content ofoxygen (%) Spinel MgO Al2O3 Sample 6
Content of
oxygen (%) Spinel MgO Al2O3
Content of Mg
is 0.001(%)
0.000001
Content of
Mg is
0.001(%)
0.000001
0.00001 x 0.00001 x
0.0001 x 0.0001 x
0.001 x x 0.001 x x
0.0036 x x 0.0036 x x
Content of Mg
is 0.0001(%)
0.000001 x
Content of
Mg is
0.0001(%)
0.000001 x
0.00001 x 0.00001 x
0.0001 x x 0.0001 x x
0.001 x x 0.001 x x
0.0036 x x 0.0036 x x
3.3 Differential scanning calorimetry
Differential scanning calorimetry was made for all
the investigated samples. The rate of heating and cooling
was 5 K/s. The results of the differential scanning
calorimetry are DSC heating and cooling curves with
determined characteristic points (solidus, liquidus) and
enthalpies of fusion and solidification. We have taken a
closer look at sample 6, and because it is the last sample
taken from the melt it is the most representative as to
which inclusions are still in the melt.
In Figure 4 we can see a heating DSC curve for
sample 6 and the process of melting. At 738.5 °C the
ferrite and pearlite begin to transform to austenite. At
1401.8 °C the austenite begins to transform to -ferrite
and to the melt, at 1509.7 °C -ferrite begins to melt,
according to thermodynamic predictions.
In Figure 5 we can see a cooling DSC curve for
sample 6 and the process of solidifying. At 1392 °C the
-ferrite begins to solidify and at 1383.3 °C the -ferrite
and the melt begin to transform or solidify into austenite.
At 861.8 °C the ferrite begins to segregate.
By comparing different DSC curves, we can see the
difference in the characteristic temperatures and energies
in more detail. In Figure 6 is a comparison of the heating
DSC curves in the temperature range from 670 °C to 850
°C, where the transformation of ferrite and pearlite to
austenite occurs. The enthalpies necessary for the eutec-
toid phase transition are given in Table 4. We can see a
significant difference in sample 1; this is due to the fact
that the eutectoid reaction is presented in a small manner.
After the treatment on the ladle furnace the enthalpy is
higher.
Table 4: The enthalpy necessary for eutectoid phase transition
Tabela 4: Entalpija, potrebna za eutektoidno fazno premeno
Sample 1 2 3 4 5 6
Energy (J/g) 3.18 35.29 38.61 38.16 47.38 47.94
In Figure 7 we can see a comparison of the DSC
curves in the temperature range from 530 °C to 1020 °C
or during the eutectoid phase transition. In Table 5 the
enthalpies during the eutectoid phase transition are
given. We can see a major difference between the
samples 2 and 4 and the rest of the samples. This can be
explained by the fact that with samples 2 and 4 the
eutectoid phase transition occurs at a lower temperature
than with the other samples and a certain part of the
energy released is related to the magnetic phase
transition, resulting in the determined enthalpy values.
Table 5: The enthalpy during the eutectoid phase transition
Tabela 5: Entalpija pri eutektoidni fazni premeni
Sample 1 2 3 4 5 6
Energy (J/g) 22.67 65.93 35.9 63.29 28.73 18.88
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Materiali in tehnologije / Materials and technology 46 (2012) 4, 361–368 365
Figure 7: Enlarged comparison of the cooling DSC curves for all 6
samples
Slika 7: Primerjava pove~anih DSC krivulj ohlajanja, za vseh 6
vzorcev
Figure 4: Heating DSC curve for sample 6
Slika 4: DSC krivulja pri segrevanju vzorca 6
Figure 5: Cooling DSC curve for sample 6
Slika 5: DSC krivulja pri ohlajanju vzorca 6
Figure 6: Enlarged comparison of the heating DSC curves for all 6
samples
Slika 6: Primerjava pove~anih DSC krivulj segrevanja vseh 6 vzorcev
In general we can see that the heating DSC curves
relate well to the microstructure of our samples. The
calculated enthalpies necessary for the eutectoid phase
transition and the calculated fraction of microstructural
phases have a high correlation; this is a result of the fast
cooling of the industrial samples. The cooling DSC
curves, on the other hand, have a much better correlation
with the phase diagrams calculated with Thermo-Calc
because of a steady cooling rate.
3.4 Microstructure analysis
The microstructure of our samples can be seen in
Figure 8. It has, with the exception of sample 1, a small
fraction of ferrite (Fe) and a large fraction of pearlite
(Fe + Fe3C). Sample 1 has a small fraction of pearlite
(Fe + Fe3C) and a large fraction of ferrite (Fe). The
precise fractions of the microstructural phases are shown
in Table 6. These fractions were calculated at 200-times
magnification.
Table 6: Fraction of microstructural phases for all samples.
Tabela 6: Dele` faz v mikrostrukturi vseh vzorcev
Sample 1 2 3 4 5 6
Ferrite (%) 77.1 7.1 3.2 4.7 1.9 2.2
Pearlite (%) 22.9 92.9 96.8 95.3 98.1 97.8
In Figure 8 the microstructures of sample 5 and 6
can be seen. There is a large fraction of pearlite (Fe +
Fe3C) and small fraction of ferrite (Fe) on the austenite
phase boundaries.
With the scanning electron microscope we found
different types of inclusions in our samples: Al2O3, MnS
and spinel-type (MgO · Al2O3) modified with calcium.
We made a point and mapping analysis. In Figure 9 we
can see an alumina inclusion in sample 6. The results of
the EDS analysis of the alumina inclusion are presented
in Figure 10.
In the sample 6 microstructure there are also MnS
inclusions. Similar inclusions where found by Lamut et
al.10 in a sample taken from the tundish on a continuous
casting machine. The EDS analysis of this inclusion is
shown in Figure 11. In our sample taken from tundish
some magnesite and alumina inclusions were also found,
but they were small and usually not larger than 1 μm.
In Figure 12 a spinel-type (MgO · Al2O3) inclusion
modified with calcium is shown. It was found in sample
5 and its chemical composition is determined with a
mapping analysis, also shown in Figure 12.
The inclusion found in sample 5 is a product of the
de-oxidation process with aluminium, during which
L. KRAJNC et al.: THERMODYNAMIC ANALYSIS OF THE FORMATION OF NON-METALLIC INCLUSIONS ...
366 Materiali in tehnologije / Materials and technology 46 (2012) 4, 361–368
Figure 10: The result of the point analysis of an alumina inclusion
Slika 10: Rezultati to~kaste analize vklju~ka Al2O3
Figure 8: Microstructure: ferrite and pearlite; a) sample 5, b) sample 6
Slika 8: Mikrostruktura: ferit in perlit; a) vzorec 5, b) vzorec 6
Figure 9: Alumina inclusion (Al2O3) and the place of the point
analysis determined in sample 6
Slika 9: Vklju~ki Al2O3 in mesto to~kaste analize na vzorcu 6
alumina inclusions were formed. We assume that
because of the high content of aluminium in the melt, it
has reduced the magnesium from the refractory and the
slag. The magnesium then reacted with the oxygen and
magnesite inclusions were formed. These have then,
because of the high content, reacted with alumina and
spinel-type inclusions were formed. With the addition of
calcium for modifying the inclusions, there was a
reaction between the spinel-type inclusions and pure
alumina with calcium. The result of this was a large
globural CaO-MgO-Al2O3 system inclusion.
In agreement with the literature7 alumina inclusions
are created during the de-oxidation process with
aluminium. Furthermore, in accordance with Jiang et
al.9, the aluminium content in the melt is increased and it
reduces the magnesium. Magnesium in turn, according
to Park et al.8, because of the high activity, reacts with
the remaining oxygen, and forms magnesite inclusions,
which then react with alumina and spinel-type inclusions
are formed7. Because there are more alumina inclusions,
some of them are trapped in the matrix of the spinel-type
inclusion and these are, with the addition of calcium in
accordance with Pires et al.6, modified into large
globular inclusions5. The EDS analysis verifies this
theory, as we see in Figure 11 that there is an angular
phase formed from MgO and Al2O3 in the middle of the
inclusion, while there is a CaO and Al2O3 phase
surrounding the first one and making the inclusion
globular.
4 CONCLUSIONS
Based on an analysis of the C45 steel sample, the
following conclusions can be drawn:
The thermodynamic calculation has shown that for
low contents of oxygen in the melt, initially magnesite
inclusions are formed, and with an increase in the
oxygen content magnesite and spinel-type inclusion are
formed, whereas at high oxygen contents spinel-type and
alumina inclusions are formed.
The thermodynamic calculation of the phase
equilibrium revealed that the formation of 3·CaO · SiO2,
MgO, CaO · Al2O3 and MgO · Al2O3 oxides in the slag is
possible. These can react with phases in the slag, steel,
refractory or atmosphere and can be reduced to steel.
With the calculation of the thermodynamic equili-
brium we found that the magnesite, alumina and
spinel-type inclusion are formed during the ladle furnace
treatment. The EDS analysis confirmed that at the end of
the ladle furnace treatment there are spinel-type
inclusions in the melt and that they are modified with
calcium. There is a possibility that the inclusions are
removed during the manipulation from the ladle furnace
to the continuous casting machine, because no large
oxide inclusion was found in our tundish sample. Small
alumina inclusions (1 μm) were found as were small
MnS inclusions.
L. KRAJNC et al.: THERMODYNAMIC ANALYSIS OF THE FORMATION OF NON-METALLIC INCLUSIONS ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 361–368 367
Figure 11: Microstructure and mapping analysis of MnS10
Slika 11: Mikrostruktura in razporeditev elementov v MnS10
Figure 12: Mapping analysis of the spinel-type inclusion modified
with calcium
Slika 12: Razporeditev elementov v vklju~ku {pinela, modificiranega
s kalcijem
5 REFERENCES
1 Y. Payandeh, M. Soltanieh, Oxide Inclusions at Different Steps of
Steel Production, Journal of Iron and Steel Research, International,
14 (2007) 5, 39–46
2 H. V. Atkinson, G. Shi, Characterization of Inclusions in Clean Steel:
a Review Including the Statistics of Extreme Methods, Progress in
Materials Science, 48 (2003), 457–520
3 V. Gontarev, Teorija metalur{kih procesov. Univerza v Ljubljani,
Naravoslovnotehni{ka fakulteta, Oddelek za materiale in metalurgijo.
Ljubljana, 2005, 129
4 A. Ghosh, Secondary Steelmaking: Principles and Aplication.
Library of Congress Cataloging-in-Publication Data, 2000, 672
5 F. Tehovnik, B. Korou{i}, V. Pre{eren, Optimizacija modifikacije
nekovinskih vklju~kov v jeklih obdelanih s Ca, Kovine Zlitine
Tehnol., 26 (1992) 1/2, 125–130
6 J. C. S. Pires, A. Garcia, Modification of Oxide Inclusions Present in
Aluminum-Killed Low Carbon Steel by Addition of Calcium. REM:
R. Esc. Minas, Ouro Preto, 57 (2004) 3, 183–189
7 D. Steiner Petrovi~, B. Arh, F. Tehovnik, M. Pirnat, Magnesium
Non-metallic Inclusions in Non-Oriented Electrical Steel Sheets, ISIJ
International, 51 (2010) 12, 2069–2075
8 J. H. Park, Formation Mechanism of Spinel-Type Inclusions in
High-Alloyed Stainless Steel Melts, The Minerals, Metals &
Materials Society and ASM International, 38B (2007), 657–663
9 M. Jiang, X. Wang, B. Chen, W. Wang, Laboratory Study on Evolu-
tion Mechanisms of Non-metallic Inclusions in High Strength
Alloyed Steel Refined by High Basicity Slag, ISIJ International, 50
(2010) 1, 95–104
10 J. Lamut, M. Knap, H. Plo{tajner, B. Sen~i~, Proces modifikacije
vklju~kov s CaSi, 13. Seminar o procesni metalurgiji jekla, 2007,
134–137
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368 Materiali in tehnologije / Materials and technology 46 (2012) 4, 361–368
S. RUGMAI et al.: SMALL-ANGLE X-RAY SCATTERING SPECTRA OF IRON-BASED MAGNETIC FLUIDS
SMALL-ANGLE X-RAY SCATTERING SPECTRA OF
IRON-BASED MAGNETIC FLUIDS
MALOKOTNO SIPANJE RENTGENSKEGA SPEKTRA
MAGNETNIH TEKO^IN NA OSNOVI @ELEZA
Supagorn Rugmai1, Chitnarong Sirisathitkul2, Komkrich Chokprasombat2,
Prawet Rangsanga2, Phimphaka Harding2, Toemsak Srikhirin3, Pongsakorn Jantaratana4
1School of Physics, Institute of Science, Suranaree University of Technology and Synchrotron Light Research Institute (Public Organization),
Nakhon Ratchasima, Thailand
2Molecular Technology Research Unit, School of Science, Walailak University, Nakhon Si Thammarat, Thailand
3Department of Physics, Faculty of Science, Mahidol University, Bangkok, Thailand
4Department of Physics, Faculty of Science, Kasetsart University, Bangkok, Thailand
schitnar@wu.ac.th
Prejem rokopisa – received: 2011-11-29; sprejem za objavo – accepted for publication: 2012-03-05
Two types of superparamagnetic nanoparticles, namely magnetite (Fe3O4) and iron-platinum (FePt), with their surface modified
by oleic acid and oleyleamine were synthesized using the the polyol process and dispersed in n-hexane. Vibrating-sample
magnetometry (VSM) was used to characterize the FePt and the magnetic susceptibility of the Fe3O4 was measured as a function
of magnetic field and frequency at room temperature. Synchrotron, small-angle, X-ray scattering (SAXS) patterns revealed the
polydispersity of the magnetic nanoparticles with average radii below 5 nm. As confirmed by transmission electron microscopy
(TEM) images and dynamic light scattering (DLS), the Fe3O4 nanoparticles had a larger average size. The DLS measurement
gave larger radii than those obtained from the SAXS because of the effect of aggregates and surfactants.
Keywords: magnetic fluid, magnetite, iron-platinum, synchrotron radiation, SAXS
Dve vrsti superparamagnetnih nanodelcev, magnetit (Fe3O4) in `elezo-platina (FePt), s povr{ino, modificirano z oleinsko kislino
in oleinaminom, sta bili sintetizirani s polyolnim postopkom in dispergirani v n-heksanu. Vibracijska magnetometrija vzorcev
(VSM) je bila uporabljena za karakterizacijo FePt in magnetna ob~utljivost Fe3O4 je bila izmerjena pri sobni temperaturi kot
funkcija magnetnega polja in frekvence. Sinhrotronsko malokotno sipanje rentgenskih `arkov (SAXS) je odkrilo polidisperznost
magnetnih nanodelcev s povpre~nim premerom, manj{im od 5 nm. Kot so potrdili posnetki s transmisijsko elektronsko
mikroskopijo (TEM) in dinami~nim sipanjem svetlobe (DLS), imajo nanodelci Fe3O4 ve~jo povpre~no velikost. Meritve DLS so
pokazale ve~je premere delcev kot meritve SAXS zaradi u~inka zdru`evanja delcev in tenzidov.
Klju~ne besede: magnetna teko~ina, magnetit, `elezo-platina, sinhrotronsko sevanje, SAXS
1 INTRODUCTION
Magnetic fluids are a special class of materials that
possess the advantages of a liquid state of the carrier and
a magnetic state of the nanoparticles.1 Either ferromagne-
tic or superparamagnetic nanoparticles can be stabilized
without agglomeration and sedimentation in a liquid by
surfactant molecules, such as oleic acid and oley-
leamine.2 In addition to an increase in the magnetization
in response to an applied magnetic field, the single-
domain, superparamagnetic particles are able to generate
heat in an alternating magnetic field. Moreover, some
product properties, such as magnetoviscosity, are exclu-
sively observed in a magnetic fluid.1 As a result,
magnetic fluids are increasingly implemented in engi-
neering and biomedical applications. Magnetite (Fe3O4)
nanoparticles in either water or oil are commercially
available. In addition to the conventional uses, such
magnetic fluids are the subjects of research and develop-
ment as drug-delivery, hyperthermia and MRI-contrast
agents.3,4 On the other hand, iron-platinum (FePt) is
investigated as a material for ultra-high-density record-
ing.5 The as-synthesized FePt nanoparticles in the
superparamagnetic state are commonly stored in the
form of a magnetic fluid.
The magnetic properties of Fe3O4 and FePt nano-
particles can be characterized in the forms of colloidal
particles or dried particles on solid substrate using either
a susceptometer or a magnetometer. Since their magnetic
properties are highly dependent on the shape and size
distribution, several techniques have been employed to
characterize such nanoparticles. Transmission electron
microscopy (TEM) provides images, but the qualitative
analysis is based on sampling. Dynamic light scattering
(DLS) gives an overall size distribution, but the size of
the aggregate is usually measured instead of the
individual particles.6 Small-angle, X-ray scattering
(SAXS) using a synchrotron source offers an alternative
characterization of the shape and the size distribution. In
addition to common uses in the case of organic
nanostructures, the technique can also characterize
magnetic nanoparticles, including ferrites7,8, maghemite9,
Fe3O4 6,10,11 and FePt.12–14
In this work, Fe3O4 and FePt nanoparticles synthe-
sized using the polyol process and dispersed in hexane
were characterized by synchrotron SAXS. The results
Materiali in tehnologije / Materials and technology 46 (2012) 4, 369–373 369
UDK 621.318.1:544.171.6 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)369(2012)
were used to complement those from DLS, TEM and
magnetic measurements.
2 EXPERIMENTAL
The polyol processes adapted from Sun et al.15 were
carried out under standard airless Schlenk line tech-
niques in a nitrogen (N2) atmosphere. For the Fe3O4
nanoparticles, a mixture of 2.0 mmol of Fe(acac)3
(99.99 %), 10 mmol of 1,2 hexadecandiol (90 %),
6 mmol of oleic acid (90 %) and 6 mmol of oleyleamine
(70 %) was added to a 100-mL Schlenk flask containing
20 mL of benzyl ether (99 %). The solution was heated
and dwelled at 200 °C for 2 h and then allowed to reflux
at 300 OC for 30 min. In the synthesis of FePt, 1.0 mmol
of Fe(acac)3 and 0.5 mmol of Pt(acac)2 (97 %) were
poured into a flask, and then 20 mL of ethylene glycol
was added. The mixture was moderately stirred, purged
with N2 for 15 min at room temperature before
increasing the temperature to 120 °C. After the addition
of 5 mmol of oleic acid and 5 mmol of oleyleamine, the
solution was heated to 205 °C and then allowed to reflux
for 2 h.
After the heating source was removed, the solution
was allowed to cool down to room temperature under N2.
The product was precipitated by adding ethanol and
separated by centrifugation. The precipitate was
dispersed in n-hexane. Extra ethanol was added in this
dispersion and the dispersion was centrifuged again for
purification. After washing the particles in ethanol for
three or more times, they were dispersed in n-hexane
with the presence of a small amount (
0.05 mL) of oleic
acid and oleyleamine, followed by purging with N2 to
remove the oxygen and then stored in glass bottles at
0–4 °C.
The morphology of the nanoparticles was inspected
by TEM. The samples were prepared by depositing a few
drops of the final hexane dispersion on a carbon-coated
grid, acting as a substrate. Magnetization as a function of
the applied magnetic field for the FePt nanoparticles was
measured by vibrating-sample magnetometry (VSM). To
characterize the magnetic properties of the Fe3O4
nanoparticles, the mass susceptibility of the magnetic
fluid in a capsule was measured as function of the static
(dc) magnetic field, as well as alternating (ac) magnetic
field and its frequency using a susceptometer at room
temperature.
The SAXS measurements were performed at BL2.2
of the Synchrotron Light Research Institute, Thailand.
The X-ray wavelength was set to 0.155 nm. A CCD
detector (Mar SX165) was used to record the 2D
scattering patterns. The sample-detector distance of
1587.46 mm, set for the experiments, was calibrated by
measuring the scattering of silver behenate powder. For
the measurements, each magnetic fluid was injected into
a sample cell with aluminum foil windows. The back-
ground subtraction was carried out by subtracting the
sample scattering patterns from that of n-hexane. The
absorptions of X-rays by the sample and the n-hexane
were taking into account in the background-subtraction
process by measuring the transmitted beam intensity,
with reference to that of an empty cell, using a photo-
diode located in front of the beamstop. The scattering
profiles for the FePt and Fe3O4 nanoparticles were then
obtained by circularly averaging the background-sub-
tracted patterns.
3 RESULTS AND DISCUSSION
The variation of the complex magnetic susceptibility
of the Fe3O4 fluid with the frequency follows the trend
previously shown by Fannin et al.16–18 The response of
the superparamagnetic fluid to an alternating magnetic
field can be modeled and experimentally fitted. In
Figure 1a, the real part is decreased gradually from 1.6
× 10–6 m3/kg at 100 Hz to 1.4 × 10–6 m3/kg at 600 Hz. It
has a linear decrease with the log of the frequency from
800 Hz to 4 000 Hz before dropping to the saturation
around 4.0 × 10–7 m3/kg at 6 000 Hz. The imaginary part
is modest and remains constant over a wider frequency
range from 100 to 1000 Hz. The peak in the literature16–18
centered around 1500–2 000 Hz is not clearly observed.
At a fixed frequency, the susceptibility does not have
S. RUGMAI et al.: SMALL-ANGLE X-RAY SCATTERING SPECTRA OF IRON-BASED MAGNETIC FLUIDS
370 Materiali in tehnologije / Materials and technology 46 (2012) 4, 369–373
Figure 1: Susceptibility of Fe3O4 in hexane as a function of: a)
frequency and b) ac magnetic field at 7 kHz
Slika 1: Ob~utljivost Fe3O4 v heksanu kot funkcija: a) frekvence in b)
izmeni~nega magnetnega polja pri 7 kHz
significant variations in response to the changes in the dc
magnetic field between –1 kA/m and 1 kA/m and the ac
magnetic field up to 56 kA/m. There is an exception in
Figure 1b: the sudden increase in the real susceptibility
from 2.75 × 10–7 to 3.5 × 10–7 m3/kg occurs by increasing
the ac field from 12 kA/m to 16 kA/m. Similar behavior
was previously observed in the case of maghemite.17, 18
According to the magnetization curve in Figure 2, the
FePt fluid also exhibits superparamagnetic behavior with
saturation from 180 kA/m.
TEM images of the as-synthesized particles are
compared in Figure 3. The morphology of the Fe3O4
nanoparticles can be described as spheroid with a rough
surface. The size distribution around 5 nm in radius and
some aggregates are observed. The FePt nanoparticles in
Figure 3b are less concentrated and generally smaller
than the Fe3O4. The DLS size distribution curves of
surfactant-coated Fe3O4 and FePt have their peaks
centered at approximately 7.96 nm and 4.28 nm in
radius, respectively.
The measured SAXS intensity profiles for FePt and
Fe3O4 are shown in Figure 4. The intensity profiles are
plotted as a function of the scattering vector q, given by
q =
⎛
⎝
⎜ ⎞
⎠
⎟4π sin


(1)
where  denotes half the scattering angle and  is the
X-ray wavelength. The measured SAXS intensity
profiles from both the FePt and Fe3O4 nanoparticles
appear to have a certain slope in the high q region and
then a transition to an exponential-like profile at small
q. Considering the characteristics of the SAXS intensity
profile of homogeneous spheres, the appearance of the
measured intensity indicates that the q-range of the
measurement may contain parts of the Guinier and the
fractal regimes, as well as the transition between the two
regimes.19
In order to analyze the particle size and size
distribution, the unified exponential-power law model of
Beaucage20 was used to fit the measured data, where the
scattering intensity is given by:
S. RUGMAI et al.: SMALL-ANGLE X-RAY SCATTERING SPECTRA OF IRON-BASED MAGNETIC FLUIDS
Materiali in tehnologije / Materials and technology 46 (2012) 4, 369–373 371
Figure 3: TEM images of: a) Fe3O4 and b) FePt nanoparticles on sub-
strates
Slika 3: TEM-posnetek: a) Fe3O4- in b) FePt-nanodelcev na podlagi
Figure 4: Measured SAXS profiles from FePt- (squares) and Fe3O4-
(triangles) nanoparticles. The black and grey solid lines are the unified
exponential power-law fits for FePt and Fe3O4, respectively, with the
contributions from the Guinier and power-law terms. The q–3 line is
shown to indicate –3 slope obtained from the fit.
Slika 4: Izmerjen SAXS-profil iz FePt- (kvadratki) in Fe3O4- (tri-
kotniki) nanodelcev. ^rna in siva linija sta prilagojeni eksponentnemu
zakonu za FePt in Fe3O4, zdru`eni s prispevkom zakonitosti Guinierja
in izrazi zakona mo~i. Prikazana je linija q–3, ki ka`e naklon –3,
dobljen iz prileganja.
Figure 2: Magnetization of FePt as a function of applied magnetic
field measured by VSM
Slika 2: Magnetizacija FePt kot funkcija uporabljenega magnetnega
polja, izmerjena z VSM
I q G
q R
B
qR
q
g g
( ) exp /= −
⎛
⎝
⎜⎜
⎞
⎠
⎟⎟ +
⎛
⎝
⎜⎜
⎞
⎠
⎟⎟
⎡
⎣
⎢
⎤
⎦
⎥
2 2 3
3 6
erf
⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟
P
(2)
The prefactors G and B, the radius of gyration Rg and
the exponent P are fit parameters. This model, including
its more general form for multi-level structure,20 has
been successfully applied to describe systems with
multiple structures.21–23 The fitting was carried out with a
least-squares minimization procedure using the software
package SASfit.24 The fitted scattering-intensity curves
are shown in Figure 4, where contributions from the
Guinier and the power law terms in Eq. (2) are also
shown. The fitted exponent P = 3 was obtained for both
samples. This value is at the limit of the mass and
surface fractal and may indicate contributions from a
uniformly dense aggregate with a very rough surface.22
To investigate the particle size and size distribution, the
lognormal distribution of the particle size is assumed,
where the lognormal distribution function is given by:25
[ ]
f R
R
R m
( ) exp
ln( / )
=
⎧
⎨
⎩
⎫
⎬
⎭
1
2
2
 
π
(3)
with m denoting the median and  is the shape para-
meter of the distribution. The mode, the value of R at
the maximum, for the lognormal distribution is
m exp(–2). The mean particle radius R can be found
from the first moment of the distribution and is related
to the two parameters of the distribution by:
R m=
⎛
⎝
⎜
⎞
⎠
⎟exp


2
(4)
where R is the mean particle radius. The standard
deviation (SD) of the distribution is also related to the
two parameters, and is given by:
SD m= − ( )1 (5)
where  exp ). Following Beaucage et al.,26 the
two parameters m and  of the lognormal distribution
can be explicitly given in terms of the ratio B/G and the
radius of gyration Rg of the SAXS intensity profile. This
enables one to extract the particle size and the distri-
bution information from the fit of the unified model of
Beaucage, Eq.(2), to the measured SAXS intensity
profile. The parameters m and  are given by:26
m
R
=
5
3 14
2
g
exp( ) 
(6)
 =
ln( / . )BR Gg
4 162
12
(7)
The value ( / . )BR Gg
4 162 in Eq.(7) also serves as an
index of polydispersity, in which the value is 1 for a
monodisperse sphere.20 The lognormal distribution
functions for both types of nanoparticles are shown in
Figure 5. The relevant fitted parameters and calculated
mean particle radii and standard deviations of the size
distribution are shown in Table 1. The results give a
smaller particle size for FePt than for Fe3O4, which
agrees with the observation from the TEM images. The
high standard deviation and index of polydispersity
relative to the size of FePt can be described by the
heterocoagulation model of FePt formation in the
modified polyol process.27 The obtained sizes of both
types of particles are smaller than those obtained by
DLS, which is a trend previously observed in the
literature.6 The reason is not only due to the exclusion of
surfactants by the SAXS but also the effect of
aggregation in the case of DLS.
Table 1: Fitted parameters for the SAXS intensity profile and
calculated mean particle radius and standard deviation compared with
the DLS results
Tabela 1: Prilagojeni parametri SAXS intenzitetnega profila in
izra~unan premer delca ter standardni odklon v primerjavi z
DLS-rezultati
Sample
SAXS DLS
BRg4/
(1.62G) Rg/nm
R /nm SD Rpeak/nm
FePt 6.55 3.67 1.71 0.70 4.28
Fe3O4 1.92 4.21 3.82 0.90 7.96
4 CONCLUSION
Synchrotron SAXS spectra revealed the polydisper-
sity of superparamagnetic Fe3O4 and FePt nanoparticles
in n-hexane synthesized using the polyol process. FePt
nanoparticles had a smaller average size, which is in a
good agreement with DLS and TEM images. The radii
obtained from the DLS measurement tended to be larger
than those from SAXS because the aggregates and
surfactants were also taken into account.
S. RUGMAI et al.: SMALL-ANGLE X-RAY SCATTERING SPECTRA OF IRON-BASED MAGNETIC FLUIDS
372 Materiali in tehnologije / Materials and technology 46 (2012) 4, 369–373
Figure 5: Lognormal size distributions for FePt (black solid line) and
Fe3O4 (grey solid line) obtained from the unified exponential
power-law fit. The positions of the mode, median and mean radii are
also shown for each distribution by dotted, dashed and dot-dashed
lines, respectively.
Slika 5: Normalna logaritemska razporeditev velikosti FePt (~rna
linija) in Fe3O4 (siva linija), dobljena iz poenotene eksponentne
odvisnosti. Prikazane so tudi pozicije vrste srednje in najve~je veli-
kosti delcev za vsako razvrstitev s pik~asto, ~rtkano in pik~asto-~rt-
kano linijo.
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S. RUGMAI et al.: SMALL-ANGLE X-RAY SCATTERING SPECTRA OF IRON-BASED MAGNETIC FLUIDS
Materiali in tehnologije / Materials and technology 46 (2012) 4, 369–373 373

I. N. ORBULOV, K. MÁJLINGER: MICROSTRUCTURE OF METAL-MATRIX COMPOSITES REINFORCED BY ...
MICROSTRUCTURE OF METAL-MATRIX COMPOSITES
REINFORCED BY CERAMIC MICROBALLOONS
MIKROSTRUKTURA KOMPOZITOV S KOVINSKO OSNOVO,
OJA^ANO S KERAMI^NIMI MIKROKROGLICAMI
Imre Norbert Orbulov, Kornél Májlinger
Department of Materials Science and Engineering, Budapest University of Technology and Economics, Bertalan Lajos utca 7.,
1111, Budapest, Hungary
orbulov@eik.bme.hu, orbulov@gmail.com
Prejem rokopisa – received: 2011-12-03; sprejem za objavo – accepted for publication: 2012-02-06
Metal-matrix composites reinforced by ceramic hollow microspheres were produced as special porous metals, called
metal-matrix syntactic foams (MMSFs). In this paper the microstructure of the ceramic hollow microspheres as reinforcing
elements is investigated in connection with the production of MMSFs with the pressure infiltration. SL150- and SL300-type
ceramic microspheres from Envirospheres Ltd (Australia) were investigated. They contained various oxide ceramics, mainly
Al2O3 and SiO2. The chemical composition and the microstructure of the microspheres had a strong effect on their infiltration
characteristics; therefore, in the view of the MMSF production it was very important to know microstructural details about the
microspheres. Due to this energy-dispersive X-ray spectroscopy, maps were recorded from the cross-sections of the
microspheres’ walls. The results showed that the Al2O3 and SiO2 distribution were not equal; the Al2O3 phase was embedded in
the surrounding mullite, while the SiO2 phase occurred in the form of needles. Line energy-dispersive X-ray spectroscopy
measurements were performed in order to investigate the possible reaction between different aluminium alloy matrices and the
ceramic microspheres. The results showed that, due to an uneven distribution of Al2O3, rich particles of the molten aluminium
could reduce the SiO2 rich parts of the microspheres and the walls of the hollow microspheres became damaged and degraded.
This chemical reaction between the microspheres and the walls could make the infiltration easier, but the resulting mechanical
properties will be reduced due to the damaged microsphere walls.
Keywords: microsphere, microballoon, energy dispersive X-ray spectroscopy, metal matrix syntactic foam, metal matrix
composite
Kompoziti s kovinsko osnovo, oja~ano s kerami~nimi votlimi mikrokroglicami, so bili izdelani kot posebna porozna kovina z
nazivom sintakti~na pena s kovinsko osnovo (MMSF). V tem ~lanku je opisana raziskava mikrostrukture kerami~nih votlih
mikrokroglic v povezavi z izdelavo MMSF z infiltracijo pod tlakom. Preiskovane so bile SL150- in SL300-kerami~ne
mikrokroglice proizvajalca Envirospheres Ltd (Avstralija). Vsebujejo razli~ne kerami~ne okside, ve~inoma Al2O3 in SiO2.
Kemijska sestava in mikrostruktura mikrokroglic imata velik vpliv na sposobnost njihove infiltracije, zato je z vidika izdelave
MMSF pomembno poznati podrobnosti mikrostrukture teh mikrokroglic. Zaradi tega je bila na prerezu stene mikrostrukturnih
kroglic posneta razporeditev elementov z energijsko disperzijsko spektroskopijo. Rezultati ka`ejo, da razporeditev Al2O3 in
SiO2 ni enaka: Al2O3-faza je vgrajena v obkro`ajo~i mulit, medtem ko se SiO2-faza pojavlja v obliki igel. Da bi preiskali
morebitne reakcije med osnovo iz razli~nih aluminijevih zlitin in kerami~nimi mikrokroglicami, je bila izvr{ena linijska
energijsko disperzijska rentgenska spektroskopija. Rezultati so pokazali, da zaradi neenakomerne razporeditve z Al2O3 bogati
delci v staljenem aluminiju lahko reducirajo s SiO2 bogata podro~ja mikrokroglic, in stene votlih mikrokroglic postanejo
po{kodovane in degradirane. Kemijske reakcije med mikrokroglicami in stenami lahko olaj{ajo infiltracijo, vendar pa se
mehanske lastnosti poslab{ajo zaradi po{kodovanih sten mikrokroglic.
Klju~ne besede: mikrokroglice, mikrobaloni, energijsko disperzijska rentgenska spektroskopija, pena, skladna s kovinsko
osnovo, kompozit s kovinsko osnovo
1 INTRODUCTION
Nowadays metallic foams are more and more impor-
tant and this is confirmed by the increasing number of
papers published on this topic. The ‘conventional’ metal-
lic foams, which contain only metallic and gas phases,
are widely covered by the literature. However, there are
still unsolved problems, for example the foaming process
of the foams1,2. The metallic foams belong to a special
class that also satisfies the definition of the particle-rein-
forced metal-matrix composites. These are the metal-ma-
trix syntactic foams (MMSFs). They were first produced
in the 1990s. The MMSFs have numerous promising ap-
plications, such as covers, hulls, castings, or they can be
used in the automotive and electromechanical industry
sectors because of their high-energy absorbing and
damping capabilities. In these porous materials the po-
rosity is ensured by incorporating ceramic hollow
microspheres3. The microspheres are commercially
available and they mainly contain various oxide
ceramics4,5. The quality of the microspheres has a strong
effect on the mechanical and other properties of the
foams. The behavior of the foams has been widely stud-
ied.
The most important properties of the foams are the
compressive strength and the absorbed energy. Wu et al.
examined the effects of the microballoon size on the
compressive strength. They found that smaller micro-
spheres ensure higher compressive strength because they
contain fewer flaws in their microstructure than the
larger ones. The damage propagation of the foams was
also investigated. The fracture was initialized in the cor-
Materiali in tehnologije / Materials and technology 46 (2012) 4, 375–382 375
UDK 66.017:666.3/.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)375(2012)
ners of the specimens by the shearing of the micro-
spheres6. Rohatgi et al. also investigated the size effect of
the microspheres, but not only in view of the compres-
sive strength, but in view of infiltration, too. Their mea-
surements showed that larger microspheres can be infil-
trated easier7. Palmer et al. proved that larger
microspheres contain more porosity in their walls and
more flaws in their microstructure than the smaller ones8.
The results of the performed upsetting tests were com-
pared to the other research on this topic. The conclusions
were the same9,10. Balch et al. performed a special upset-
ting test. The loading was applied in small steps and af-
ter each test, X-ray or neutron-diffraction measurements
were carried out. The main aim of the work was to inves-
tigate the load transfer from the matrix to the
microspheres. They found a chemical reaction between
the microspheres and the matrix materials, which has a
detrimental effect on the load transfer and, through that,
on the mechanical properties of the foams11. In their pre-
vious work Balch et al. found that the microspheres have
at least the same importance in the syntactic foams as the
matrix material. The fracture strength and the yield
strength of the matrix determine the failure stress of the
syntactic foams. Therefore, the investigation of the
microstructure and the quality of the microspheres is
very important9. In addition to the compressive strength,
other mechanical properties, such as the tensile strength,
or the hardness of the syntactic foams were investi-
gated12,13. The sliding behavior of the syntactic foams
was also examined because the ceramic microspheres
have high hardness and therefore the composites show a
better wear behaviour than the pure matrix14,15.
As it can be seen in the previous paragraph, the qual-
ity and chemical composition of the microspheres influ-
ence many properties of the syntactic foams. And they
also have a strong influence on the production of the syn-
tactic foams. The foams are usually produced by the
mixing technique and gravitational casting or by the
pressure infiltration. In all these cases the contact angle
between the ceramic microspheres and the metal matrix
has a detrimental effect on the infiltration characteristics
and on the threshold pressure (in the case of pressure in-
filtration)16–18. The contact angle is influenced by many
parameters like the chemical composition and the possi-
ble reaction between the reinforcement and the matrix
material. Therefore, the microspheres should be pre-
cisely investigated on the microstructure’s scale.
The ultimate method for this purpose is the scanning
electron microscopy and the energy-dispersive X-ray
spectroscopy (EDS). EDS is sensitive to the chemical
composition and able to investigate a point, a line or
even an area. These possibilities give extremely good op-
portunities for acquiring detailed information about the
microstructure of the microspheres and about the distri-
bution of their constituents. According to the published
works mentioned above, the main aim of this work was
to investigate the microstructure and the distribution of
the constituents in the ceramic hollow microspheres and
to investigate the interface between different alu-
minium-alloy matrices and the ceramic microspheres in
order to provide information about the metal-matrix syn-
tactic foam production and their expectable mechanical
behaviour.
2 MATERIALS AND METHODS
The investigated materials were the SL150- and
SL300-type microspheres provided by Envirospheres Ltd
(Australia)4. Their main parameters are listed in Table 1.
The phase composition was determined by X-ray diffrac-
tion measurements. For this purpose a Phillips X-Pert-
type diffractometer with 35-mA cathode heating current
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Table 1: Morphological properties and the phase constitution of the applied hollow ceramic spheres
Tabela 1: Morfolo{ke zna~ilnosti fazne sestave uporabljenih votlih kerami~nih kroglic
Type
Average
diameter Size range
Specific
surface Al2O3
Amorphous
SiO2
Mullite Quartz Other
(μm) (μm) (μm–1) w/%
SL150 100 56–183 0.060
30–35 45–50 19 1 bal.
SL300 150 101–330 0.040
Table 2: Calculated and measured density and porosity values of the MMSFs3
Tabela 2: Izra~unane in izmerjene vrednosti za gostoto in poroznost MMSF3
Specimen
Density (g cm–3) Porosity (%)
Theoretical Measured Particle Matrix Total
Al99.5-SL150 1.34 1.43 50.9 –6.2 44.7
Al99.5-SL300 1.42 1.52 48.2 –7.2 41.0
AlSi12-SL150 1.32 1.31 50.9 1.1 52.0
AlSi12-SL300 1.40 1.37 48.2 1.9 50.1
AlMgSi1-SL150 1.34 1.52 50.9 –13.4 37.5
AlMgSi1-SL300 1.42 1.57 48.2 –10.5 37.7
AlCu5-SL150 1.37 1.53 50.9 –11.6 39.3
AlCu5-SL300 1.44 1.62 48.2 –12.2 36.0
and a copper anode (Cu K,  = 0.154186 nm) with a
40-kV voltage were used. The rotating speed of the
goniometer was 0.04° s–1. The Scanning Electron Micro-
scope (SEM) tests were performed by a Phillips
XL-30-type electron microscope equipped with an
EDAX Genesis EDS analyzer. The ceramic microspheres
were coated with carbon in order to put a conductive
layer on them, but basically they were investigated in
as-received condition. The excitation was 15 kV and
EDS maps were recorded from the surface of the micro-
spheres. Five typical microspheres from each SL-type
microsphere group were investigated.
Later the microspheres were incorporated in pure
aluminium (Al 99.5), aluminium-silicon (AlSi12),
aluminium-magnesium-silicon (AlMgSi) or aluminium-
copper (AlCu5) alloys to create MMSFs. The foams
were designated according to their matrices and
reinforcement. For example, Al99.5-SL150 denotes a
pure aluminium-matrix syntactic foam with an
SL15-microsphere reinforcement. The volume fraction
of the microspheres was maintained at a relatively high
volume fraction (
60 %) level and the production
method is described, in details, elsewhere3. The density
and the porosity values of MMSFs are listed in Table 2,
while the chemical compositions of MMSFs are shown
in Table 3. In Table 2 theoretical density and particle
porosity were calculated from the geometrical
parameters of the microspheres. The matrix porosities
were calculated as the difference between the theoretical
and the measured density divided by the theoretical
density. The negative matrix porosity refers to the
infiltrated microspheres (the particle porosity should be
decreased). However, the values of the matrix porosity
were always below 8 %, so the infiltration can be
qualified as good enough. The values of Table 3 were
determined with the XRD measurements described
above. EDS line measurements were carried out on the
foams to characterize the elemental distribution at the
interfaces, where the hollow microspheres are in contact
with the matrix. All EDS-line analyses were performed
on metallographically polished surfaces, and the
preparation steps for the EDS analysis with an automatic
grinding / polishing machine are listed in Table 4. The
excitation voltage for the EDS analysis was 20 kV. The
line EDS measurement started on the matrix material and
crossed the wall of a hollow sphere. One hundred points
were measured along each line; each point was excited
for 20 s with a 35-μs detector acquisition rate.
3 RESULTS AND DISCUSSION
3.1 Investigation of the microspheres’ walls (map EDS
measurements)
In Figure 1 a typical site from the outer surface of an
SL150-type microsphere is shown, while in Figure 2, a
SEM micrograph from the cross-section of a micro-
sphere in an Al99.5-matrix foam sample is presented. In
both images needle-like structures can be clearly ob-
served. They are densely situated and they do not have
any distinguished direction. The different gray scale of
the needles on the back-scattered electron (BSE) images
indicates a somewhat different chemical composition.
The needles are very small, their length is between 5 μm
and 10 μm, while their diameter is smaller than 0.5 μm.
EDS maps provide information that is much more useful
than the single EDS spot measurements, because single
spots can be largely effected by the surrounding matrix.
This effect can be decreased, or avoided, with the EDS
maps, which can show the distribution of the elements.
Because of this reason the EDS-map analyses were
performed on the cross-sections of the microspheres to
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Materiali in tehnologije / Materials and technology 46 (2012) 4, 375–382 377
Table 3: Phase constitution of aluminum-matrix syntactic foams according to the XRD measurements (in mass fractions, w/%)
Tabela 3: Fazna sestava pen, skladnih s kovinsko osnovo, izvr{ena z XRD-meritvami (v masnih dele`ih, w/%)
Specimen Al Si Mullite -Al2O3 -Al2O3 Amorph. CuAl2
Al99.5-SL150 67 8 11 3 11 0 -
Al99.5-SL300 78 0 11 0 0 11 -
AlSi12-SL150 72 7 13 0 0 8 -
AlSi12-SL300 72 7 12 0 0 8 -
AlMgSi1-SL150 60 7 8 0 25 0 -
AlMgSi1-SL300 60 6 6 0 28 0 -
AlCu5-SL150 60 6 8 8 12 0 6
AlCu5-SL300 60 5 10 7 12 0 6
Table 4: Sample-preparation steps for the EDS analysis with an automatic grinding / polishing machine
Tabela 4: Stopnje priprave vzorcev za EDS-analizo z avtomatskim brusilno-polirnim sistemom
Abrasive material Duration of grinding/polishing (min)
Grinding/polishing force
(N)
Rotation speed
(r/min) Rotation direction
P 320 SiC 1 22 220 counter
6 μm diamond 15 27 150 counter
3 μm diamond 6 27 150 counter
0.05 μm SiO 3 27 125 compliance
get additional information about the element distribution
in a microsphere’s wall. To illustrate this analysis, the
EDS maps of an SL300-type microsphere are presented
(Figure 3). It is important to emphasize that all of the
other microsphere types (SL150 and SL300) showed the
same features. Figure 3a shows a SEM image of the
investigated surface. The needle-like structure is again
very clear. Figure 3b shows the distribution of
aluminium. It is evident that the needles contain more Al
than the surrounding environment, but Al can be found
everywhere, not only in the needles. From the XRD
measurements it is also clear that Al2O3 and SiO2 form
mullite (3Al2O3·2SiO2), therefore – in accordance with
the EDS map and Table 1 – the wall of a microsphere is
built up from the mixture of mullite and amorphous
SiO2. This means that the distribution of Al2O3 is uneven;
it can be found as a part of mullite in the wall and as
Al2O3 needles embedded in this wall matrix. Figure 3d
definitely confirms this conclusion by showing the
distribution of silicon. Si can be found everywhere on the
surface (mainly amorphous SiO2 mixed with mullite)
except in the needles (the needles appear black in this
picture). This indicates that the needles do not contain Si
and, therefore, they are really Al2O3 needles. Finally, as
it is expected, the oxygen distribution (in Figure 3c) is
totally balanced; it is built in both Al2O3 and SiO2.
The analysis above shows the formation of the Al2O3
rich zones and it also indicates the presence of the amor-
phous SiO2 rich zones. The amorphous SiO2 is undesir-
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378 Materiali in tehnologije / Materials and technology 46 (2012) 4, 375–382
Figure 3: SEM image a) from the surface of an SL300-type
microsphere and EDS maps of this area showing element distributions
of: b) Al, c) O and d) Si
Slika 3: SEM-posnetek a) povr{ine SL300 mikrokroglice in EDS
posnetek razporeditve elementov: b) Al, c) O in d) Si
Figure 1: SEM image from the surface of an SL150-type ceramic
hollow microsphere
Slika 1: SEM-posnetek povr{ine SL150 votle kerami~ne mikro-
kroglice
Figure 4: Micrographs of: a) damaged and b) undamaged micro-
spheres in the a) Al99.5 and b) AlSi12 matrices
Slika 4: Mikroposnetek: a) po{kodovane in b) nepo{kodovanih mikro-
kroglic v a) Al99,5 in b) AlSi12 osnovi
Figure 2: SEM image of the cross-section of an SL150-type ceramic
hollow microsphere’s wall
Slika 2: SEM-posnetek pre~nega prereza stene SL150 votle kerami~ne
mikrokroglice
able, because – opposite to Al2O3 and mullite – the
chemical stability of SiO2 at an elevated temperature is
not good enough. During the production of MMSFs the
molten aluminium can reduce SiO2 according to the fol-
lowing chemical reaction:
4Al(liq) + 3SiO2 (sol) 
 2Al2O3 (sol) + 3Si(sol)
At first sight this reaction is advantageous, because it
forms Al2O3 (with better properties) from amorphous
SiO2. But this is a diffusion-controlled reaction leading
to a degradation of the microspheres’ walls as it is shown
in Figure 4a. This indicates a drastic drop in the com-
pressive strength and other mechanical properties as it
was shown in the previous papers3,19,20. As indicated by
the XRD results of Table 3, the reaction mainly pro-
duced -Al2O3 and it took place only in the case of a pure
aluminium matrix reinforced with the SL150- and
SL300-type microspheres. In the case of the SL300-type
microspheres the reaction was suppressed with a lower
infiltration temperature. The infiltration temperature had
a strong influence on the kinetics of the change reaction.
That is why there was no reaction in the case of the
Al99.5-SL300-type syntactic foams infiltrated at a lower
(690 °C) temperature3. This indicates that there is a limit
relating to the temperature. Above this temperature the
reaction is intensive, but below the temperature limit, the
reaction will not occur (for details see3). The reaction did
not take place in the case of the AlSi12-matrix material,
because the driving force of the diffusion-controlled
chemical reaction is the Si difference between the molten
matrix and the solid microspheres. In the case of the
AlSi12 metal, the large Si content decreased the driving
force; the reaction became suppressed and did not take
place. The walls of the microspheres remained unharmed
as it can be observed in Figure 4b. Most of -Al2O3 was
found in the MMSFs with the AlMgSi1 matrix and the
above mentioned reaction also took place in the samples
with the AlCu5 matrix. In the case of the MMSFs with
the AlCu5 matrix, the thermodynamic conditions also
enabled the formation of the CuAl2 phase.
To summarize, the microspheres form the Al2O3 nee-
dles embedded in the mixture of the amorphous SiO2 and
mullite. The chemically reactive molten aluminium can
damage the microspheres by reducing their amorphous
SiO2-rich parts. This is undesirable with respect to me-
chanical properties.
3.2 Investigation of the interface layer between micro-
spheres and the matrix
After a detailed examination of the microspheres’
surfaces, the investigations of MMSFs were carried out.
First, EDS overview-map measurements were
performed in the microsphere-matrix region of the
MMSFs with different matrices. For example, in the case
of an Al99.5 matrix with SL300 microspheres, the
EDS-map results are shown in Figure 5. The area of the
SEM image (Figure 5a) was investigated for the distri-
bution of the alloying elements (Figures 5b to 5d). The
microstructure of the matrix-grain boundaries formed
after the solidification can be seen clearly in the con-
centration differences between aluminium (Figure 5b)
and silicon (Figure 5d) in the EDS images covering all
the samples. Oxygen in a higher concentration was only
detected in the walls of the microspheres (Figure 5c), for
all of the samples. In the case of the AlSi12 matrix larger
areas of primer silicon were detected as expected. In the
samples with the AlCu5 matrix copper-rich precipi-
tations can be observed (with a dimension of approx.
1 μm × 20 μm in Figure 6); according to the XRD
measurements they are CuAl2 particles. In the samples
with the AlMgSi1 matrix magnesium showed a uniform
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Materiali in tehnologije / Materials and technology 46 (2012) 4, 375–382 379
Figure 6: EDS map of an AlCu5-matrix MMSF with an SL150-type
microsphere showing copper distribution and precipitations with the
CuAl2 phase
Slika 6: EDS-razporeditev bakra in izlo~ki CuAl2 faze v AlCu-osnovi
MMSF z mikrokroglico SL150
Figure 5: SEM image: a) from an Al99.5-matrix MMSF with an
SL300-type microsphere and EDS maps of this area showing element
distributions of: b) Al, c) O and d) Si
Slika 5: SEM-posnetek: a) Al99,5 MMSF osnove z SL300 mikro-
kroglico in EDS razporeditev elementov: b) Al, c) O in d) Si
distribution in the aluminium areas, and no magnesium
was detected in the primer silicon precipitations at all.
Magnesium enrichment was detected alongside the
microspheres’ outer walls (Figure 7) indicating the
solution of magnesium in the microspheres walls. For a
more detailed investigation of the microsphere
interface-matrix region, EDS line measurements were
done at a higher resolution.
EDS line measurements were performed on polished
specimens perpendicular to the interface layer between
the microspheres and the matrix material. The interface
layer is very important, because it is responsible for the
load transfer from the matrix to the microspheres. The
line-scan profiles showed the alternation of chemical
elements along the line. The biggest advantage of the
EDS line method is that it offers a very good opportunity
to examine the interface layer and the changes in a
hollow-microsphere wall in the matrix. Examples for the
AlMgSi1 and the AlCu5 matrices are shown in Figures
8 and 9, respectively.
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380 Materiali in tehnologije / Materials and technology 46 (2012) 4, 375–382
Figure 9: BSE image and EDS line-scan profiles of the AlCu5-SL300
syntactic foam
Slika 9: BSE-posnetek in profil EDS-linijske analize pene, skladne z
AlCu5-SL300
Figure 8: BSE image and EDS line-scan profiles of the
AlMgSi-SL150 syntactic foam
Slika 8: BSE-posnetek in profil EDS-linijske analize pene, skladne z
AlMgSi-SL50
Figure 7: EDS map of an AlMgSi-matrix MMSF with an SL300-type
microsphere showing Mg distribution and Mg enrichment in the
microsphere’s outer wall
Slika 7: EDS-razporeditev in obogatitev Mg v zunanji steni mikro-
kroglice v AlMgSi-osnovi MMSF z mikrokroglicami SL300
In Figure 8, the wall of an SL150-type hollow
microsphere can be observed in the high-magnification
BSE image. It can be seen that the outer edge of the wall
is not very well defined. This means that the surface of
the microsphere is degraded. As mentioned above, the
exchange reaction produces Al2O3, which is advanta-
geous, but not at the price of degrading the hollow
microsphere’s wall. The diffusion process made a rela-
tively wide interface layer on the outer surface of the
hollow microspheres (i.e. from point B to point C). The
width of the interface layer was relatively wide 
2–4 μm
(in previous research it was 6 μm 20) and it was measured
between the significant changes of the derivation of the
fitted curves showing the changing of the Al. Along the
interface the Si and Al contents increased and decreased,
respectively, with a moderate slope. After point C the Al
content was alternated according to the actual composi-
tion of the wall. From point D the measurement is not re-
liable because of the curvature of the inner surface of the
hollow microsphere.
In all of the aluminium and Al-alloy matrices the
alternation of the Al and Si contents was observed
because of the presence of the primer-Si precipitations in
the alloy. In the case of the AlMgSi1 matrix (Figure 8)
such primer-Si precipitation can be seen between points
A and B. Points B and C were close to each other. This
indicates that there is a very narrow (i.e. less than 3 μm)
observable interface layer. In the concentration-sensitive
BSE image, the lighter needle-like phases can be
observed again in the walls of the SL150 spheres. The
increase of the Mg content after point B confirms the
presence of the magnesium solution in the microsphere
wall, less than 4 μm deep (this was also indicated by the
EDS map measurements). In the case of the AlCu5
matrix (Figure 9) a precipitation (the CuAl2 phase) can
be observed between points A and B, while the interface
layer was about 3 μm on both SL150 and SL300
microspheres. Parts of the microspheres’ walls were
covered with the Cu precipitations, which can be clearly
seen on the BSE image, but according to the EDS line
measurements (between point B and C) this is not a
solution in the microspheres’ walls, but just a precipi-
tation on them. In the case of the AlSi12 matrix the outer
surface of the microsphere was seemingly unharmed as
the exchange reaction was suppressed by the large Si
content of the matrix. The interface layer was less than
1.5 μm for both SL-type microspheres.
According to this, a well-defined Al and Si alterna-
tion occurred in the line of the EDS analysis diagram. In
the lighter areas, the Si content decreased, while the Al
content increased. There was no big fluctuation in the
oxygen content within the matrix and within the
microspheres’ walls. This implies that, again, the lighter
phases in the walls of the hollow microspheres are the
Al2O3 particles embedded in the SiO2 and mullite matrix.
At point D, on the inner side of the hollow microsphere’s
wall, a very narrow Fe, Mg and K rich zone can be ob-
served (originated from the various oxides of the hollow
microspheres).
Considering the possibility of chemical reactions be-
tween the microspheres and the matrix materials, it is
worth mentioning that the ceramic microspheres can be
used as a source of alloying elements in the MMSF sys-
tems. The appropriate choice and concentration of an al-
loying element in the wall material, or in the surface
coating8 of the microspheres, can result in precipitation,
grain refinement, and microstructure as per the designed
scheme. These alloying elements can enhance the me-
chanical and/or other properties of the MMSFs. How-
ever, there is an effective range of the alloying elements
due to the limited time after the infiltration and before
the cooling. If the effective distance between the
microspheres is small enough, it is possible to guarantee
a homogenous microstructure. If the distance is larger,
then the alloying effects are only in the vicinity of the
microspheres.
To summarize, the EDS line measurements are appli-
cable when investigating the interface layer between the
microspheres and the matrix material. The investigations
proved the presence of the Al2O3 needles in the walls and
an exchange reaction between the SiO2 content of the
microsphere and the molten aluminium during the pro-
duction. The suppressing effect of a high Si content was
also confirmed. The alloying, or the coating, of the
microspheres offers a great opportunity to influence the
microstructure and the properties of MMSFs.
The authors assume that the microspheres also have
an effect on the orientation and the size of the ma-
trix-material grains. Therefore, the following paper will
deal with the electron backscatter diffraction (EBSD) in-
vestigation of MMSFs.
4 CONCLUSIONS
From the results of the above mentioned and dis-
cussed EDS and XRD measurements, the following con-
clusions can be drawn:
The microstructures of the microspheres, their vol-
ume fraction and properties have a strong effect on the
properties of MMSFs. The SEM and EDS investigations
showed that Al2O3 needles can be found in the walls of
the hollow ceramic microspheres. These needles are
packed densely and they are embedded in the mixture of
mullite and SiO2.
Due to the uneven distribution of Al2O3, SiO2 rich
zones were formed on the surfaces of the microspheres.
During the MMSF production the molten aluminium
chemically attacked these zones and this reductive chem-
ical reaction resulted in a severe damage of the micro-
spheres’ walls.
In the case of the Al99.5, AlCu5 and AlMgSi1 matri-
ces the reaction was intensive. The driving force of the
diffusion-controlled reaction was the Si concentration
gradient between the microspheres and the matrix. In the
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Materiali in tehnologije / Materials and technology 46 (2012) 4, 375–382 381
case of the AlSi12-matrix syntactic foams, the reaction
was suppressed by a considerable amount of Si in the
matrix alloy.
In the case of the AlCu5 matrix, local-copper precipi-
tations were found on the microspheres’ walls, while in
the case of the AlMgSi1matrix, magnesium was solved
into the outer regions of the microspheres’ walls.
In addition to the concrete conclusions above, it is
worth mentioning that the ceramic microspheres can be
used as a source of alloying elements in the MMSF sys-
tems. The appropriate choice and concentration of an al-
loying element in the wall material, or in a surface coat-
ing, of the microspheres can enhance the properties of
MMSFs.
Acknowledgements
Special thanks to Professor J. T. Blucher for his kind
support and to I. Sajó for XRD measurements. The
Metal Matrix Composites Laboratory is supported by
Grant # GVOP 3.2.1-2004-04-0145/3.0. This paper was
supported by the János Bolyai Research Scholarship of
the Hungarian Academy of Sciences. The investigations
were supported by the Hungarian Research Fund,
NKTH-OTKA PD 83687. This work is connected to the
scientific programme of the "Development of quality-ori-
ented and harmonized R+D+I strategy and functional
model at BME" project. This project is supported by the
New Széchenyi Plan (Project ID: TÁMOP-4.2.1/
B-09/1/KMR-2010-0002).
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THE END MILLING OF THE AA6013 ALUMINUM
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KON^NEM REZKANJU ALUMINIJEVE ZLITINE AA6013
Hülya Durmuþ
Celal Bayar University, Engineering Faculty, Department of Materials Engineering, Muradiye, Manisa-Turkey
hulya.kacar@bayar.edu.tr
Prejem rokopisa – received: 2011-12-19; sprejem za objavo – accepted for publication: 2011-12-21
In the study the effects of multi-process parameters on the surface roughness during the machining of the AA6013 aluminum
alloy with non-coated cemented carbide milling cutters were investigated by using Taguchi’s experimental design method.
Experiments were conducted on the basis of Taguchi’s L16 orthogonal array (OA) at different levels of the factors relating to the
aging process, the cutting speed, the feed rate, the axial depth of a cut and the radial depth of a cut; then they were evaluated
through the signal/noise ratio (S/N), ANOVA and the main effects’ graphic. Accordingly, the lowest surface roughness was
estimated to be 0.436 μm and, finally, Taguchi’s method allowed the optimization of the system for the verification of the tests.
Keywords: milling, surface roughness, Taguchi’s method, aging, experimental design
Raziskan je vpliv multiprocesnih parametrov na hrapavost povr{ine aluminijeve zlitine AA6013, obdelane s karbidnim
rezkarjem in z na~rtovanjem preizkusov po Taguchijevi metodi. Preizkusi so bili opravljeni na podlagi Taguchijeve ortogonalne
porazdelitve L16 po razli~nem staranju zlitine, pri razli~nih hitrostih rezkanja, podajanja, aksialne in radialne globine rezov in
ocenjeni na podlagi razmerja signal/hrup (S/N), ANOVE in grafi~nih zapisov glavnih vplivov. Na podagi rezultatov je bila
ocenjena kot najmanj{a hrapavost povr{ine 0,436 μm, s Taguchijevo metodo pa so bili verificirani preizkusi in optimiziran
sistem.
Klju~ne besede: rezkanje, hrapavost povr{ine, Taguchijeva metoda, staranje, na~rtovanje preizkusov
1 INTRODUCTION
Aluminum alloys have been extensively used in
aviation, automotive, plastic injection molding and
defense industries due to their high resistance/weight
ratios, good corrosion/fatigue resistance and high feed
rates in recent years. Therefore, studies of the machining
of these materials have become more important1. Most of
the studies have been conducted on milling these
materials, frequently focusing on the effects of cutting
parameters and tool geometries on surface roughness.
The studies conducted on the variation in surface
roughness in end milling aluminum are evaluated below.
This is a method, in which effects of only quantitative
parameters can be obtained in a small experiment
number compared with full factorial experiment design
and both the statistical model and the optimization can
be obtained. Whang and Chang2 modeled the effect of
cutting parameters and cutter geometry on the surface
roughness in milling Al2014-T6 under dry and wet
conditions with RSM. It was determined that the factors
of the cutting speed, the feed rate, the concavity and the
axial gap angle were significant for the variation in the
surface roughness under dry conditions, while the feed
rate and the concavity angle were significant under wet
conditions. Erzurumlu and Öktem3 modeled the surface
roughness of a mold made of AA7075 material with the
help of the methods of RSM and artificial neural
networks (ANN) and they found that the model based on
ANN was higher in accuracy. Routara et al.4 modeled the
machining of the materials of AA6061-T4, AISI 1040
and UNS C34000 with a CVD-coated carbide end mill
with RSM. They modeled it as the second-order regres-
sion equation for five surface roughness parameters of
the depth of a cut, the spindle speed and the feed rate,
which were reference parameters in this area, and they
accomplished the optimization of the system.
Artificial neural networks, the neural fuzzy logic,
fuzzy networks and genetic programming methods are
also preferred in this area. Lou and Chen5 investigated
the variation in the surface roughness during the machin-
ing of the AA6061 aluminum alloy by considering the
cutting parameters (the spindle speed, the feed rate and
the depth of a cut) and the vibration criteria through the
neural-fuzzy method. Their model predicted the surface
roughness at the accuracy rate of 96 %. Chen and
Savage6 predicted the effects of the factors of the spindle
speed, the feed rate, the type of the tool material, the tool
diameter, and the vibration on the surface roughness in
Materiali in tehnologije / Materials and technology 46 (2012) 4, 383–388 383
UDK 621.9.025.7:669.715:620.179.11 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(4)383(2012)
the milling of AA6061 and AISI 1018 materials with the
neural-fuzzy approach. The suggested method predicted
surface roughness during cutting with a correlation of
90 %. Brezocnik et al.7 modeled the effect of the process
parameters of the spindle speed, the feed rate, the depth
of a cut and the vibration on the surface roughness in the
end milling of AA6061 alloy with a HSS plain-end mill
by using genetic programming.
Taguchi’s method can determine effects of the factors
on the quality characteristic with the smallest experiment
number without any need for complex calculations,
accomplishing the optimization of the system at a certain
confidence level. The method is very preferable in this
area because it requires only a small experiment number
and, as a result, a lower cost and shorter time. Yang and
Chen8 determined the effects of the factors of the spindle
speed, the feed rate and the tool diameter on the surface
roughness during the machining of AA6061 with a HSS
plain-end mill and they accomplished the optimization of
the system through Taguchi’s method. Lo et al.9 investi-
gated the effects of the factors of the spindle speed, the
feed rate, the depth of a cut and the tool material in the
high-speed milling of AA6061 by using Taguchi’s
method. They evaluated the results of the experiments
belonging to the L9 orthogonal array of Taguchi through
the S/N ratio and the variance analysis, and presented the
significance of the factors and their orders. Öktem et al.10
evaluated the data obtained in a 3-numbered study by
using Taguchi’s method and the full factorial experi-
ment-design methods for this time. They modeled the
surface roughness at the correlation coefficient of 0.96
by using regression analysis. They found, in the analyses
conducted with Taguchi, that the machining tolerance
was the most significant factor influencing the rough-
ness, followed by the radial depth of a cut, the axial
depth of a cut, the feed rate and the cutting speed per
thread. Pinar et al.1,11,12 investigated the effects of the
cutting speed, the feed rate, the depth of a cut and the
machining pattern on the surface quality in the milling of
the alloys of AA5083, 6013 and 7075 by using Taguchi’s
method. The experiments conducted on the basis of the
standard L27 orthogonal array of Taguchi were assessed
by using S/N, ANOVA and the main effects’ graphics so
that the system was optimized.
In this study, unlike other studies, the effect of the
aging process, which was imposed on the material, was
added to the effects of the other cutting factors (the
cutting speed, the feed rate, the axial depth of a cut and
the radial depth of a cut) influencing the surface
roughness to be analyzed by Taguchi’s method so that
the system was optimized.
2 MATERIAL AND METHOD
The AA 6013 alloy, whose chemical composition is
given in Table 1, was used as a test sample. The samples
to which precipitation hardening had, or had not, been
applied were used in the experiments. The pre-tests were
conducted to find the situation, in which the mechanical
properties were at their maximum. During the first step
of the precipitation hardening, the samples were solution
treated at 560 °C for 5 h. Then the quenching step was
completed. Finally, the samples were aged processed at
190 °C and their hardness was measured for 42 h once
every 2 h to find the optimum aging time. The hardness
tests were conducted by using Brinell hardness device.
The aging time achieving the maximum hardness was
determined to be 32 h (Figure 1). The end-milling tests
of the aged and non-aged samples were conducted in a
vertical machining center of First MCV 300 CNC
operating in line with Fanuc control unit. The single-
insert non-cemented carbide end mill with a 90-degree
incidence, a 14-mm diameter and a 0.8-mm insert
diameter, recommend by Mitsubishi Company in end
milling of aluminum, was used as the cutting tool
H. DURMUÞ: OPTIMIZATION OF MULTI-PROCESS PARAMETERS ACCORDING TO THE SURFACE ...
384 Materiali in tehnologije / Materials and technology 46 (2012) 4, 383–388
Figure 2: Cutting tool employed in the experiments
Slika 2: Rezno orodje, uporabljeno pri preizkusih
Table 1: Chemical composition of AA6013 experiment specimen
Tabela 1: Kemi~na sestava AA6013 raziskanega vzorca
Si Fe Cu Mn Mg Zn Cr Al
1.2 0.339 0.16 0.799 1.01 0.21 0.059 Geri kalan
Figure 1: Hardness graph employed in the determination of aging
time
Slika 1: Zapis trdote, uporabljen za dolo~itev ~asa staranja
(Figure 2). Machining operations were conducted by up
milling with liquid cooling using Ecocool 2030 MB with
a 5 % oil emission.
The CNC part programs of the samples were
obtained in MasterCAM V10 software as 2½-axial con-
tour machining. The cutting parameters were specified
on the basis of the manufacturer’s catalog data and bench
capacity. The surface-roughness measurements were
repeated three times perpendicular to the machining
traces through Mitutoyo Surftest SJ 311 profilometer
operating on the basis of the insert principle. Their
arithmetical means were used in the statistical analyses.
3 EXPERIMENTAL DESIGN AND STATISTICAL
ANALYSIS
Taguchi, which has become prevalently preferable in
recent years, only requiring experiments in a small
number and allowing system optimization without the
need for complex mathematical calculations, was used
for the experimental design and the statistical-analysis
method13,14.
The system design is the part, in which the factors
affecting the quality characteristic to be investigated and
their degrees are determined and it requires technical
data related to engineering knowledge in this area. The
parameter design is the most comprehensive and
important step of Taguchi’s method. In this phase, the
experiments conducted on the basis of the specified
experiment plan are analyzed, the optimum levels for the
factors are determined and the optimum dependent
variable is predicted for these levels. The final step is the
tolerance design, in which the average result of the
verification experiments is checked to find out whether
or not it is within the specified confidence interval12.
4 APPLICATION OF THE OPTIMIZATION
PROCESS
In the present study, the effects of the process para-
meters, which are the cutting speed, the feed rate, the
radial/axial depth of a cut and the aging, as well as their
dual interactions on an average surface roughness
(Ra/μm) are evaluated and the system is optimized. Table
2 shows the factors used in the experimental system and
their levels.
The OA establishing the experimental plan is selected
on the basis of the degree of freedom. Accordingly, the
freedom degree of the OA to be selected should be
higher than, or equal to, the freedom degree of the
system1. The freedom degree (FD) of the experimental
system is obtained by summing up the freedom degrees
of the factors and their interactions. The freedom degree
of the experimental system is specified according to the
number of the factor level interaction. The freedom
degree of a factor is specified as the level number of that
factor-1. Accordingly, because there are 5 factors, the
freedom degree of the factors is 5 × 1 = 5. In case of
interactions, individual freedom degrees of the factors
establishing interactions are multiplied to calculate the
overall freedom degree. Accordingly, because there are
10 interactions, their freedom degree is 10 × 1 = 10. The
freedom degree of the system is 15 being the summation
of 5 + 10. According to this data, L16 OR with the
freedom degree of 15, having 15 columns and 16 rows
was selected. If the same factors were investigated by
using the full factorial experimental design method, 25
experimental studies would be required. Thus, the saving
in terms of the number of experiments is 50 %.
Consequently, savings were obtained in time, sample
cutters and the power consumed by the bench. The
factors and their interactions were transferred to the
columns of OA by using the linear graphic method
(Figure 3). In line with this, the first column was
assigned to the cutting speed, while the second was
assigned to the feed rate. The forth column was assigned
to the axial depth of a cut, the eighth to the radial depth
of a cut, the fifteenth to the aging and the remaining
columns were assigned to the interactions.
H. DURMUÞ: OPTIMIZATION OF MULTI-PROCESS PARAMETERS ACCORDING TO THE SURFACE ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 383–388 385
Figure 3: Assignment of the factors and interactions to L27 OA by
means of linear graph
Slika 3: Porazdelitev faktorjev in interakcij za L27 OA z linearno
grafiko
Table 2: Choosen factors and their levels
Tabela 2: Izbrani faktorji in njihovi nivoji
Factors
Level
1 2
Cutting speed (A) m/min 100 300
Feed (B) mm/min 100 1200
Axial depth of cut (C) mm 0.75 2
Radial depth of cut (D) mm 5 10
Ageing treatment (E) – No Yes
4.1 Analyzing the data
The experimental results given in Table 3 were
evaluated by using the signal/noise ratio (S/N), the
variance analysis and the main effects’ graphic, with the
help of Minitab software, at the confidence level of
95 %.
4.1.1 S/N ratio analysis
Taguchi’s method uses the signal/noise ratio to
measure the existing variation. Description of the S/N
ratio varies according to the objective function or, in
other words, quality characteristic to be investigated.
Three different S/N ratios are used in the analyses as the
nominal values are the best (NB), smaller is better (SB)
and bigger is better. Since the average surface roughness
was investigated in the study as the quality characteristic
and the objective was a minimization of these para-
meters, the SB was selected as the NB. The following
equation refers to it.
S/N
n
y i
i
n
= −
⎡
⎣⎢
⎤
⎦⎥=
∑10 1 2
1
lg (1)
Herein, n represents the number of measurements
and yi is the measured roughness value. The S/N ratio
unit is a decibel. Table 3 shows the results of the experi-
ments conducted on the basis of the L16 orthogonal
array, the means and the corresponding S/N ratios.
4.1.2 ANOVA
In the study, the influence of the process parameters
on the quality characteristic and percentage distributions
were established through ANOVA. Tables 4 and 5 give
ANOVA results corresponding to the means and S/N
ratios. The significance of the process parameters and
their interactions are determined in ANOVA by com-
paring the F values (the ratio of the relevant parameter’s
variance to the error variance) in the fifth column of the
tables with the corresponding Ftable. If a process or an
H. DURMUÞ: OPTIMIZATION OF MULTI-PROCESS PARAMETERS ACCORDING TO THE SURFACE ...
386 Materiali in tehnologije / Materials and technology 46 (2012) 4, 383–388
Table 3: Experimental results, means and corresponding S/N ratios
Tabela 3: Eksperimentalni rezultati, povpre~ja in ustrezna razmerja S/N
Exp.
Number
Cutting
speed (A) Feed (B)
Axial
depth of
cut (C)
Radial
depth of
cut (D)
Ageing
treatment
(E)
Ra1 Ra2 Ra3 Raort S/N
1 100 100 0.75 5 No 0.58 0.58 0.58 0.580 4.731
2 100 100 0.75 10 Yes 0.76 0.76 0.8 0.773 2.230
3 100 100 2 5 Yes 0.8 0.8 0.8 0.800 1.938
4 100 100 2 10 No 0.84 0.81 0.81 0.820 1.722
5 100 1200 0.75 5 Yes 0.9 0.94 0.91 0.917 0.754
6 100 1200 0.75 10 No 0.92 0.95 0.94 0.937 0.568
7 100 1200 2 5 No 1.11 1.12 1.13 1.120 -0.985
8 100 1200 2 10 Yes 1.2 1.2 1.26 1.220 -1.730
9 300 100 0.75 5 Yes 0.38 0.38 0.38 0.38 8.404
10 300 100 0.75 10 No 0.4 0.4 0.4 0.4 7.959
11 300 100 2 5 No 0.62 0.64 0.65 0.637 3.920
12 300 100 2 10 Yes 0.7 0.66 0.67 0.677 3.390
13 300 1200 0.75 5 No 0.78 0.78 0.82 0.793 2.008
14 300 1200 0.75 10 Yes 0.8 0.8 0.84 0.813 1.792
15 300 1200 2 5 Yes 0.9 0.89 0.9 0.897 0.947
16 300 1200 2 10 No 0.9 0.9 0.92 0.907 0.851
Table 4: ANOVA results for means
Tabela 4: ANOVA rezultati za povpre~ja
Source SD SS V F Ftable KT’ P
Cutting speed (A) 1 0.172917 0.172917 44.94 4.75 0.169069 22.28
Feed (B) 1 0.402167 0.402167 104.51 4.75 0.398319 52.49
Axial depth of cut (C) 1 0.137517 0.137517 35.74 4.75 0.133669 17.62
Radial depth of cut (D) (1) (0.011201) Pooled – Pooled –
Ageing treatment (E) (1) (0.005017) Pooled – Pooled –
AxB (1) (0.000584) Pooled – Pooled –
AxC (1) (0.000034) Pooled – Pooled –
AxD (1) (0.003701) Pooled – Pooled –
BxC (1) (0.000851) Pooled – Pooled –
BxD (1) (0.000951) Pooled – Pooled –
Error 12 0.046176 0.003848 0.05772 7.61
Total 15 0.758777 0.758777 100
SD: Degree of freedom, SS: Sum of squares, V: Variance, KT’: Pure sum of squares, P: Percent of contribution
interaction parameter has a F value higher than this
value, it is deemed significant. Thus, it was observed that
the process parameters of the cutting speed, the feed rate
and the axial depth of a cut were significant, while their
interactions did not have a significant effect on surface
roughness in the cases of the two ANOVAs. Percentage
distributions are given in the last column of the ANOVA
table indicating significance degree of each of the
process parameters and dual interactions. On the basis of
this data, it can be established that the most significant
process parameter is the feed rate, followed by the
cutting speed and the axial depth of a cut. In Taguchi’s
method the factor levels indicating the optimum surface
roughness are determined by considering ANOVAs, the
main effects’ and interactions’ graphics. Significant
parameters are considered in both ANOVAs in the calcu-
lation of the optimum surface roughness. According to
this, the means of the parameters of the feed rate, the
cutting speed and the axial depth of a cut will be
considered. The levels to be selected are determined by
using the main effects’ and interactions’ graphics. As
none of the interactions in the experimental system is
significant, the main effects’ graphic will be sufficient in
this phase. Figure 4 shows the main effects’ graphic.
According to this graphic, surface roughness varies
depending on the cutting speed in inverse proportion and
also on the feed rate, the axial depth of a cut and the
radial depth of a cut in direct proportion. In the case of
the aging-state parameter, it was seen that surface rough-
ness increased a bit during the machining of the aged
samples. According to this data, the optimum surface
roughness was achieved at the cutting rate of 300 m/min,
a feed rate of 100 mm/min, an axial depth of a cut of
0.75 mm and a radial depth of a cut of 5 mm in non-aged
samples. The optimum surface roughness is obtained
through the following equation:
Pr = OA2 + OB1 + OC1 – 2K (2)
Herein, OA3 stands for the means from the experi-
ments conducted at the second level of the cutting speed.
OB1 stands for the means from the experiments con-
ducted at the first level of the feed rate. OC1 stands for the
means from the experiments conducted at the first level
of the axial depth of a cut. K stands for the means from
all experiments. Table 6 shows the relevant means.
According to Equation 2, the optimum surface roughness
was predicted to be 0.436 μm.
Table 6: Means of factor levels
Tabela 6: Nivoji povpre~nih faktorjev
Level
Cutting
speed
(A)
Feed
rate
(B)
Axial
depth of
cut (C)
Radial
depth of
cut (D)
Ageing
treatment
(E)
1 0.8958 0.6333 0.6992 0.7654 0.7742
2 0.6879 0.9504 0.8846 0.8183 0.8096
Delta 0.2079 0.3171 0.1854 0.0529 0.0354
Rank 2 1 3 4 5
Overall average (K)=0.792μm
In the last step of Taguchi, called the tolerance
design, the confidence interval is determined, in which
the upper and the lower limits of the existing predicted
surface roughness are determined. The means of the
roughness results from the conducted verification tests
should be within this interval. The following equation is
used to calculate the confidence interval12,14.
H. DURMUÞ: OPTIMIZATION OF MULTI-PROCESS PARAMETERS ACCORDING TO THE SURFACE ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 383–388 387
Figure 4: Main effects plot of the factors
Slika 4: Glavne povezave faktorjev
Table 5: ANOVA results for S/N ratios
Tabela 5: ANOVA rezultati za razmerje S/N
Source SD SS V F Ftablo SS’ P
Cutting speed (A) 1 25.104 25.1044 27.42 4.75 24.188 20.85
Feed (B) 1 56.584 56.5838 61.80 4.7 55.668 48.00
Axial depth of cut (C) 1 21.144 21.1439 23.09 4.75 20.228 17.44
Radial depth of cut (D) (1) 1.524 - Pooled -
Ageing treatment (E) (1) 0.581 - Pooled -
AXB (1) 2.295 - Pooled -
AXC (1) 0.865 - Pooled -
AXD (1) 0.348 - Pooled -
BXC (1) 2.493 - Pooled -
BXD (1) 0.375 - Pooled -
Error 12 10.988 0.9157 15.903 13.71
Total 15 115.987 115.987 100
SD: Degree of freedom, SS: Sum of squares, V: Variance, KT’: Pure sum of squares, P: Percent of contribution
CI F V
n R
= ⋅ +
⎡
⎣⎢
⎤
⎦⎥
⎛
⎝
⎜
⎞
⎠
⎟
CI e
eff
1 1
1 2/
(3)
Herein, FCI is obtained from the standard tables at the
significance level of 0.05 according to the freedom
degree of 1 including an error and it is 4.75. Ve, is the
variance value of the error in the ANOVA table and
calculated as 0.003848 (Table 4). R is how many times
the verification tests are repeated under optimum con-
ditions (3) and finally, neff is the effective repeat number.
This was obtained by using the following equation.
n
N
Veff t
=
+1
(4)
Herein, N is the total number of the experiments in
the experimental plan (16), while Vt is the total freedom
degree of the factors existing in the equation, through
which the optimum surface roughness is predicted (3).
Accordingly, CI was calculated to be 0.103 at the
confidence level of 95 % or at the significance level of
0.05. The mean of the results of the verification tests
repeated three times should be within the interval of
0.333 < Ra < 0.539. The average surface roughness was
determined to be 0.343 μm (0.35, 0.34 and 0.34 μm
respectively) in the verification tests conducted under the
optimum conditions. According to this data, it is clearly
seen that Taguchi’s method optimized the system
successfully.
5 CONCLUSION
In the present study, the effects of multi-process
parameters (the cutting speed, the feed rate, the axial
depth of a cut, the radial depth of a cut and the aging
state) and their dual interactions on the surface
roughness in the end-milling operations were evaluated
and the following conclusions were achieved:
Surface roughness varies depending on the cutting
speed in inverse proportion and on the feed rate, the axial
depth of a cut and the radial depth of a cut in direct
proportion. It was seen that surface roughness increased
a bit during the machining of aged samples.
Only the cutting speed, the feed rate and the axial
depth of a cut are significant for both ANOVAs. It is seen
that the most significant process parameter is the feed
rate followed by the cutting speed and the axial depth of
a cut.
The optimum surface roughness was predicted to be
0.436 μm found in the case of the second cutting speed
(A = 300 m/min), the first feed rate (B = 100 mm/min),
the first axial depth of a cut (C = 0.75 mm), the first
radial depth of a cut (D = 5) and under the non-aging
conditions. The average surface roughness was found to
be 0.343 μm in the verification tests that were repeated
three times, and it existed within the specified confi-
dence interval. Thus, the system was optimized success-
fully. Furthermore, the minimum surface roughness was
found to be 0.38 μm in the nine-number experiment in
the experimental plan (A = 300 m/min, B = 100 mm/min,
C = 0.75 mm, D = 5 mm E: yes). With the result
obtained during the verification tests an improvement of
9.74 % was obtained.
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K. STRANSKY et al.: NUMERICAL AND EXPERIMENTAL ANALYSES OF THE CHEMICAL HETEROGENEITY ...
NUMERICAL AND EXPERIMENTAL ANALYSES OF THE
CHEMICAL HETEROGENEITY OF A SOLIDIFYING
HEAVY DUCTILE-CAST-IRON ROLLER
NUMERI^NA IN EKSPERIMENTALNA ANALIZA KEMIJSKE
HETEROGENOSTI STRJEVANJA LITEGA TE@KO GNETLJIVEGA
@ELEZNEGA VALJA
Karel Stransky1, Frantisek Kavicka1, Bohumil Sekanina1, Jana Dobrovska2,
Vasilij Gontarev3
1Brno University of Technology, Technická 2, 616 69 Brno, Czech Republic
2VSB TU Ostrava, 17.listopadu 15, 708 33 Ostrava, Czech Republic
3University of Ljubljana, A{ker~eva 12, 1000 Ljubljana, Slovenia
stransky@fme.vutbr.cz
Prejem rokopisa – received: 2011-10-17; sprejem za objavo – accepted for publication: 2012-02-24
The quality of the working rollers used for rolling rails is determined by the chemical and structural compositions of the
material of the rollers and the production technology. It is necessary to cast rollers with significantly improved utility properties,
i.e., mainly a high wear resistance and optimal mechanical and structural properties. It is, therefore, necessary to find and ensure
the optimal relationships between the matrix structure and the resulting values of the mechanical properties of the rollers in
order to maximize their life time. The requirements introduced here cannot be ensured without a knowledge of the kinetics of
the solidification. Therefore, numerical and experimental investigations of the temperature field of the solidifying roller were
conducted. The kinetics of the solidification has a measurable and non-negligible influence on the chemical and structural
heterogeneity of the investigated type of ductile cast-iron. Linking to the results of the model of the temperature field of the cast
rollers, an original methodology was developed for the measurement of chemical micro-heterogeneity. The structure of this
cast-iron is created by a large amount of the transition form of graphite and a small amount of globular graphite, and also the
lamellar graphite and cementite, whereas the structure of the metal matrix is perlitic. The volume amounts of the structural
components were determined using a quantitative metallographic analysis, according to which the places for the analysis of the
element composition using X-ray energy-dispersive spectral micro-analysis were selected. The chemical and structural
heterogeneity of the cast roller is, therefore, a significant function of the method of melting, modification and inoculation and
the successive procedures of risering, casting and crystallization after cooling.
Keywords: spheroidal graphite cast iron, roller, solidification, chemical and structural heterogeneity, methodology for
measurement
Kakovost delovnih valjev za valjanje tirnic je dolo~ena s kemijsko in strukturno sestavo materiala za valje in s tehnologijo
proizvodnje. Potrebno je ulivati valje z znatno pove~animi koristnimi lastnostmi, kot so obrabna odpornost in optimalne
mehanske in strukturne lastnosti. Tako je treba ugotoviti in zagotoviti optimalne odnose med strukturo matrice in rezultirajo~imi
vrednostmi mehanskih lastnosti valjev za maksimiranje trajnostne dobe. Zahteve, ki so tu predstavljene, ne morejo biti
zagotovljene brez znanja kinetike strjevanja. Zato je bila vpeljana numeri~na in eksperimentalna raziskava temperaturnega polja
valja pri strjevanju. Kinetika strjevanja ima merljiv in ne brezpomemben vpliv na kemijsko in strukturno heterogenost
raziskovanega gnetljivega litega `eleza. Glede na rezultate modela temperaturnega polja ulitih valjev je bila razvita originalna
metodologija za merjenje kemijske mikroheterogenosti. Struktura tega ulitega `eleza je bila ustvarjena z veliko koli~ino
prehodnih oblik grafita in manj{ih koli~in kroglastega grafita in cementita, medtem ko je struktura kovinske osnove perlitna.
Volumenska koli~ina strukturnih komponent je bila dolo~ena s kvantitativno metalografsko analizo, po kateri so bila izbrana
mesta za rentgensko analizo sestave ter spektrografsko mikroanalizo. Kemijska in strukturna heterogenost ulitih valjev je
pomembna funkcija pri taljenju in modificiranju ter za uspe{no ulivanje in kristalizacijo po ohlajevanju.
Klju~ne besede: siva litina s kroglastim grafitom, valj, strjevanje, kemijska in strukturna heterogenost, metodologija meritev
1 INTRODUCTION
The kinetics of solidification has a non-negligible
influence on the chemical and structural heterogeneity of
the cast-iron in question1. An original methodology for
the measurement of the micro-heterogeneity was
developed, based on the results of the model of the
temperature field of the cast rollers. The chemical and
structural heterogeneity of the cast roller has proven to
be a significant function of the method of melting, modi-
fication and inoculation and the successive procedures of
risering, casting and crystallization after cooling.
2 STRUCTURAL AND CHEMICAL
HETEROGENEITY OF THE ROLLER
The final mechanical properties of the rollers – of the
spheroidal graphite cast-iron – are determined, not
primarily by the chemical composition, but mainly by
their structural and chemical heterogeneity, which occurs
during casting, crystallization and successive cooling of
the material. Some defects occurring in this way can be
corrected by heat treatment; however, the quality of the
pouring structure is very important, especially with
graphite cast-iron.
Materiali in tehnologije / Materials and technology 46 (2012) 4, 389–392 389
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The chemical composition is given in Table 1.
The samples used for determining the structural and
chemical heterogeneity were taken from the spheroidal
graphite cast-iron roller with the chemical composition
given in Table 1 from the upper part (Section III) and
from the bottom part (Section I). Figure 1 shows how
the four samples were taken, with the upper surfaces –
marked "X" – being analyzed. In addition, they were
specified as follows:
• The upper part of the roller on the outer surface (R1 =
540 mm) marked TR1, TR2
• The upper part of the roller on the inner surface (R4 =
400 mm) marked TR3, TR4
• The lower part of the roller on the outer surface
marked BR1, BR2
• The lower part of the roller on the inner surface
marked BR3, BR4
The actual analyses were conducted at a specialized
workplace2,3. The structural analysis of the cast-iron
samples was conducted using a Neophot 32 metallo-
graphic microscope and an Olympus digital camera. The
measurements of the metallographic parameters of the
graphite and the evaluation of the volume part of the
structural components were conducted on the Olympus
CUE4 image analyzer. The basic statistical parameters of
the graphite particles and also the stereological estimate
of certain other parameters2 were calculated from the
data files.
The JSM-840 (JEOL) electron scanning microscope
was used for determining the chemical heterogeneity
together with the LINK AN 10/85S X-ray energy-disper-
sive microanalyzer. Based on the results of the
metallographic analysis2, two pairs of graphite grains
were selected on the ground surface of every sample in
order for the line of measurement between the first pair
to pass through the perlite (marked "a") and the cemen-
tite (marked "b"). The third line was selected through
two boundaries created by particles of lamellar graphite,
where the measurement passes through the basic perlite
(marked "c"). The individual elements (Si, P, Mo, Cr,
Mn, Fe, Ni, and Cu) were analyzed by means of a point
X-ray microanalysis in a direct line with a step of 3 μm 3,4.
Figure 2 illustrates the material around the selected line
of the measurement on sample TR1 through the scanning
electron microscope and the X-ray energy-dispersive
microanalyzer. Following the chemical (elemental)
analysis, the samples were etched using 2 % nital in
order to make the contamination traces visible and
display the line of measured points, including their
connection to the sample microstructure.
2.1 Structural heterogeneity
Based on experience, it can be assumed that the
structure of the material of the casting, spheroidized
graphite cast iron (named "ductile cast-iron" for working
purposes) will, besides the globular graphite that is
characteristic of ductile cast-iron, also contain a mixture
of a certain amount of transition forms of graphite and
lamellar graphite found between the globular and
lamellar graphite. Table 2 shows the results from the
quantitative metallographic analyses of the structures of
the samples in Figure 1. The procedure of the quanti-
tative structural measurements, including the results, is
described in detail in a special report2.
K. STRANSKY et al.: NUMERICAL AND EXPERIMENTAL ANALYSES OF THE CHEMICAL HETEROGENEITY ...
390 Materiali in tehnologije / Materials and technology 46 (2012) 4, 389–392
Table 1: The chemical composition of ductile cast-iron B10 (w/%)
Tabela 1: Kemijska sestava sive litine B10, z globularnim grafitom (w/%)
Cast-iron B10
Element C Mn Si P S Cr Ni Mo Mg Cu
Content (%) 3.31 0.65 0.70 0.105 0.005 0.35 2.59 0.59 0.04 1.48
Table 2: The results from the quantitative measurements of the volume part of the components (in volume fractions, /%)
Tabela 2: Rezultati kvantitativnih meritev volumenskega dela komponent (volumenski dele`i, /%)
Sample All graphite, incl.micro-shrinkages Globular graphite
Transition forms of graphite
and lamellar graphite Cementite Perlite
BR1 10.5 ± 2.8 1.3 9.2 2.2 ± 1.7 87.3
BR3 13.0 ± 5.4 1.1 11.9 1.2 ± 1.2 85.8
TR1 13.4 ± 5.6 1.1 12.3 2.2 ± 1.8 84.4
TR3 12.6 ± 5.7 1.1 11.5 1.3 ± 1.4 86.1
Figure 1: The set-up of a vertically cast roller and sampling
Slika 1: Pogled na navpi~no ulit valj in vzor~enje
2.2 Chemical heterogeneity
The distribution of the concentration of the eight
elements, measured three times in one line for each of
the samples (Figures 2a–c), was evaluated using
statistics: the arithmetic mean of the concentration of the
element in the selected interval cst, the standard deviation
n–1 of the measured concentration, the minimum
concentration cmin and the maximum cmax – measured
each time in the selected interval of the sample (in mass
fractions, w/%). Besides these basic concentration
parameters, the segregation index of each measured
element was determined and defined as IS = cmax/cst, i.e.,
as the quotient of the maximum concentration of
elements measured on the given interval and the arith-
metic mean of its concentration in the same interval. The
tendency of the element to segregate – based on
experience – is expressed in the values of the segregation
indexes. A detailed description of the measurement
procedure and the corresponding results can be found in
the reports3,4.
Figure 3 illustrates the values of this parameter for
individual elements and samples measured in one line
across the lamellar graphite – c. Similarly, the segre-
gation index values can be plotted for the area between
the globular graphite in a line across perlite – a, and also
in the area between the globular graphite in a line across
cementite – b. Figure 4 shows the segregation index
values for the elements calculated as the average of the
values for all three areas – a, b, c. Furthermore, the
graphs in Figures 3 and 4 also indicate that the
segregation indexes of the highest values – from the set
of measured elements – belong to phosphorus and
molybdenum, and the lowest values refer to nickel and
iron, which make up the matrix.
3 CONCLUSION
This article introduces an original methodology for
the measurement of the chemical heterogeneity of
cast-iron of the composition given in Table 1. The
structure of this cast-iron is created by a large amount of
the transition form of graphite and small amount of
K. STRANSKY et al.: NUMERICAL AND EXPERIMENTAL ANALYSES OF THE CHEMICAL HETEROGENEITY ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 389–392 391
Figure 2: The structures in the area of sample TR1 selected for analysis: a) between two spheres of graphite through perlite, b) between the two
spheres of graphite through cementite, c) through lamellar graphite
Slika 2: Strukture v obmo~ju vzorca TR1 za analizo: a) med dvema kroglama grafita v perlitu, b) med dvema kroglama grafita v cementitu, c) v
lamelnem grafitu
Figure 4: The average index of segregation of the analyzed elements
in individual samples
Slika 4: Povpre~ni indeks izcejanja analiziranih elementov v posa-
meznih vzorcih
Figure 3: The index of the segregation of elements for individual
samples in the area of the transition graphite in the measured area
across lamellar graphite
Slika 3: Indeks izcejanja elementov posameznih vzorcev v obmo~ju
prehoda grafita v merjeno podro~je lamelnega grafita
globular graphite and also lamellar graphite and cemen-
tite, whereas the structure of the metal matrix is perlitic.
The volume amounts of the structural components were
determined using a quantitative metallographic analysis,
which was simultaneously the basis for the selection of
the places for the analysis of the element composition
using an X-ray energy-dispersive spectral microanalysis.
Inside the microstructure of four samples, taken from the
outer and inner parts of the ring at the top and bottom of
the roller (Figure 1), in a total of twelve areas, the linear
concentration dispersion was measured on the eight
elements which make up the basic constitution of the
cast-iron: Si, P, Mo, Cr, Mn, Fe, Ni, and Cu. A statistical
analysis was used to determine the distribution characte-
ristics of the concentration of the individual elements,
including the segregation indexes and the relationships
among them. In this way it was possible to assess the
influence of the changing kinetics of the temperature
field on the resultant structural and chemical hetero-
geneity of the roller. The kinetics of the solidification of
the cast-iron roller had a non-negligible influence on the
chemical and structural heterogeneity of the investigated
type of cast-iron. It is obvious that the chemical and
structural heterogeneity of the cast roller is a significant
function of the method of melting, modification and
inoculation and the successive procedures of risering,
casting and crystallization after cooling.
Acknowledgements
This analysis was conducted using a program devised
within the framework of the GACR projects No.
106/08/0606, 106/09/0940, 106/09/0969 and P107/11/
1566.
4 REFERENCES
1 F. Kavicka et al.: Numerical optimization of the method of cooling of
a massive casting of ductile cast-iron. In: Book of Abstracts and CD
ROM of the 13th International Heat Transfer Conference, Sydney,
Australia, August 2006, 27
2 J. Belko, K. Stransky, Structural analysis of cast-iron samples (in
Czech). Research report, VTUO Brno, 2005
3 Z. Winkler, K. Stransky, Microheterogeneity of the composition
ductile cast-iron samples alloyed with Mn, Cu, Ni and Mo (in
Czech). Research report, VTUO Brno, 2005
4 R. Kamensky, et al.: The influence of silicon on the thickness of the
hardened surface layer of cast-iron castings (in Czech). Slevarenstvi,
16 (1968) 1, 35–39
K. STRANSKY et al.: NUMERICAL AND EXPERIMENTAL ANALYSES OF THE CHEMICAL HETEROGENEITY ...
392 Materiali in tehnologije / Materials and technology 46 (2012) 4, 389–392
B. SARIKAN et al.: INVESTIGATION ON THE AGING BEHAVIOUR OF THE FUNCTIONALLY GRADIENT ...
INVESTIGATION ON THE AGING BEHAVIOUR OF THE
FUNCTIONALLY GRADIENT MATERIAL CONSISTING
OF BORON CARBIDE AND AN ALUMINUM ALLOY
RAZISKAVA PONA[ANJA PRI STARANJU FUNKCIONALNIH
GRADIENTNIH MATERIALOV IZ BOROVEGA KARBIDA IN
ALUMINIJEVE ZLITINE
Bertan Sarýkan1, Erhan Balcý1, Mustafa Übeyli2, Necip Camuºcu1
1TOBB University of Economics and Technology, Mechanical Engineering, Sö

gütözü Cad. No:43 06560 Ankara, Turkey
2Osmaniye Korkut Ata University, Mechanical Engineering, Karacao

glan Yerleºkesi, 80000 Osmanýye
mubeyli@gmail.com
Prejem rokopisa – received: 2011-10-18; sprejem za objavo – accepted for publication: 2012-03-05
In the current study the effect of different temperatures on the aging behaviour of a functionally gradient material was
investigated to observe the variation of hardness with respect to the aging time. To do this, the functionally-gradient-material
(FGM) specimens, containing boron carbide and aluminum alloy, were produced with hot pressing. Three different layers were
used in the FGM samples. Then the macro and micro examinations were carried out to observe the interface between the layers,
the porosity, the grain size and probable cracks in these samples. After that the specimens were solutionized at 470 °C for 1 h.
Next, the artificial aging at 100 °C, 120 °C and 150 °C was applied to the FGM specimens for 96 h. During the aging treatment,
the Brinell hardness measurements were made on the samples at certain intervals. Moreover, three-point bending tests were also
carried out to clarify the influence of the aging treatment on the strength of the FGM. Experimental results indicated that the
highest hardness values were obtained at 120 °C after the aging period of 48–65 h. The aged specimens exhibited higher
bending strength than the solutionized specimens.
Keywords: metallic materials, inorganic materials, aging, three-point bending test
V tej {tudiji je prikazan vpliv razli~nih temperatur na staranje funkcionalnega gradientnega materiala in spreminjanje trdote v
odvisnosti od ~asa staranja. Vzorci funkcionalno gradientnih materialov (FGM), ki vsebujejo borov karbid in aluminijevo
zlitino, so bili izdelani s stiskanjem v vro~em. Trije razli~ni sloji so bili uporabljeni za FGM-vzorce. Na vzorcih so bile izvr{ene
makro- in mikropreiskave stikov med sloji, dolo~ene so bile poroznost, velikost zrn in prisotnost morebitnih razpok. Potem so
bili vzorci 1 h raztopno `arjeni na 470 °C, nato pa umetno starani 96 h na temperaturah 100 °C, 120 °C in 150 °C. Med
postopkom staranja so bile v dolo~enih intervalih izvr{ene meritve trdote po Brinellu. Dodatno je bil izvr{en {e trito~kovni
upogibni preizkus, da bi opredelili vpliv staranja na trdnost FGM. Rezultati preizkusov ka`ejo, da je bila dose`ena najvi{ja
trdnost pri staranju med 48 h in 65 h na temperaturi 120 °C. Starani vzorci so pokazali ve~jo upogibno trdnost kot raztopno
`arjeni vzorci.
Klju~ne besede: kovinski materiali, anorganski materiali, staranje, trito~kovni upogibni preizkus
1 INTRODUCTION
A new type of material called functionally gradient
material (FGM) has been considered to be used for
special purposes. In this type of material, at least two
different types of layers are chemically used in a gradient
form. In this way the properties change throughout the
cross-section of the material. By definition, an FGM
consists of one material being on one side and another
material on the other side1–21. FGMs have been used in
structural, electrical, chemical, optical, nuclear and
biomedical applications1. The main methods for
producing these materials are powder metallurgy,
thermal spraying, coating, laser coating and other
thermo-mechanical processes1,2. Among these methods,
powder metallurgy techniques are very suitable for the
FGM production1,2. The main advantages of chemically
forming a gradient in a material are the creation of a
metallurgical bond between the layers and different
mechanical, nuclear, electrical and/or thermal properties
of individual layers1–20. Producing a hard outer layer and
a soft inner layer in the same material has been of key
importance in the armour applications1–3. Therefore, the
FGMs containing a hard front layer and a tough inner
layer can be suitable for these applications. However,
there is a large gap in the literature related to the
production and mechanical properties of FGMs that need
to be investigated extensively with respect to the
objectives of materials science and manufacturing
methods. In this paper an FGM with three layers,
chemically changing from the top to the bottom, was
produced using the powder metallurgy technique. In this
material the bottom layer was chosen to be the aluminum
alloy (AA) 7075, whereas the boron-carbide particles
reinforcing the AA 7075 alloy were used, in different
proportions, for the middle and the top layers. The
influence of the artificial-aging treatments on the
hardness of the layers was examined at different aging
temperatures of 100 °C, 120 °C and 150 °C. In addition,
the bending strength of the FGMs was determined for the
solutionized and the aged samples.
Materiali in tehnologije / Materials and technology 46 (2012) 4, 393–397 393
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2 EXPERIMENTAL PROCEDURE
The FGM samples with three layers were produced
with the powder metallurgy technique. The powders used
as starting materials were supplied from the market. The
mean powder sizes of aluminum, boron carbide, chro-
mium, copper, iron, silicon, zinc and magnesium were
measured to be 10.22 μm, 4.04 μm, 33.42 μm, 19.17 μm,
6.66 μm, 3.53 μm, 6.7 μm and 45.08 μm, respectively.
Furthermore, the powders had very high purities
alternating between 99.5 % and 100 %. In the investi-
gated aluminum alloy, the nominal chemical composition
was taken as mass fractions 5.5 % Zn, 2.5 % Mg, 1.6 % Cu,
0.5 % Fe, 0.4 % Si, 0.23 % Cr. The powder size is very
important when considering the sintering operation. A
small powder size results in a higher surface area, which
effectively enhances the sintering operation22. The FGM
samples were based on the AA 7075 that can be
hardened to a very high level by an artificial aging
treatment23–25.
In the investigated FGM samples, the bottom layer
was considered to be a monolithic AA 7075 layer, while
20 % and 40 % B4C (in volume) of the particle reinforce-
ments in the AA 7075 matrix were used in the middle
and top layers, respectively. Initially, the layers in the
suitable chemical proportions were pre-shaped
separately by cold pressing. Then they were transferred
into the mould of a uni-axial hot press to be sintered at
590 °C for 20 min. After the sintering operation, the
artificial aging treatments were performed to get the
hardness-aging time curves for the produced FGM.
Before the aging process, the samples were solutionized
at 470 °C for 1 h and water quenched at 20 °C. The
aging temperatures were 100 °C, 120 °C and 150 °C.
Furthermore, the macro and the micro examinations on
the samples were done to observe probable production
defects such as cracks, pores and separations between
the layers, as well as distribution of ceramic particles in
the matrix. For the metallographic analysis, the samples
were prepared with the standard methods and etched
using the etchant, a modified Keller’s reagent. The
hardness tests and the three-point bending tests were
made on the solutionized and peak-aged samples accord-
ing to the standards of ASTM E 10-84 26 and ASTM B
528-05 27, respectively. For the hardness testing, the
samples were taken out of the furnace and water
quenched for the hardness measurement at certain
intervals during the aging treatment (every half an hour
during the first 6 hours and every hour during the period
between the 7th and the 96th hour). Each time the mean
values of five measurements were recorded.
3 RESULTS AND DISCUSSION
3.1 Macro- and microstructural observations
Figure 1 shows a typical macro view of the cross
section for a produced sample with the total thickness of
15 mm. The three layers can be seen clearly in this
figure. There are no macro-level cracks or separations
observed between the layers of the produced sample. The
adhesion of the layers with various compositions appears
to be smooth. Figures 2 and 3 illustrate the microstruc-
ture of the produced FGM taken by an optical micro-
scope. A uniform distribution of ceramic particles can be
observed on these figures. In addition, there is a good
transition-and-interface formation between the layers.
Although the formation of interfaces and the
distribution of ceramic particles throughout the samples
are found to be uniform, some porosity is detected. The
porosity level, detected by an image-analysis program, is
estimated to be 5 %, 6 % and 8 % for the bottom, middle
and top layers, respectively. Moreover, the average grain
B. SARIKAN et al.: INVESTIGATION ON THE AGING BEHAVIOUR OF THE FUNCTIONALLY GRADIENT ...
394 Materiali in tehnologije / Materials and technology 46 (2012) 4, 393–397
Figure 2: Micro view of: a) the middle and b) the top layers
Slika 2: Mikroposnetek: a) sredine in b) vrhnjih plasti
Figure 1: Typical cross sectional view of a produced FGM sample
with the total thickness of 15 mm
Slika 1: Zna~ilen prerez FGM-vzorca s skupno debelino 15 mm
size was measured to be 24 μm according to the ASTM
E 112 28.
3.2 Aging effects on the hardness
Figure 4 gives the hardness/aging time curves for the
sample aged at 100 °C. Initially, the hardness values for
the bottom, middle and top layers are HB (104, 148 and
218), respectively. The hardness of all the layers
increases after 30 min. Later there are some fluctuations
in the curves in the aging period of the first 6 h. Finally,
the hardness is stabilized for all three layers sometime
between the 6th and the 96th hour. The highest hardness
values recorded for the bottom, middle and top layers are
HB (142, 172 and 242), respectively.
The aging curves for the sample aged at 120 °C are
given in Figure 5. As expected, the highest hardness
value belongs to the top layer. The general tendency in
the shape of the curve is very similar for all three
different layers. After the solutionizing treatment, the
hardness is measured to be HB (104, 149 and 218) for
the bottom, middle and top layers, respectively. During
the aging period of the first 4–5 h, there are some
fluctuations in the hardness profile. However, later the
hardness starts to increase reaching the peak levels
between the 48th and the 66th hour for all the layers.
During the period between the 66th and the 96th hour, a
gradual reduction in the hardness is observed due to over
aging. The highest hardness values measured for the
bottom, middle and top layers are HB (145, 181 and
247), respectively.
Figure 6 depicts the change in the hardness of the
sample with respect to time at the aging temperature of
150 °C. As in the aging treatments at 100 °C and 120 °C,
there is no change in the hardness profile of the alloy
including the B4C ceramic particles. After the aging
period of 12 hours, the hardness is recorded to be HB
(135, 171 and 239) for the bottom, middle and top
layers, respectively. The hardness begins to decrease
after the 12 h aging period. At the end of the 66th hour,
the hardness becomes HB (116, 142 and 210) for the
same layers, respectively. When aging the sample at 150
B. SARIKAN et al.: INVESTIGATION ON THE AGING BEHAVIOUR OF THE FUNCTIONALLY GRADIENT ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 393–397 395
Figure 5: Hardness versus the aging time at 120 °C for: a) bottom, b)
middle and c) top layers
Slika 5: Trdota v odvisnosti od ~asa staranja pri 120 °C za: a) spodnji,
b) srednji in c) zgornji sloj
Figure 3: Interfaces between: a) the bottom and the middle, b) the
middle and the top layers
Slika 3: Stik med: a) spodnjo in srednjo, b) srednjo in zgornjo plastjo
Figure 6: Hardness versus the aging time at 150 °C for: a) bottom, b)
middle and c) top layers
Slika 6: Trdota v odvisnosti od ~asa staranja pri 150 °C za: a) spodnji,
b) srednji in c) zgornji sloj
Figure 4: Hardness versus the aging time at 100 °C for: a) bottom, b)
middle and c) top layers
Slika 4: Trdota v odvisnosti od ~asa staranja pri 100 °C za: a) spodnji,
b) srednji in c) zgornji sloj
°C, the peak hardness values are obtained in shorter
times, but they are lower than those obtained at the
temperature of 120 °C.
The AA 7075 is a heat-treatable alloy23 and its
precipitation sequence can be given as follows24:
Super saturated solid solution ;
; GP zones ; ’ ; -MgZn2
It is considered that the highest hardness is reached
when the metastable phase of ’ is formed during the
aging23,25. The GP zones and ’ form at the beginning of
the aging treatment. The size and density of the preci-
pitates very much affect the mechanical properties23,25. In
the current work the higher hardness levels are obtained
at the aging temperature of 120 °C. Therefore, it is
thought that an extensive formation of the ’ phase takes
place at this temperature. Although the precipitation
kinetics seem to be higher at 150 °C, the hardness values
recorded at this temperature were lower than those
obtained at 120 °C.
In a previous study29 on the aging of AA 7075, the
peak hardness value (
HV 190) was reached at 121 °C
and after 48 h. The general trend of the hardness profile
and the hardness values found in this study are similar to
those found in ref.29. Furthermore, in another study30 the
highest hardness value was recorded to be HV 170 at
120°C after 48 h for the AA 7075 produced with powder
metallurgy. This result is also consistent with the current
data given in this work.
3.3 Bending strength
During the bending test the force was applied through
the layer with the 40 % B4C particle reinforcement. The
bending strength of the solutionized sample was found to
be 456 MPa, whereas the bending strength of the aged
sample was measured to be 527 MPa. A remarkable
effect of the aging on the bending strength of the FGM
was observed. During the testing a major crack was
formed through the main axis that was exposed to the
bending force and some particle fractures occurred on
the bottom layer of the non-aged sample that has lower
hardness (Figure 7). On the other hand, no big crack
formation was observed on the aged sample, as only a
fracture of some small particles were detected on its
bottom layer. The rigidity of this sample increased with
an increase in its hardness.
4 CONCLUSIONS
The peak hardness values for all the layers of the
investigated FGMs were found at the aging temperature
of 120 °C. The artificial aging of the FGM at 120 °C
allowed about 40 %, 21 % and 13 % increase in the
hardness of the bottom, middle and top layers,
respectively. The addition of the B4C ceramic particles to
the AA 7075 matrix had no significant effect on the
aging behaviour of the alloy. The aged specimen
exhibited a 15 % higher bending resistance than the
non-aged one. In order to enhance the hardness and
strength of the investigated FGMs, the porosity level
should be reduced or eliminated using higher pressure
and smaller powder sizes.
Acknowledgement
This work was supported by the Research Fund of
Tübitak, Project # 110M034. The authors are thankful to
Tübitak for its support.
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functionally graded material by powder metallurgy, Mater. Chem.
Phys., 68 (2001), 130–135
7 L. M. Zhang, J. Liu, R. Z. Yuan, T. Hirai, Properties of TiC-Ni3Al
composites and structural optimization of TiC-Ni3Al functionally
gradient materials, Mat. Sci. Eng. A, 203 (1995), 272–277
8 C. Y. Lin, H. B. Bathias, H. B. Mcshane, R. D. Rawlings, Production
of silicon carbide Al2124 alloy functionally graded materials by
mechanical powder metallurgy technique, Powder Metall., 42 (1999),
29–34
9 Y. Miyamoto, Z. Li, Y. S. Kang, K. Tanihata, Fabrication and pro-
perties of alumina-based hyperfunctional ceramics with symmetri-
cally compositional gradient structures, Powder Metall., 42 (1995),
933–938
10 Y. G. Jung, U. Paik, S. C. Choi, Influence of the particle size and
phase type of zirconia on the fabrication and residual stress of
zirconia/stainless steel 304 functionally gradient material, J. Mater.
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Figure 7: Macro view of the non-aged sample with the total thickness
of 15 mm after the three-point bending test
Slika 7: Makrovidez nestaranega vzorca s skupno debelino 15 mm po
trito~kovnem upogibnem preizkusu
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multilayered mullite/Mo functionally graded materials fabricated by
powder metallurgy processing, Mater. Chem. Phys., 89 (2005),
238–243
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functionally graded materials, Mat. Sci. Eng. A, 362 (2003), 81–105
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Materiali in tehnologije / Materials and technology 46 (2012) 4, 393–397 397

M. TKADLECKOVA et al.: SETTING A NUMERICAL SIMULATION OF FILLING AND SOLIDIFICATION ...
SETTING A NUMERICAL SIMULATION OF FILLING
AND SOLIDIFICATION OF HEAVY STEEL INGOTS
BASED ON REAL CASTING CONDITIONS
POSTAVITEV NUMERI^NE SIMULACIJE POLNJENJA IN
STRJEVANJA VELIKIH JEKLENIH INGOTOV NA PODLAGI
REALNIH RAZMER PRI ULIVANJU
Marketa Tkadleckova1, Pavel Machovcak2, Karel Gryc1, Petr Klus1,
Karel Michalek1, Ladislav Socha1, Marek Kovac3
1V[B – Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of Metallurgy and
Regional Materials Science and Technology Centre, Ostrava-Poruba, Czech Republic
2VÍTKOVICE HEAVY MACHINERY, a. s., Ostrava-Vítkovice, Czech Republic
3MECAS ESI, s. r. o., Brno, Czech Republic
marketa.tkadleckova@vsb.cz
Prejem rokopisa – received: 2011-10-22; sprejem za objavo – accepted for publication: 2012-01-06
The paper is devoted to new experiences with the setting of a numerical simulation of filling and solidification of a 90-ton steel
ingot in the ProCAST simulation programme. The aim of the numerical modelling realized under the conditions of the
Department of Metallurgy and Regional Materials Science and Technology Centre (RMSTC) at VSB-TU Ostrava is the
verification and optimization of the production technology for the heavy-steel ingots produced in VÍTKOVICE HEAVY
MACHINERY a.s. The input parameters of the computation were determined with the real conditions of casting a 90-ton steel
ingot. The ingot geometry was created in the CAD system SolidWorks. Before the computational grid generation of finite
elements in the Visual-Mesh module, the geometry was subjected to an analysis of the topology. The material properties of the
individual components of the ingot-casting system were defined with the Computherm calculating module selecting the
materials from its own database of ProCast. In addition, the thermodynamic properties were determined by using the datasheets
of the refractory materials of the manufacturer, and finally checked with the equations generally used to determine liquidus and
solidus temperatures, density and enthalpy, etc. The boundary conditions and the heat transfer were also defined. In parallel with
the numerical simulation, the operational experimental casting of a 90-ton ingot was carried out. To obtain more complete
information about the temperature fields of the ingot-casting system and of the data about the values of the heat flow, the
process of filling and solidification was monitored by using thermal imaging cameras. The conclusion summarizes the main
knowledge obtained on the basis of the primary results of the computation and gives a guideline for further research.
Keywords: numerical simulation, heavy-steel ingot, heat flow, thermal measuring
^lanek obravnava nove izku{nje z numeri~no simulacijo polnjenja in strjevanja 90-tonskega jeklenega ingota s
ProCAST-programom za simulacijo. Namen izvr{enega numeri~nega modeliranja v Department of Metallurgy in Regional
Materials Science and Technology Centre (RMSTC) na VSB-TU Ostrava je preverjanje in optimiranje proizvodne tehnologije
te`kih jeklenih ingotov, proizvedenih v Vítkovice Heavy Machinery, a. s. Vhodni parametri za izra~un so bili dolo~eni v realnih
razmerah pri ulivanju 90-tonskega jeklenega ingota. Geometrija ingota je bila postavljena v CAD-sistemu SolidWorks. Pred
postavitvijo mre`e kon~nih elementov z Visual-Mesh-modulom je bila geometrija obdelana s topolo{ko analizo. Lastnosti
materiala posameznih komponent livnega sistema ingota so bile dolo~ene z ra~unskim modulom Computherm in z izbiro
materiala iz datoteke ProCast. Termodinamske lastnosti so bile dolo~ene z uporabo datotek proizvajalcev ognjevzdr`nih
materialov ter kon~no preverjene z izra~unom po ena~bah, ki se uporabljajo za dolo~anje temperature likvidusa in solidusa,
gostote, entalpije itd. Dolo~eni so bili robni pogoji in prenos toplote. Vzporedno z numeri~no simulacijo je bilo izvr{eno
eksperimentalno ulivanje 90-tonskega ingota. Da bi dobili bolj popolno informacijo o temperaturnih poljih ulivnega sistema
ingota in podatke o vrednostih toplotnega toka, je bil postopek polnjenja in strjevanja posnet tudi s termovizijsko kamero. Sklepi
povzemajo glavne ugotovitve primarnih rezultatov izra~unov in dajejo napotke za nadaljnje raziskave.
Klju~ne besede: numeri~na simulacija, te`ki jekleni ingot, toplotni tok, termografske meritve
1 INTRODUCTION
The numerical modelling of filling and solidification
of a steel ingot was performed within the Department of
Metallurgy and RMSTC using the ProCAST software.
The software allows a 3D fully dimensional numerical
simulation of filling and solidification of steel including
a prediction of the ingot-volume defects such as porosity
and shrinkage. Due to the Flow and Stress modules, it is
possible to take into account the effects of natural
convection and of the formation of the air gaps between
the ingot body and the inner wall of the mould during the
calculation, and to predict the emergence of the internal
stress that can ultimately lead to cracks and rupture1.
In general, the numerical solution of each task is
divided into three stages:
1. Pre-processing: includes the geometric modelling,
the computational mesh-generation process, and the
definition of the calculation.
2. Processing: involves the computation in the solver.
3. Post-processing: focuses on the evaluation of the
results.
The conditions for the numerical model setting were
based on the real conditions of an experimentally cast,
Materiali in tehnologije / Materials and technology 46 (2012) 4, 399–402 399
UDK 621.74.043(437.3):519.872 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 46(4)399(2012)
90-ton steel ingot produced in VÍTKOVICE HEAVY
MACHINERY, a. s.
2 GEOMETRY AND THE COMPUTATIONAL
MESH
The computational-mesh geometry and the gene-
ration of the casting system for the 90-ton steel ingot
were made in cooperation with MECAS ESI, s. r. o. The
comparison between the real and the CAD geometry of
the casting system is shown in Figure 12. Figure 2
shows the final computation mesh of the casting system.
The volume ingot mesh consists of 348 794 nodes. The
total volume of the tetra elements is 1 766 041. The
mould has a rough mesh. For a more appropriate descrip-
tion of the geometry, the details, the insulation, and the
ceramic parts have a better mesh.
The final mesh of geometry was saved in *. mesh
format and subsequently loaded in the ProCAST module
for a calculation entry, also called the PreCAST. It was
necessary to define:
– the material properties of the individual parts of the
casting system,
– the heat-transfer coefficients on the interface of the
geometry elements,
– the boundary conditions, such as the casting tem-
perature, the casting speed, the conditions of heat
losses through the surface of the ingot,
– the operating conditions (such as gravity),
– the initial conditions for the calculation,
– the calculation parameters – the so-called RUN
PARAMETERS.
Among the other things, the RUN PARAMETERS
define the conditions for the calculation termination, the
so-called STOP criteria. The stop criteria also include
the attainment of a certain temperature in the ingot or the
termination of the calculation at a particular time after
the filling. The number of the steps of the calculation is
also specified, as well as the length of the time step and
the frequency of storing the results regarding the
temperature field and/or the heat flux2,3.
The computational time of one variant is around 96 h
when two processor cores are used. However, the time
needed for the preparation of the simulation and for the
evaluation of the achieved results has to be added to the
time taken by the computation.
3 MONITORING OF THE HEAT FLUX AND OF
THE TEMPERATURE DURING THE
EXPERIMENTAL CASTING
As mentioned above, the course of the changes in the
temperature field on the surface of the casting system
during the experimental casting of a 90-ton ingot was
also monitored by using the thermal-imaging cameras.
The thermography measurements were performed with
AVIO TVS 700 cameras. According to the tables of the
emission coefficients, the value of emissivity was set at
0.85. At the time of the measurement, an average
ambient temperature was 26 °C corresponding to the
conditions of the operation of a steel plant. The
measurement was not affected by the increased dusting
or drafts, and it was done with a tripod from a constant
distance of 6.5 m. Figure 3 shows the workers making
the thermography measurements.
During the thermography measurements images were
being taken in 10-minute intervals, from the moment
when the casting system got preheated, during the
casting itself and until after the stripping, therefore for
24 h. The temperatures were monitored on the surface of
the head, the mould, the washers, and the casting stake.
During the measurements, the temperature of the pit
wall, in which the casting system was located, was
scanned twice. After the completion of the filling, an
image of the temperature-field distribution on the surface
of steel, incurred after the addition of the insulation
backfill, was recorded. A self-evaluation of images was
M. TKADLECKOVA et al.: SETTING A NUMERICAL SIMULATION OF FILLING AND SOLIDIFICATION ...
400 Materiali in tehnologije / Materials and technology 46 (2012) 4, 399–402
Figure 1: Comparison between the real and the CAD geometry of the
casting system for a 90-ton ingot2
Slika 1: Primerjava realne in CAD-geometrije ulivnega sistema
90-tonskega ingota2
Figure 2: View of the final computational mesh of the casting system
Slika 2: Videz kon~ne ra~unske mre`e ulivnega sistema
performed in the SW GORATEC Thermography Studio
v4.5. When evaluating the images, the temperature
distribution, the temperature profile and the heat flux in
the selected area were monitored. The values of the
thermal flow and the temperature information were used
to specify the parameters of the setting of the numerical
simulation of filling and solidification of the steel ingot.
4 DISCUSSION OF THE ONGOING RESULTS
During the tuning of the setting of the numerical
simulation of the filling and solidification of a 90-ton
ingot, 9 configurations have been calculated up to now.
The difference between the temperature profiles on the
surface of the ingot obtained for the first variant and for
9th variant, 20 min after the initiation of the filling and at
the end of the filling, when the thermography-measure-
ment results were used, is captured in Figure 4.
The software for the image evaluation has different
options for setting the displayed scale of temperatures;
thus, the authors of the paper used the possibilities of the
Visual Viewer of ProCAST and adjusted the scale of the
numerical-simulation results to the scale of thermo-
graphy. The selection and editing of the colours were
dependent on the colour sense of the user. The advantage
of the Visual-Viewer postprocessor is the display of the
temperature in precise contours – changes to the mini-
mum and maximum values can provide accurate infor-
mation about the temperature in the chosen geometric
location. In its first column Figure 4 shows the tempe-
rature evolution on the mould surface for the original
setting, where the values of the heat-transfer coefficients
ranged from 100 W m–2 K–1 to 1000 W m–2 K–1 as is typi-
cal in these simulations. However, it is obvious that the
temperature values on the mould surface are below the
values measured during the thermography measurements
(the second column). Therefore, it was necessary to adjust
the coefficients according to the measured values
depending on time. The third column already captures the
evolution of the temperature on the mould surface after
the adjustment of the heat-transfer coefficients. The
M. TKADLECKOVA et al.: SETTING A NUMERICAL SIMULATION OF FILLING AND SOLIDIFICATION ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 399–402 401
Figure 4: Comparison of the temperatures on the ingot mould surface: a) 20 min after the initiation of the filling, b) after the filling
Slika 4: Primerjava temperature povr{ine kokile: a) 20 min po za~etku ulivanja, b) po koncu ulivanja
Figure 3: View of the workers making the thermography measure-
ments
Slika 3: Delavci med izvajanjem termografskih meritev
results of the thermography measurements and the nume-
rical simulation are already converging.
5 CONCLUSION
The setting of the numerical model for filling and
solidification of a heavy 90-ton steel ingot was based on
real experimental conditions of casting. The numerical
modelling was carried out under the conditions of the
Department of Metallurgy and RMSTC at VSB-TU
Ostrava in ProCAST SW. During the solution, an import
of the specific material properties was planned. For
example, solidus or liquidus temperatures were deter-
mined for the material (steel) using the instrumentation
for a thermal analysis (STA 449 Jupiter of NETZSCH
Company) placed in the laboratory of RMSTC4. During
the experimental ingot casting, the temperatures and the
heat flows on the surface of the casting system were
monitored with a thermal imaging camera. Once the
process of the temperature profile on the mould surface
of the numerical-simulation results corresponds to the
results of the thermography measurements from the start
of the filling to the actual stripping, the researchers will
proceed to further stages of numerical simulations based
on the distribution of the segregation of chemical
elements through the cross section of the ingot. At the
same time, the analysis of the segregation distribution
will be evaluated on the experimental cut ingot. The aim
of the numerical simulations will be to test the changes
in the boundary conditions of casting that should help us
not only to minimize the volume of defects, but also to
minimize the occurrence of the segregation in the heavy
forging of ingots produced by VÍTKOVICE HEAVY
MACHINERY, a. s.
Acknowledgments
This paper was created within the project TIP No.
FR-TI3/243, the Ministry of Industry and Trade of the
Czech Republic, and within the project No. CZ.1.05/
2.1.00/01.0040, "Regional Materials Science and Tech-
nology Centre" (ED0040/01/01), the operation programme
"Research and Development for Innovations" financed
by the Structural Funds and by the state budget of the
Czech Republic.
6 REFERENCES
1 M. Tkadle~ková, K. Michalek, P. Klus, K. Gryc, V. Sikora, Testing of
numerical model setting for simulation of steel ingot casting and
solidification. In 20th Anniversary International Conference of
Metallurgy and Materials Metal 2011. May 18th-20th, Brno, Czech
Republic, EU. ©Tanger, s. r. o. Ostrava. pp. 6
2 M. Tkadle~ková, K. Gryc, K. Michalek, P. Klus, L. Socha, P.
Machov~ák, M. Ková~, Verifikace vlivu okrajových parametrù lití
ocelových ingotù na velikost objemových vad. In XXI. International
Scientific Conference Iron and Steelmaking. 19.–21. øíjna 2011,
Horní Be~va, Beskydy, ^eská republika. s. 5
3 ProCAST 2009, Release Notes & Installation Guide. 2009 ESI
Group
4 K. Gryc, B. Smetana, K. Michalek, V. Sikora, M. Tkadle~ková, S.
Zlá, M. @aludová, Linking basic and apllied research in thermo
physical study of the properties of steel, slag and ferro-alloys. In 20th
Anniversary International Conference of Metallurgy and Materials
Metal 2011. May 18th–20th, Brno, Czech Republic, EU. ©Tanger, s. r.
o., Ostrava. pp. 6
M. TKADLECKOVA et al.: SETTING A NUMERICAL SIMULATION OF FILLING AND SOLIDIFICATION ...
402 Materiali in tehnologije / Materials and technology 46 (2012) 4, 399–402
K. GRYC et al.: A STUDY OF THE HIGH-TEMPERATURE INTERACTION BETWEEN SYNTHETIC SLAGS AND STEEL
A STUDY OF THE HIGH-TEMPERATURE INTERACTION
BETWEEN SYNTHETIC SLAGS AND STEEL
[TUDIJ VISOKOTEMPERATURNE INTERAKCIJE MED
SINTETI^NO @LINDRO IN JEKLOM
Karel Gryc1, Karel Stránský2, Karel Michalek1, Zdenìk Winkler3, Jan Morávka4,
Markéta Tkadle~ková1, Ladislav Socha1, Jiøí Ba`an1, Jana Dobrovská5, Simona Zlá5
1V[B – Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of Metallurgy and Foundry, Czech Re-
public
2Brno University of Technology, Faculty of Mechanical Engineering, Institute of Materials Science and Engineering, Czech Republic
3Military Technical Institute of Protection in Brno, Czech Republic
4MATERIÁLOVÝ A METALURGICKÝ VÝZKUM s. r. o., Czech Republic
5V[B – Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of Physical Chemistry and Theory of
Technological Processes, Czech Republic
karel.gryc@vsb.cz
Prejem rokopisa – received: 2011-10-27; sprejem za objavo – accepted for publication: 2012-02-15
The paper is devoted to selected aspects of the current issues of the high-temperature interaction of the synthetic
multicomponent oxidic systems and selected grades of steel. Analysing the consequences of the interaction between slag and
molten steel is an integral part of modern research that can be connected with subsequent, applied research in the field of
producing and refining steel. The interactions between the synthetic slag and the molten metal is realized after the processing of
the liquid steel in the primary metallurgical aggregates, i.e., at the beginning of refining processes of secondary metallurgy. The
addition of these synthetic systems (slag) significantly affects many technological processes and indirectly affects the final
quality of the cast steel. This influence can be seen on two levels: the metallurgical and the metallographic. Both of these levels
of evaluation complement each other.
For the evaluation of the metallurgical aspects of refining processes a number of thermodynamic relations were used,
supplemented by empirically established formulas and coefficients. Metallographic analyses utilize modern tools in the study of
the structure and the chemical composition of materials, i.e., from light microscopy to sophisticated systems of microanalysis of
elemental composition using a scanning electron microscope combined with an energy-dispersive X-ray micro-analyser. It is
obvious that this area deserves more focused research and more attention, particularly in the context of the on-going needs for
the identification and quantification of phenomena taking place in the following metallurgical innovations.
Keywords: synthetic slag, interaction with steel, saturation
^lanek obravnava izbrane vidike aktualnih vpra{anj o visokotemperaturni interakciji sinteti~nih ve~komponentnih oksidnih
sistemov v izbranih vrstah jekel. Analiziranje posledic interakcije med `lindro in talino je sestavni del potekajo~ih raziskav in se
lahko pove`e z uporabnimi raziskavami na podro~ju proizvodnje in rafinacije jekla. Do interakcije med sinteti~no `lindro in
staljeno kovino pride po izdelavi taline jekla v primarni metalur{ki pe~i, na za~etku postopka rafinacije v sekundarni
metalurgiji. Dodatek sinteti~ne `lindre mo~no vpliva na {tevilne tehnolo{ke procese in neposredno vpliva na kon~no kvaliteto
ulitega jekla. Ta u~inek se ka`e na dveh nivojih: na metalur{kem in metalografskem. Oba nivoja ocene sta med seboj
komplementarna.
Za oceno metalur{kih vidikov postopka rafinacije je bilo uporabljeno ve~ termodinamskih odvisnosti, dopolnjenih z empiri~no
dolo~enimi ena~bami in koeficienti. Metalografska analiza uporablja sodobna orodja za {tudij strukture in kemijske sestave
materiala – od svetlobne mikroskopije do zapletenih sistemov mikroanalize, elementne sestave z uporabo vrsti~nega
elektronskega mikroskopa v kombinaciji z energijsko disperzijsko rentgensko spektroskopijo. O~itno to podro~je zahteva bolj
usmerjene raziskave in ve~jo pozornost, posebno v kontekstu zahtev, identifikacije in kvantifikacije pojavov, ki so posledica
metalur{kih inovacij.
Klju~ne besede: sinteti~na `lindra, interakcija z jeklom, nasi~enje
1 INTRODUCTION
It is necessary to continuously optimize the produc-
tion of steel in steel mills. Such an optimization of the
metallurgical processes is often connected with opera-
tional challenges that could be solved by changes in the
slag regime. An integral part of the steel-making process
is the use of different types of synthetic slag. These, to-
gether with other slag materials, facilitate the successful
course of the metallurgical reactions that are necessary to
achieve the desired chemical and metallurgical purity of
the steels.
In addition to plant experiments, which are an
essential and major part of the slag-optimization regimes
and whose results are crucial for the innovations in the
final steel production process, there is the possibility of
studying the interaction of metal and oxidic systems,
e.g., under laboratory conditions.1–4
This paper focuses on a discussion of the results of
the oriented research that was realized within the grant
project ID No. 106/09/0969 with the financial support of
the (GACR) Czech Science Foundation. The presented
results are not directly tied to any specific operating
conditions in a steel plant.
The studied synthetic slags react in real conditions
with steel, especially together with pre-existing slag
and/or with currently added slag-forming materials. Al-
Materiali in tehnologije / Materials and technology 46 (2012) 4, 403–406 403
UDK 669.1 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 46(4)403(2012)
though this study of the high-temperature interactions
between synthetic oxidic melts and steels lies outside the
field of applied research, its results can extend current
knowledge about the mechanisms associated with these
systems of melts. The methods used could provide guid-
ance to address the challenges associated with solving
the practical aspects of steelmaking technology.
A series of interesting results was obtained in the
course of this project. That is why the focus of this paper
will be only on an analysis of the changes in the chemi-
cal composition of the studied synthetic oxidic mixtures,
depending on the chemical composition of steel, with
which these compounds reacted during the laboratory
heats.
2 EXPERIMENTAL CONDITIONS AND
CHARACTERISATION OF THE STUDIED
MELTS
Heats designed for the study of the interaction be-
tween liquid metal and synthetic oxidic melts were car-
ried out under laboratory conditions at the Department of
Metallurgy in an induction furnace (Figure 1) connected
to a GV 22 high-frequency generator.
Each 300 g sample of steel (Table 1) was melted in a
corundum crucible under the continuous maintenance of
a protective atmosphere (Ar, flow rate of 0.5 L min–1).
One of the following synthetic oxidic mixtures (Table 2)
was added (30 g) after reaching the temperature of
1600 °C. The melting process of oxidic mixtures
requires approximately 2.5 min. However, a stable tem-
perature (1600 °C) was kept for an additional 10 min.
Then, the generator was stopped. The solidified steel and
slag were removed from the crucible after cooling down.
Subsequently, samples were prepared to perform the ap-
propriate analyses.
Table 1: The content of the monitored elements in the selected steels
Tabela 1: Vsebnost analiziranih elementov v izbranih jeklih
Identification
of steel
Chemical composition in mass fractions, w/%
C Mn Si P S Cr
1 0.19 1.16 0.14 0.015 0.033 1.100
2 0.21 1.04 0.13 0.033 0.170 0.054
3 0.12 1.20 <0.01 0.060 0.340 0.080
4 0.19 1.34 0.18 0.015 0.015 0.055
Table 2: Content of selected oxidic mixture components
Tabela 2: Vsebnost komponent v izbranih me{anicah oksidov
Identifi-
cation of
oxidic
mixture
Chemical composition in mass fractions, w/%
Fetot CaO Al2O3 SiO2 S P2O5 MgO MnO Cr2O3
A 0.42 83 11 1 0.05 0.44 1.91 0.81 0.49
B 0.79 48 46 2 0.12 0.39 1.08 0.2 0.41
C 1.96 34 47 9 0.19 0.55 1.58 0.23 0.31
D 1.21 49 39 5 0.32 0.39 1.68 0.07 –
E 0.63 31 56 5 0.67 0.12 3.84 0.02 0.02
It is obvious (Table 1) that the steels were selected
based on such a chemical composition in order to assess
the impact of significant changes in the content of the
significant elements on the result of the metallurgical
refining process.
3 DISCUSSION
The contents of the elements in the steel and slag
components listed in this paper in Table 1 and 2 were
determined using standard methods.
K. GRYC et al.: A STUDY OF THE HIGH-TEMPERATURE INTERACTION BETWEEN SYNTHETIC SLAGS AND STEEL
404 Materiali in tehnologije / Materials and technology 46 (2012) 4, 403–406
Figure 2: Ternary diagram for the starting oxidic mixture (A) and the
slag obtained after the reaction with molten steels
Slika 2: Ternarni diagram za~etne me{anice oksidov (A) in `lindra, ki
je nastala po interakciji s staljenimi jekli
Figure 1: Diagram of the experimental equipment; 1 – protective
cylinder of SiO2, 2 – water-cooled inductor, 3 – graphite block, 4 –
working crucible, 5 – protective Al2O3 powder, 6 – PtRh6 – PtRh30
thermocouple, 7 – molten metal, 8 – protective cover, 9 – supply of
argon, 10 – TERMOVIT cotton, 11 – insulating brick, 12 – fireclay
base
Slika 1: Shematski prikaz eksperimentalne opreme: 1 – za{~itna
obloga iz SiO2, 2 – vodno hlajen induktor, 3 – grafitni blok, 4 – talini
lonec, 5 – za{~itni Al2O3 pra{ek, 6 – PtRh6 – PtRh30 termo~len, 7 –
staljena kovina, 8 – za{~itni pokrov, 9 – dovod argona, 10 –
TERMOVIT preja, 11 – izolacijska opeka, 12 – ognjevarna obloga
Analyses of the elements in the steel:
• C, S using a CS 230 LECO combustion analyser,
• Mn, Si, P, Cr using a PW 1400 Philips X-ray spec-
trometer.
The analysis of the content of the monitored slag
components was implemented using the analytic com-
plex formed by the EDAX PHILIPS energy-dispersive
micro-analyser in conjunction with the PHILIPS scan-
ning electron microscope. A detailed description of the
apparatus and the applied methods can be, together with
the results of the analysis, found in the report5.
The results of the analyses are summarized in the
simplified CaO-Al2O3-SiO2 ternary diagrams (Figures 2
to 6). It is obvious that in terms of securing good kinetic
conditions of the refining processes (low viscosity of the
oxidic systems), it is advantageous to have a chemical
composition of oxidic systems with a melting point
below the working temperature of steel during its
processing in a given technology node.
The ternary diagrams (Figures 2 to 6) show the
chemical composition of the initial oxidic mixture
(green) and the slag after laboratory heats with different
steels (1 – dark blue, 2 – red, 3 – light blue, 4 – black).
Figure 2 shows a ternary diagram with plotted areas
for the oxidic mixture A (83 % CaO-11 % Al2O3-1 %
SiO2) and the slag after the reaction with molten steels 1
to 3. It is obvious that there is a significant change in the
chemical composition due to the interaction of an oxidic
mixture with steels. The resulting oxidic melt contains
mass fractions of Al2O3 from 51 % to 60 %. The chemi-
cal composition of the resulting slag, in the case of an in-
teraction with the steels 1 and 3, ensures that the slag
was liquid at temperatures above 1600 °C. In contrast,
the resulting slag for the interaction with steel No. 2 oc-
curs in an area that does not guarantee its liquid state.
The initial, synthetic slag B (48 % CaO-46 % Al2O3-
2 % SiO2) occurs in the ternary diagram area that is
characterized by melting temperatures in the interval
K. GRYC et al.: A STUDY OF THE HIGH-TEMPERATURE INTERACTION BETWEEN SYNTHETIC SLAGS AND STEEL
Materiali in tehnologije / Materials and technology 46 (2012) 4, 403–406 405
Figure 6: Ternary diagram for the starting oxidic mixture (E) and the
slag obtained after the reaction with molten steels
Slika 6: Ternarni diagram za~etne me{anice oksidov (E) in `lindra, ki
je nastala po interakciji s staljenimi jekli
Figure 4: Ternary diagram for the starting oxidic mixture (C) and the
slag obtained after the reaction with molten steels
Slika 4: Ternarni diagram za~etne me{anice oksidov (C) in `lindra, ki
je nastala po interakciji s staljenimi jekli
Figure 3: Ternary diagram for the starting oxidic mixture (B) and the
slag obtained after the reaction with molten steels
Slika 3: Ternarni diagram za~etne me{anice oksidov (B) in `lindra, ki
je nastala po interakciji s staljenimi jekli
Figure 5: Ternary diagram for the starting oxidic mixture (D) and the
slag obtained after the reaction with molten steels
Slika 5: Ternarni diagram za~etne me{anice oksidov (D) in `lindra, ki
je nastala po interakciji s staljenimi jekli
1300 °C to 1400 °C (Figure 3). The chemical compo-
sition of the reaction products, is after the interaction of
the (B) oxidic mixture with steels No. 1 to 3, close to the
original one. However, due to the increasing Al2O3 con-
tent, the steels have increased their melting temperature,
to the 1400 °C to 1600 °C interval, which still guarantees
their high reactivity in order to maintain their relatively
low viscosity.
The initial oxidic mixture C (34 % CaO-47 %
Al2O3-9 % SiO2 – Figure 4) contains the most silica
compared with the previous slag (Figures 2 and 3). The
CaO-Al2O3-SiO2 final slag remains in the ternary
diagram, mostly in the same area, which is characterised
by the melting-temperature interval from 1400 °C to
1600 °C after the reaction with steels 1 to 3. The melting
temperature exceeds 1600 °C only partially.
The initial chemical composition of the oxidic mix-
ture D (49 % CaO-39 % Al2O3-5 % SiO2 – Figure 5) is
very close to the chemical composition of the mixture B
(Figure 3). Its position in the ternary diagram is also lo-
cated in areas with a relatively low melting point
(1300–1400 °C). The oxidic interaction products have,
compared to the oxidic mixture B (w = 2–4 %) a higher
content of silica (w = 5–7 %).
The oxidic mixture E (31 % CaO-56 % Al2O3-5 %
SiO2 – Figure 6) with its initial chemical composition,
differs from the previous ones (B, C, D) in terms of a
higher content of Al2O3 to the extent that it already ap-
pears in the area with a relatively high melting point
(above 1600 °C) in the ternary diagram. The products of
the mixture interaction with the metal melt no longer
contain the higher content of Al2O3, and are located at
the identical area in the ternary diagram.
4 CONCLUSION
The analysis of the interaction of synthetic oxidic mix-
tures with steels having different contents of the major met-
allurgical elements C, Mn, Si, P, S, Cr resulted in the fol-
lowing conclusions.
The interaction of the above-discussed oxidic sys-
tems and selected steels led to a final content of Al2O3 in
the interval (w = 46–60 %). Although the content of
Al2O3 in the initial oxidic mixture (A) was the lowest (11
wt.%), in comparison with the other mixtures the Al2O3
content was the highest of all (w = 60 %) after its inter-
action with the steel (2). These facts lead to the hypothe-
sis that the settings of the experiments enabled the satu-
ration of the resulting slag with Al2O3 content in the
interval 46 % to 60 %. This saturation occurs within 10
min of the dissolution of the added oxidic mixtures. Fur-
thermore, we can say that a clear dependence of the re-
sulting chemical composition of the oxidic mixture on
steels, with which they interacted, has not been proved.
The content of SiO2 did not increase above 9 % in all the
studied cases and its influence on the change of the melt-
ing temperature of the oxidic mixtures was not signifi-
cant.
Acknowledgement
The presented research focused on the high-tempera-
ture interactions between synthetic slag and steel was
carried out with financial support from GACR (Czech
Science Foundation) under project ID No. 106/09/0969.
5 REFERENCES
1 B. Smetana, S. Zlá, M. @aludová, J. Dobrovská, P. Kozelský,
Application of High Temperature DTA to Micro-Alloyed Steels,
Metalurgija, 51 (2012) 1, 121–124
2 B. Smetana, S. Zlá, J. Dobrovská, P. Kozelský, Phase transformation
temperatures of pure iron and low alloyed steels in the low
temperature region using DTA, International Journal of Materials
Research, 101 (2010) 3, 398–408
3 R. Dudek, ¼. Dobrovský, J. Dobrovská, Interpretation of Inorganic
Melts Surface Properties on The Basis of Chemical Status and
Structural Relations, International Journal of Materials Research, 99
(2008) 12, 1369–1374
4 K. Gryc, B. Smetana, K. Michalek, V. Sikora, M. Tkadle~ková, S.
Zlá, M. @aludová, J. Dobrovská, Connecting of Basic and Applied
Research in the Field of Thermo-Physical Study of the Properties of
Steels, Slag and Ferro-Alloys, In 20th Anniversary International
Conference on Metallurgy and Materials – METAL 2011, 125–131
5 K. Stránský, J. Janová, Chemical analysis of slags realized in
laboratory of Brno University of Technology, FME in 2010.
Technical report. BUT, FME, Brno 2010. [in Czech]
K. GRYC et al.: A STUDY OF THE HIGH-TEMPERATURE INTERACTION BETWEEN SYNTHETIC SLAGS AND STEEL
406 Materiali in tehnologije / Materials and technology 46 (2012) 4, 403–406
L. GLOBA et al.: DEVELOPMENT OF A MODEL FOR THE INTERNET PORTAL "STRENGTH OF MATERIALS"
DEVELOPMENT OF A MODEL FOR THE INTERNET
PORTAL "STRENGTH OF MATERIALS"
RAZVOJ MODELA ZA INTERNETNI PORTAL "TRDNOST
MATERIALOV"
Larisa Globa1, Rina Novogrudska1, Ilija Mamuzi}2
1National Technical University of Ukraine "Kyiv Polytechnic Institute", Institute of Telecommunication Systems,
37 Pobeda ave, 03056 Kyiv, Ukraine
2University of Zagreb, Faculty of Metallurgy, Berislavi}eva 6, 10 000 Zagreb, Croatia
lgloba@its.kpi.ua
Prejem rokopisa – received: 2011-11-09; sprejem za objavo – accepted for publication: 2012-03-01
We present an approach to the development of a specialized-knowledge Internet portal for work with large quantities of
information and computational resources in the field of strength of materials. The ontology-based portal provides an information
basis for the design of alloys represented by the chosen fields of science, which may be difficult for formalizing, and it allows
the use of web services for the realization of engineering tasks with the Internet and material-science knowledge that is
accumulated in the databases.
Keywords: Internet portal, knowledge representation, model, ontology, strength of materials
Opisan je razvoj specializiranega znanja za internetni portal za uporabo velike koli~ine informacijskih in podatkovnih virov s
podro~ja trdnosti materialov. Ontolo{ka podlaga portala zagotavlja informacijsko osnovo za na~rtovanje zlitin, ki obsegajo
podro~ja znanosti, ki jih je te`ko formalizirati, in omogo~a uporabo web-storitev za realizacijo in`enirskih nalog z uporabo
interneta in znanosti o materialih, zbranih v bazah podatkov.
Klju~ne besede: internetni portal, predstavitev znanja, model, ontologija, trdnost materialov
1 INTRODUCTION
Nowadays, there is a wide variety of technical
resources and software solutions that can be used to
solve various engineering tasks. However, their use is
sometimes restricted for two reasons: the software and
technical resources of this type are either very expensive
or kept hidden for commercial purposes. Also, the new
theoretical and practical results obtained by researchers
in numerous institutions may be concentrated in
institutions, meaning their external use is limited. Thus,
providing access to the information for as many users as
possible is actually becoming a real challenge.
On the other hand, large quantities of information are
already stored on the Internet, but it may be poorly
structured and systematized and distributed across
different sites, electronic libraries and archives. This may
prevent the rapid and easy access to specific knowledge.
To solve these problems a specialized-knowledge
Internet portal for work with a large quantity of infor-
mation and computational resources in a defined
technical sphere is proposed. Such a portal cannot only
provide the possibility to search and systematize the
information, but it can also help to realize specific
computational tasks for the users.
2 KNOWLEDGE REPRESENTATION IN THE
FIELD OF THE STRENGTH OF MATERIALS
When we introduce a formal description of the
subject field in the form of object classes and their
mutual relations, the portal’s ontology gives the
structures that present real data and their inter-
connections. The use of an ontology to construct the
informational basis of the portal gives an integral
presentation of the technical fields that are considered to
be difficult to formalize, and also to allow the auto-
mation of the processes of acquiring information and its
storage on any chosen field. Such a conceptual model
makes it possible to uniformly present the knowledge
data and the semantic coherence.
The knowledge-portal ontology in the field of the
strength of materials was constructed on the basis of the
above descriptions. Formally, the ontology may be
specified as O = {C, A, R, T, F, D}. Here, C is the set of
classes that describes the notions of a subject field; A is
the set of attributes that describes the features of the
notions and relations; R is the set of relations specified
for the classes R R R R RAS IA N CD= { }, , , with RAS – associ-
ative relation, RIA – relation "is-are", RN – relation of
"heredity", RCD – relation "class-data"; T – the set of
standard types of attribute values; F – set of limitations
for values of attribute notions and relations; and D is the
set of class exemplar1,2.
Materiali in tehnologije / Materials and technology 46 (2012) 4, 407–410 407
UDK 004.738.5:539.4 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 46(4)407(2012)
Such an ontology may serve to present the notions
that are necessary for describing the knowledge in the
field of the strength of materials as well as for the
engineering activities performed in this context.
3 ONTOLOGY OF THE PORTAL "STRENGTH
OF MATERIALS"
The ontology of the portal in the field of the strength
of materials includes four ontologies: engineering-
activity ontology, engineering-knowledge ontology,
engineering-computations ontology and subject-field
ontology (Figure 1).
The engineering-activity ontology (EAO) includes
the general classes of notions related to the organization
of scientific activities. The engineering-knowledge
ontology (EKO) includes the meta-notions that specify
the structures to describe the problem field, the engi-
neering-computations ontology (ECO) groups classes
that describe the portal’s calculation abilities and the
subject-field ontology (SFO) represents the general
knowledge of the subject field, such as the hierarchy of
notions classes and their semantic relations.
In Figure 2 the EAO, EKO and ECO classes, as well
as their specified relations, are shown. The classes are
mapped as an oval line forms in the definite rectangle
corresponding to the ontology the class belongs to. For
example, the classes "Person", "Organization", "Acti-
vity" and etc. are elements of the EAO and are placed in
the first rectangle. The classes "Research method",
"Research object", "Research result" etc. belong to EKO
and are located in the second rectangle, while the classes
"Calculation", "Service", "Parameters" and "Result" are
part of the ECO and are located in the third rectangle.
The classes enumerated are related to each other in
single ontology and to classes of other ontologies by
associative relations. For example, the classes of the
EAO "Person" and "Organization" are related through
the associative relations "Be a member of". It means that
in real life a person may be a member of some organi-
zation. Associative relations may correlate the classes of
a single ontology and the classes that belong to a diffe-
rent ontology, also. For example, the class "Literature"
being a class of the EAO is associatively correlated by
the relation "Describe" with the ECO class "Research
result". The associative relations allow the understanding
of the correlation of notions that are described in one
class of ontology with another class notion in reality. In
addition to the associative relations, in working up the
portal ontologies the relations of the type "is-are" to
relations of subclasses with their parent classes are used.
For example, the class "Literature" is related to "is-are",
to the classes "Documents", "Training materials", and
"Published materials". It means that the class "Litera-
ture" is a parent class for its subclasses "Documents",
L. GLOBA et al.: DEVELOPMENT OF A MODEL FOR THE INTERNET PORTAL "STRENGTH OF MATERIALS"
408 Materiali in tehnologije / Materials and technology 46 (2012) 4, 407–410
Figure 2: Elements of the portal ontology
Slika 2: Elementi ontologije portala
Figure 1: Portal ontology model
Slika 1: Portal ontologije modela
"Training materials" and "Published materials". Such
types of relations play an important role in presenting the
hierarchical structure of a model of engineering know-
ledge.
4 ONTOLOGY OF THE SUBJECT FIELD
The ontology of the subject field describes the
strength of materials as a whole as science and its parts,
notions and their connections. These notions are the
realizations of meta-notions of the EKO and may be put
in the order into the hierarchy "is-are". For example,
"Research methods" (class of the EKO) correspond to
such methods as the methods of strain, the discharge
method, the stress distribution method, etc. in the
strength of materials3,4. The "Research objects" are
materials, material groups or specific material properties
(Figure 3). The main class of the ECO "Calculation"
corresponds to such notions from the field of strength of
materials as the limit state design, deformation analysis,
stress calculation, etc.
The classification of materials was carried out on the
basis of the classification system (a hierarchy that was
constructed according to the properties of materials and
their features) and is shown below5:
Level 1 – types of materials:
Steel, aluminum alloys, titanium alloys, copper, etc.
Level 2 – purpose: (if above steel was chosen)
Structural steel, tool steel, cast steel.
Level 3 – composition (if above steel –> steel structural
were selected)
carbon steel, carbon steel of high quality, low-alloy
steel, alloy steel, etc.
Level 4 – description according to the:
Bars, plates and sheets, mass fraction of elements,
etc.
Level 5 – general properties:
Mass content of elements, temperature of critical
points, assignment, etc.
Level 6 – mechanical properties.
Level 7- technological characteristics,
5 OPTIONS OF THE PORTAL
The portal information is systematized in the
following areas:
• investigation of the mechanical properties of
materials, the construction components and the
structures:
– types of research,
– methodologies,
– techniques and research tools,
– processing of experiments results and their
statistical analysis,
– software,
– certification of testing equipment and research
laboratories,
– regulation in the field,
– etc.
• mechanical properties of materials:
– static loading,
– sustained static loading,
– dynamic loading,
– cyclic loading,
– combined loading,
– etc.
• strength calculation
– calculation methods,
– evaluation of stress-strain state of structures,
– software resources,
– specifications,
– etc.
The proposed portal will provide access to databases,
reference books, manuals, express information, network
resources, etc. for the browsing of different types of
theoretical information and practical results collected
and stored for years in different institutions. By means of
this portal it will be possible to solve different calcu-
lation and computational tasks according to the user’s
needs. For example, the portal will make it possible: to
L. GLOBA et al.: DEVELOPMENT OF A MODEL FOR THE INTERNET PORTAL "STRENGTH OF MATERIALS"
Materiali in tehnologije / Materials and technology 46 (2012) 4, 407–410 409
Figure 4: Portal options
Slika 4: Mo`nosti portala
Figure 3: Ontology of the subject field
Slika 3: Ontologija polj subjektov
find a material according to some criteria (characte-
ristics), to plot different graphs, diagrams and depen-
dences, to compute the data of various types according to
the user’s needs. The portal will make it possible to
browse blocks of news, information about conferences
(upcoming and past), competitions and grants, as well as
other information about various events relevant to this
area of knowledge (Figure 4). An important step in the
portal’s construction is the structuring and systemati-
zation of the information and knowledge in the field of
strength of materials that will allow users to browse and
search specific information in the area chosen. The
placement of information on the portal is organized
comfortably for end-user implementing problem-
oriented navigation and search tools. Thus the search for
information is organized to give the user the possibility
to specify a search request, not only using keywords, but
using well known terms of the portal subject field as
well.
6 CONCLUSIONS
The possibility of a "strength of materials" portal
design is described. The portal’s informational model is
presented by means of an ontology that allows us to
systematize and structure the information and to organize
an effective search and navigation through the infor-
mation space of the engineering-knowledge portal.
7 REFERENCES
1 O. A. Andreeva, O. I. Borovikova, Y. A. Zagogulko et al.: Archeo-
logical portal of knowledge: substantial access to knowledge and
informational resources / 1st national conference of artificial
intelligence KII 2006. – M.: PhithMathLit., 2006, 832–840
2 Y. A. Zagogulko, O. I. Borovikova, The technology of ontologies
constructing for the portals of scientific knowledge, Journal of NGU.
Series: Informational technologies, 5 (2007) 2, 12–15
3 Ya. B. Fridman, Mechanical properties of metals – M.: Mechanical
engineering, 1974, 2p, 368
4 Strength of materials and constructions / editorial board: V. T. Tro-
shenko (editor-in-chief) and others. – K.: Akademperiodica, 2005,
1088p
5 A. C. Zubchenko: Book of steels and alloys. – M. - Mechanical
engineering, 2001, 663p
L. GLOBA et al.: DEVELOPMENT OF A MODEL FOR THE INTERNET PORTAL "STRENGTH OF MATERIALS"
410 Materiali in tehnologije / Materials and technology 46 (2012) 4, 407–410
A. MILOSAVLJEVIC et al.: THE INFLUENCE OF THE HEAT-TREATMENT REGIME ON A FRACTURE SURFACE ...
THE INFLUENCE OF THE HEAT-TREATMENT REGIME
ON A FRACTURE SURFACE OF NICKEL-BASED
SUPPERALLOYS
VPLIV TOPLOTNE OBDELAVE NA POVR[INO PRELOMA
SUPERZLITIN NA OSNOVI NIKLJA
Andjelka Milosavljevic1, Sanja Petronic2, Suzana Polic-Radovanovic3,
Jasmina Babic4, Darko Bajic5
1Faculty of Mechanical Engineering, Belgrade, Serbia
2Innovation Center of the Faculty of Mechanical Engineering, Belgrade, Serbia
3Central Institute for Conservation, Belgrade, Serbia
4Military Technical Institute, Belgrade, Serbia
5Faculty of Mechanical Engineering, Podgorica, Montenegro
sanjapetronic@yahoo.com
Prejem rokopisa – received: 2012-01-11; sprejem za objavo – accepted for publication: 2012-02-23
Nickel-based superalloys are distinguished from other materials by their excellent mechanical and physical properties. As they
are used at high temperatures and pressures, as well as in aggressive environments, their characteristics need constant
improvement. An adequate choice of their chemical composition and the heat-treatment regime contributes to the improvement
of the chemical, physical and mechanical properties of these nickel-based superalloy materials. During the heat treatments of the
superalloys Nimonic 263 and Hastelloy S some changes in their microstructures were observed. In this paper the changes in the
microstructures after various regimes of the heat treatment were analysed on the fractured surfaces. The fractured surfaces were
observed using light microscopy, scanning electron microscopy (SEM), and energy-dispersive X-ray spectrometry (EDS).
Keywords: superalloy, heat treatment, microstructure, EDS, SEM
Superzlitine na osnovi niklja se razlikujejo od drugih materialov zaradi svojih odli~nih mehanskih in fizikalnih lastnosti. Ker se
jih uporablja pri visokih temperaturah in tlakih ter v agresivnem okolju, je potrebno stalno izbolj{evati njihove lastnosti.
Primerna izbira kemijske sestave in na~ina toplotne obdelave prispevata k izbolj{anju kemijskih, fizikalnih in mehanskih
lastnosti superzlitine na osnovi niklja. Med toplotno obdelavo superzlitine Nimonic 263 in Hastelloy S se opazijo spremembe v
mikrostrukturi. V tem ~lanku so analizirane spremembe mikrostrukture na prelomnih povr{inah po razli~nih toplotnih
obdelavah. Povr{ina prelomov je bila opazovana s svetlobno mikroskopijo, vrsti~no elektronsko mikroskopijo (SEM) in
rentgensko disperzijsko spektroskopijo (EDS).
Klju~ne besede: superzlitina, toplotna obdelava, mikrostruktura, EDS, SEM
1 INTRODUCTION
Modern industry has a high demand for superalloys
due to their enhanced technological features, such as
good mechanical strength and hardness, corrosion
resistance, heat resistance, wear resistance and surface
degradation1–3.
The good tensile strength of superalloys is based on
the principle of a stable face-centered cubic matrix com-
bined with precipitation hardening and/or solid-solution
strengthening4.
Precipitation hardening produces a high strength with
finely dispersed precipitates formed during the heat
treatment and deposited in the elastic matrix. These
particles can be obstacles to the movement of dislo-
cations through the crystal structure, thus reinforcing the
heat-treated alloy.
The effect of the superalloys’ strengthening depends
on the type of particles. The best results can be achieved
when the coherent and partially coherent particles finely
disperse within the matrix.
Precipitation hardening increases the mechanical
properties, especially the strength of the materials, with
the precipitation from supersaturated solid solutions. The
principal strengtheners in the nickel-based superalloys
are complex precipitates of ’Ni3(Al, Ti) and ”
Ni3(Nb, Al, Ti) and the carbide particles. Other phases
have a negligible effect on increasing the tensile strength,
but a significant effect on increasing the creep and
fracture strength, and the segregation5.
Some unwanted phases can precipitate during the
heat treatment, plastic deformation and/or long-time
service. By selecting the optimal heat-treatment time and
the optimal heat-treatment temperature, these phases
could be avoided5,6.
In addition, a proper selection of the chemical
composition and the heat treatment contributes to an
improvement of the chemical, physical and mechanical
properties of the nickel-based superalloys7.
During the heat treatment of the superalloys Nimonic
263 and Hastelloy S the phase transformations occurred8.
The influence of the heat treatment on the micro-
structure of various metals and their alloys has been
Materiali in tehnologije / Materials and technology 46 (2012) 4, 411–417 411
UDK 669.245:621.785 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 46(4)411(2012)
investigated so far5,9–12, but, bearing in mind the import-
ance and applicability of these superalloys in industry, it
is undoubtedly important to further investigate the
influence of heat treatment on the microstructure of these
superalloys.
In this paper, the influence of heat treatment on the
microstructure of two nickel-based superalloys, which
have been strengthened with different strengthening
mechanisms, is investigated. Hastelloy S is a solid-solu-
tion strengthened alloy with a high content of molybde-
num. Nimonic 263 is a precipitation-hardenable alloy
with an addition of molybdenum for the solid-solution
strengthening. These two superalloys are very often used
in domestic industry.
2 EXPERIMENTAL
In this paper, the experimental investigations are
carried out on the commercial, nickel-based superalloys
Nimonic 263 and Hastelloy S. The samples are cut from
the sheets, thickness of 1.2 mm. The chemical compo-
sition is determined by the gravimetric method and listed
in Table 1.
The homogenization heat treatment of the superalloys
Nimonic 263 and Hastelloy S is performed in a vacuum
at a temperature of 1050 °C and for 16 h, with the aim to
achieve a homogeneous structure.
The samples of Nimonic 263 superalloy are subjected
to a two-stage heat treatment:
1) solid solution at 1150 °C and cooling in water,
2) aging at 800 °C/8 h and cooling the air.
The solid solution time at a temperature of 1150 °C
for one group of Nimonic 263 samples is 10 min (regime
– R1). As the results obtained after regime R1 suggest
that the solid-solution time is short, the second group of
samples is subjected to a solid solution treatment for 60
min (regime – R2), due to the assumption that longer
periods of solid solution treatment could bring about the
formation of coarse particles. The regime of precipitation
hardening for both groups is the same – 800 °C/8 h
cooled air.
The heat treatment of the superalloy Hastelloy S,
described according to literature data and experience5,13:
• Solution treatment at 1080 °C/1 h and rapidly cooled
in water to room temperature,
• Primary precipitation hardening at 840 °C/4 h and
cooling to room temperature,
• Secondary precipitation hardening at 760 °C/3 h and
cooling in air to room temperature.
After the heat treatment carried out according to the
described regime, the superalloy samples were prepared
for light microscopy. After the polishing of the samples
the etching was performed for the superalloys Nimonic
263 and Hastelloy S in a solution with the following
composition: 30 % H2O, 20 % HNO3, 10 % HF, 20 %
H3PO4 and 10 % CH3COOH.
The fractured surfaces were observed with a light
microscope (model KEYENCE VH-Z100), a scanning
electron microscope (model JOEL JSM-5800) and
analysed using energy-dispersive X-ray spectrometry.
The tensile tests were performed with a mechanical
universal testing machine (Schenck-Trebel, RM400) and
the 0.2 % offset yield strength, tensile strength and elon-
gation were determined. The tensile tests were carried
out at room temperature according to the en 10002-1
standard. The hardness tests were made with a semi-
automatic Hauser 249A and the HV30 was measured.
3 RESULTS AND DISCUSSION
The first step in the heat treatment of superalloys is
usually a solid-solution treatment. The solid-solution
temperature depends on the required characteristics.
Higher temperatures are used for optimum creep-
fracture properties, and produce a higher yield and a
more extensive carbide dissolution. Lower temperatures
result in the optimal yield strength at elevated tempera-
ture and a resistance to fatigue14.
By observing the microstructure of the Nimonic 263
samples, prepared for light microscopy by polishing and
etching, at 225-times magnification, there is a difference
in grain size, i.e., a finer grain size is achieved for the
heat treatment regime R2 (Figure 1b) than for the
heat-treatment regime R1 (Figure 1a). The average grain
size is calculated using the method of the circle15 and
their values are as following: Fm = 314.78 μm2 and Fm =
237.68 μm2 for the regimes R1 and R2 applied,
respectively.
The values measured by Vickers hardness with a load
of 30 N are:
• HV30 = 283 regime R1,
• HV30 = 302 for regime R2.
The finer grain structure and the higher hardness
values imply that the heat treatment carried out by
regime R2 has a more favourable impact on the pro-
perties of the superalloy Nimonic 263, compared to the
heat treatment by regime R1.
After the heat treatment, regime R2, which was
preceded by homogenization, the microstructure of the
A. MILOSAVLJEVIC et al.: THE INFLUENCE OF THE HEAT-TREATMENT REGIME ON A FRACTURE SURFACE ...
412 Materiali in tehnologije / Materials and technology 46 (2012) 4, 411–417
Table 1: Chemical composition of Nimonic 263 and Hastelloy S superalloys in mass fractions (w/%)
Tabela 1: Kemijska sestava superzlitin Nimonic 263 in Hastelloy S v masnih dele`ih (w/%)
Element C Si Mn Al Co Cr Cu Fe Mo Ti Ni
Nimonic 263 0.06 0.3 0.5 0.5 20 20 0.1 0.5 5.9 2.2 49.94
Hastelloy S 0.3 0.5 0.5 0.3 – 15.3 – 1.34 14.4 – 67
superalloy Nimonic 263 consists of the following:  solid
solution, ’ intermetallic compounds – ’Ni3(Al, Ti),
carbides M23C6, MC carbide and a number of annealing
twins. The volume fraction of ’ for the start of the
thermal deposition is about 10 % 16, and its fraction
increases with the deposition time, and is unevenly
distributed in the  solid solution. The M23C6 carbides are
densely distributed at the grain boundaries, while the
amount of MC carbide is small and mainly formed
during solidification and precipitation. The MC carbides
are mainly Ti carbides, and Ti is a ’ former as well. This
explains the creation of ’-free zones at the grain
boundaries, as Ti forms carbides, and depletes Ti in the 
solid solution near the grain boundaries.
Figure 2a shows a micrograph of the microstructure
of superalloy Nimonic 263, after heat treatment – the
regime R2. The results of the EDS analysis are listed in
Table 2. It is clear that the Cr and Ti carbides precipitate
at the grain boundaries, which confirms the increased
content of Cr in spectrum 1 and the increased content of
Ti in spectrum 2. An analysis of the grain (spectrums 3,
4, 5) indicates that the chemical composition is generally
close to the average one. For chromium carbides we
could not exactly determine the type of carbide. Accord-
ing to the literature3,5,17,18 they could be the M23C6 type
carbide, which are densely distributed at the grain boun-
daries. The size of these carbides is up to 600 nm. Their
shape is elliptical, and they are uniformly distributed at
the grain boundaries. The size, morphology, distribution
and location of these carbides are favourable for the
straightening of superalloy. Comparing to16 where time
of the solid solution was shorter, these carbides are finer
and more uniformly distributed at the grain boundaries.
A. MILOSAVLJEVIC et al.: THE INFLUENCE OF THE HEAT-TREATMENT REGIME ON A FRACTURE SURFACE ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 411–417 413
Table 2: Results of EDS-analysis of the spectrums in Figure 2a (w/%)
Tabela 2: Rezultati EDS-analize spektrov s slike 2a (w/%)
Spectrum Al Si Ti Cr Mn Fe Co Ni Mo
Spectrum 1 0.2 0.11 1.99 32.42 0.32 0.25 13.87 44.82 6.02
Spectrum 2 0.22 0.21 6.73 19.35 0.34 0.26 17.82 49.25 5.82
Spectrum 3 0.51 0.28 2.01 19.89 0.5 0.48 20 50.4 5.93
Spectrum 4 0.47 0.32 2.2 20.03 0.46 0.52 20.32 49.85 5.83
Spectrum 5 0.54 0.27 2.18 19.95 0.53 0.51 19.98 50.14 5.9
Figure 2: a) Carbides precipitated at the grain boundaries after the
heat treatment regime R2, b) TiC carbide precipitated by solution
treatment of regime R1 (SEM)
Slika 2: a) Karbidni izlo~ki po mejah zrn po toplotni obdelavi z
na~inom R2, b) TiC karbidni izlo~ki po raztopnem `arjenju po na~inu
R1 (SEM)
Figure 1: The microstructure of the superalloy Nimonic 263 after heat
treatment by: a) regime R1, b) regime R2, taken with an optical
microscope
Slika 1: Mikrostruktura superzlitine Nimonic 263 po toplotni obde-
lavi: a) na~in R1, b) na~in R2, opti~ni mikroskop
According to Figure 2b and Table 3, it can be con-
cluded that in spectrum 1, at the grain boundaries, the
carbide TiC precipitated. The grain boundaries are
’-free, and the EDS analysis in spectrum 3 and Figure
2b indicates that Ti carbides are deposited at the grains.
It is believed that this is the reason for the relatively
small amount of ’ phase. The carbides at the grains are
fine and considered as being favourable for the
strengthening of the structure.
However, the Ti carbide in spectrum 1 is due to the
short time of the solution treatment, and then the rapid
cooling. It is considered to be too large to be beneficial
to the microstructure. The carbide is hexagonal, up to
6.15 μm in size, and locally precipitated at the triple
grain boundary. Due to its size, this carbide can be a
convenient place for the deposition of a topologically
close-packed (TCP) phase, and may be the initiator of
the appearance of microcracks.
The characterization of the microstructural changes
of the fracture surfaces occurring in the stated super-
alloys for the regimes R1 and R2 was made by scanning
electron microscopy and energy-dispersive X–ray spec-
trometry analysis, as shown in Figures 3 and 4.
A. MILOSAVLJEVIC et al.: THE INFLUENCE OF THE HEAT-TREATMENT REGIME ON A FRACTURE SURFACE ...
414 Materiali in tehnologije / Materials and technology 46 (2012) 4, 411–417
Table 3: Results of EDS analysis of the spectrums in Figure 2b (w/%)
Tabela 3: Rezultati EDS-analize spektrov s slike 2b (w/%)
Element Al Si Ti Cr Mn Fe Co Ni Mo
Spectrum 1 – – 80.99 3.83 – – 0.35 14.83 –
Spectrum 2 0.34 0.35 2.32 18.9 0.29 0.47 19.95 51.28 6.1
Spectrum 3 0.44 0.38 18.58 16.15 0.29 0.52 15.15 48.49 –
Table 4: Results of EDS analysis of the spectrums in Figure 3 (w/%)
Tabela 4: Rezultati EDS-analize spektrov s slike 3 (w/%)
Al Si Ti Cr Mn Fe Co Ni Mo
area in Fig. 3.a) 0.45 1.33 2.80 19.88 0.90 0.52 17.72 49.04 6.47
Spec1–Fig. 3.b) 0.47 0.48 8.79 19.13 0.39 0.80 16.40 46.17 6.37
Spec2–Fig. 3.b) 0.58 0.63 7.30 20.34 0.28 0.48 17.47 46.60 6.31
Figure 4: a) The appearance of a fracture surface of the Nimonic 263
superalloy after heat treatment R1, b) detail from Figure 4a –
unfavourable Ti carbide
Slika 4: a) Videz povr{ine preloma superzlitine Nimonic 263 po
toplotni obdelavi z na~inom R1, b) detajl s slike 4a – neza`eleni
Ti-karbidi
Figure 3: a) The appearance of a fracture surface after the heat
treatment R2, b) carbides at the fracture surface
Slika 3: a) Videz povr{ine preloma po toplotni obdelavi z na~inom
R2, b) karbidi na povr{ini preloma
The fractured surfaces were obtained from the tensile
tests and the results of the tensile strength, yield strength
(0.2 % offset) and elongation, together with the results
obtained for the superalloy Hastelloy S, are given in
Table 7.
Figures 3a and 3b show the fracture surfaces of the
superalloy Nimonic 263, after the heat treatment with
regime R2. The fractured surface is relatively homo-
geneous and the dimples are relatively uniform. In
Figure 3b the same fracture is presented at a higher
magnification. The deposition of Ti carbides can be
observed, as confirmed by the EDS analyses listed in
Table 4. The size of carbides is up to 1.16 μm and it is
believed that these carbides are not the cause of the
fracture.
A. MILOSAVLJEVIC et al.: THE INFLUENCE OF THE HEAT-TREATMENT REGIME ON A FRACTURE SURFACE ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 411–417 415
Table 5: Results of EDS-analysis of the spectrums in Figure 4 (w/%)
Tabela 5: Rezultati EDS-analize spektrov s slike 4 (w/%)
Al Si Ti Cr Mn Fe Co Ni Mo
The area in Fig. 4.a) 0.25 1.33 2.80 19.88 0.90 – 17.72 51.04 6.47
Spectrum 1 – Fig. 4.b) – – 46.55 12.01 0.51 0.21 10.20 28.47 2.16
Table 6: Results of EDS-analysis of the area in Figure 6a (w/%)
Tabela 6: Rezultati EDS-analize podro~ja s slike 6a (w/%)
Al Si Cr Mn Fe Ni Mo
Whole area in Fig. 6a 0.4 0.46 14.86 0.54 1.08 64.82 18.6
Whole area in Fig. 6b 0.38 0.48 15.04 0.56 1.33 65.41 17.8
Table 7: Mechanical properties of superalloys Nimonic 263 and Hastelloy S after applied heat treatments
Tabela 7: Mehanske lastnosti toplotno obdelanih zlitin Nimonic 263 in Hastelloy S
Mechanical
properties
R0.2 /(N/mm2) Rm /(N/mm2) A5/(%)
I II III I II III I II III
Nimonic 263 R1 550 573 545 820 835 832 39 38 39
Nimonic 263 R2 582 587 593 972 975 979 39 40 39
Hastelloy S 464 450 466 845 839 837 49 48 50
Figure 6: Fracture surfaces of the superalloy Hastelloy S taken with a
SEM
Slika 6: Povr{ina preloma superzlitine Hastelloy S (SEM)
Figure 5: Molybdenum carbides precipitated in the superalloy
Hastelloy S after heat treatment, taken with a light microscope
Slika 5: Karbidi molibdena, izlo~eni v toplotno obdelani superzlitini
Hastelloy S; svetlobni mikroskop
Figures 4a and 4b show the fractured surfaces of the
Nimonic 263 superalloy after heat treatment – regime
R1. In Figure 4b is a detail from Figure 4a – an
unfavourable precipitated carbide. The EDS analysis
results listed in Table 5 confirm the assumption that it is
a Ti carbide, and the place, morphology (irregular, hexa-
gonal-like) and size (3.95 μm) suggest it has a negative
influence on the mechanical characteristics of the
material. It is believed that the large carbides contributed
greatly to the breaking of the material.
Based on our observations of the samples of the
superalloy Hastelloy S using a light microscope, the dark
phases are visible, for which the energy-dispersive X-ray
spectrometry analysis shows they are the molybdenum
carbides. Their formation contributed the increased
content of molybdenum in the alloy, as well, compared
to the prescribed one.
The molybdenum carbides are segregated into arrays
and nests – Figures 5a and 5b. It is believed that these
carbides, together with the Cr carbides that occur in
these alloys, strengthened the alloy after the heat-treat-
ment process.
Figures 6a and 6b show the fractured surfaces of the
superalloy Hastelloy S. In Table 6 are the results of the
EDS analysis of the whole areas presented in Figure 6.
A higher content of molybdenum was observed, and this
is consistent with light-micrograph analyses. The struc-
ture of the fractures is homogeneous and uniform, as
confirmed by the EDS analysis conducted at the 150x
and 5 000-times magnifications (Table 6). The chemical
compositions of the two investigated surfaces differ very
little, which implies the homogeneity of the structure.
The mechanical properties obtained with the tensile
tests are listed in Table 7. Relatively high values of the
tensile strength, the yield strength (0.2 % offset) and the
elongation are in favour of a good heat treatment being
applied. In this paper, the short-solid solution time
(regime R1) results in lower values of the ultimate
tensile strength and yield strength, while the elongation
remains the same. The results listed in16, where the
authors applied different aging times for three groups of
samples, show that the various heat treatments have a
large effect on all the tensile properties of the material.
Compared to the literature data4,19 obtained for Nimonic
263 (R0.2 = 580 MPa, Rm = 970 MPa and A5 = 39 % ) we
observed slightly higher R0.2 and Rm values. According to
the literature4,20 the tensile-test characteristics of the
Hastelloy S alloy are: R0.2 = 444 MPa, Rm= 844 MPa and
A5 = 49 %. In this paper, the higher value of R0.2, the
yield strength (0.2 % offset), is obtained – the important
characteristic of the material used in design projects.
Based on the results in Table 7, the fracture appearance
and the EDS analysis, it can be assumed that during the
heat treatment the creation of undesirable phases has not
taken place, which would help in the breaking of the
material.
4 CONCLUSIONS
The microstructure and mechanical characteristics of
the superalloys Nimonic 263 and Hastelloy S, besides
processing route and chemical composition, depend to a
large extent on the applied heat treatment. The heat-
treatment processes of the superalloys Nimonic 263 and
Hastelloy S, carried out according to the regimes
described in this paper, indicate their considerable
influence on the microstructure transformations, and in
this way on the material properties. During the applied
heat-treatment regimes, the various micro-constituents
formed with a major influence on the properties of the
materials.
The heat treatment of the superalloy Nimonic 263,
which included 60 min of solid solution time (R2), rather
than the 10 min solid solution time (R1), resulted in a
better microstructure and better mechanical properties of
the stated superalloy. During both regimes of heat treat-
ment, Ti carbides participated, but according to the size,
the morphology and the distribution in regime R2 they
are favourable, while in regime R1 they are not. The
grains are finer and more uniformly distributed after the
heat treatment of regime R2 than after regime R1, i.e.,
with a longer solid-solution time. Also, the values of
mechanical properties: the hardness, the ultimate tensile
strength and the yield strength (0.2 % offset), are higher
when the time of the solid solution treatment is longer.
After the heat treatment of the nickel-based super-
alloy Hastelloy S, applied in this study, the molybdenum
carbides segregated. As a result, a higher value of the
yield strength is obtained. This characteristic is critical in
design projects, which makes these results very useful
for engineering practice.
Acknowledgements
This work was supported by the Ministry of Science
of the Republic of Serbia under contract number
TR-35040 and TR 34028.
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Grundstoffindustrie – Leipzig, 1975, 41
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Kim, C. Y. Jo, Materials Science and Engineering A, 523 (2009),
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A. MILOSAVLJEVIC et al.: THE INFLUENCE OF THE HEAT-TREATMENT REGIME ON A FRACTURE SURFACE ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 411–417 417

M. AMBRO@I^, L. GORJAN: UPOGIBNA TRDNOST KORUNDNE KERAMIKE: PRIMERJAVA RAZLI^NIH ...
UPOGIBNA TRDNOST KORUNDNE KERAMIKE:
PRIMERJAVA RAZLI^NIH TEORETI^NIH
PORAZDELITEV NA OSNOVI EKSPERIMENTALNIH
PODATKOV
BEND STRENGTH OF ALUMINA CERAMICS: A COMPARISON
OF DIFFERENT THEORETICAL DISTRIBUTIONS ON THE BASIS
OF EXPERIMENTAL DATA
Milan Ambro`i~1, Lovro Gorjan2,3
1Fakulteta za naravoslovje in matematiko, Univerza v Mariboru, 2000 Maribor, Slovenija
2Institut "Jo`ef Stefan", Jamova 39, 1000 Ljubljana, Slovenija
3Hidria AET d. o. o., Poljubinj 89A, 5220 Tolmin, Slovenija
lovro.gorjan@hidria.com
Prejem rokopisa – received: 2012-01-27; sprejem za objavo – accepted for publication: 2012-03-02
Statisti~no smo ovrednotili 5100 eksperimentalnih vrednosti upogibnih trdnosti testnih vzorcev iz redne proizvodnje korundnih
kerami~nih izdelkov. Primerjali smo teoreti~no izra~unano Weibullovo porazdelitev z dvema drugima pogosto uporabljenima
dvoparametri~nima porazdelitvama, normalno in log-normalno, da bi ugotovili, katera se najbolj sklada z meritvami. Za izra~un
ustreznih prostih parametrov smo uporabili metodo najve~je verjetnosti (maximum-likelihood method). Potem smo za
primerjavo rezultatov uporabili Q – Q-diagrame. Potrdili smo domnevo, da se z eksperimentalnimi vrednostmi trdnosti najbolj
ujema Weibullova porazdelitev.
Klju~ne besede: Weibullova porazdelitev, metoda najve~je verjetnosti, Q – Q-diagram
We have performed a statistical evaluation of 5100 experimental values of the bend strength of test pieces from a serial
production of alumina products. The Weibull distribution was compared with two other commonly used 2-parametric
distributions, i.e., normal and log-normal, in order to reveal which of them best matches the experiments. The
maximum-likelihood method was used to evaluate the corresponding parameters, and then a Q–Q plot was used for all the
statistics. We confirmed that the Weibull distribution describes the experimental strengths most accurately.
Keywords: Weibull distribution, maximum-likelihood method, Q–Q diagram
1 UVOD
Neenakost mno`ice merjenih vrednosti trdnosti
kerami~nih in drugih krhkih materialov v standardnih
mehanskih preskusih je posledica naklju~ne porazdelitve
napak v vzorcih, najve~krat pa te trdnosti opi{emo z
Weibullovo statisti~no porazdelitvijo1–4. Zanesljivost
teoreti~nih napovedi na osnovi Weibullove ali drugih
pogosto uporabljenih porazdelitvenih funkcij, ki jih izra-
~unamo za majhno {tevilo vzorcev, so veliko raziskovali
teoreti~no in eksperimentalno (najve~krat pa kombini-
rano) za razli~ne pogoje merjenja5–19. Znano je na primer,
da razli~ni ra~unski postopki sistemati~no dajejo
prevelike ali pa premajhne vrednosti prostih parametrov
raznih statisti~nih porazdelitev, ~e je {tevilo zlomljenih
vzorcev relativno majhno, npr. 30 ali {e manj7,8. Metoda
najve~je verjetnosti (ML = maximum-likelihood method)
se zaradi svoje u~inkovitosti pogosto uporablja za izra-
~un parametrov statisti~nih porazdelitev20,21.
Veliko avtorjev se je `e vpra{alo o upravi~enosti upo-
rabe Weibullove porazdelitve, zato so zanesljivost njenih
napovedi primerjali tudi z drugimi porazdelitvami: nor-
malno (Gaussovo), log-normalno itd.7,22. Za majhno
{tevilo testnih vzorcev namre~ ne moremo zanesljivo
izklju~iti veljavnosti vseh teh porazdelitev, posebno ne
Gaussove22. Eden od zanimivih na~inov vizualne pri-
merjave eksperimentalnih trdnosti vzorcev s teoreti~nimi
vrednostmi iz ustrezne porazdelitvene funkcije so Q –
Q-diagrami (Q = quantile)23,24. Ujemanje med teorijo in
eksperimentom je dobro, ~e so to~ke v diagramu zelo
blizu premici, ki oklepa z vodoravno osjo kot 45° (sime-
trala prvega kvadranta).
V ~lanku opisujemo statisti~no obdelavo 5100 iz-
merjenih vrednosti upogibne trdnosti vzorcev iz redne
proizvodnje korundnih kerami~nih izdelkov. Za oceno
parametrov treh razli~nih teoreti~nih porazdelitev smo
uporabili ML-metodo, potem pa med drugim primerjali
ujemanje teoreti~no izra~unanih in eksperimentalnih
trdnosti s Q – Q-diagrami.
2 MERITVE TRDNOSTI
Kerami~ne testne vzorce v obliki kvadrov dimenzij 4
mm × 3 mm × 45 mm smo izdelali v podjetju Hidria
AET, d. o. o., s tehniko nizkotla~nega brizganja v kalupe.
V enem letu se nam je nabralo 5100 meritev trdnosti
vzorcev iz 425 proizvodnih serij, po 12 zlomljenih vzor-
cev iz vsake serije. Vzorci so bili iz korundne keramike
Materiali in tehnologije / Materials and technology 46 (2012) 4, 419–422 419
UDK 519.61/.64:666.3/.7 ISSN 1580-2949
Strokovni ~lanek/Professional article MTAEC9, 46(4)419(2012)
Al2O3 s 95-odstotno teoreti~no gostoto, pripravljeni s
sintranjem tri ure pri temperaturi 1640 °C. Ta material se
v glavnem uporablja za elektri~no izolacijo in se zanj ne
zahtevajo najvi{je trdnosti, tako da ga v podjetju
ozna~ujemo kot "Al2O3 keramiko srednjih trdnosti".
Trdnosti smo izra~unali z ena~bo za 4-to~kovni upogibni
test21:
 =
−3
2
2 1
2
F L L
ah
( )
(1)
kjer je  upogibna trdnost vzorca, F zlomna sila, L2 = 40
mm in L1 = 20 mm sta zunanji in notranji razmik
podpor, a = 4 mm je {irina vzorca, h = 3 mm pa njegova
debelina.
3 STATISTI^NI MODEL IN GRAFI^NA
PONAZORITEV
Na{a naklju~na (statisti~na) spremenljivka je upo-
gibna trdnost . V ra~unih uporabljamo obe porazde-
litveni funkciji: verjetnostno gostoto p() in kumulativno
verjetnostno funkcijo P p x x
o
( ) ( )

= ∫ d . Primerjali bomo
ujemanje med eksperimentalnimi podatki in tremi napo-
vedanimi 2-parametri~nimi porazdelitvami: Weibullovo,
normalno (Gaussovo) in log-normalno.
3.1 Postopek za ocenitev stopnje ujemanja teorije in
eksperimenta
Podroben postopek za oceno dveh prostih parametrov
in vizualne ponazoritve ujemanja dane teoreti~ne poraz-
delitvene funkcije z eksperimentalnimi podatki sestoji iz
naslednjih korakov:
a) Za vsako od treh porazdelitev posebej dobimo
najprimernej{i vrednosti prostih parametrov z metodo
najve~je verjetnosti (ML). N (= 5100) izmerjenih
trdnosti, i, i = 1 do N, vstavimo v verjetnostno
gostoto p(a, b; ), kjer sta a in b ustrezna parametra,
npr., a  m in b  0W za Weibullovo porazdelitev
(gl. spodaj). Z metodo ML poi{~emo najve~jo
mogo~o vrednost naslednje funkcije glede na vse
mogo~e vrednosti a in b:
Y p a b p a bi
i
N
i
i
N
= =
= =
∏ ∑ln ln( ( , ; ) ( , ; ) 
1 1
(2)
V ustreznih (dveh) ena~bah sta odvoda funkcije Y po
parametrih a in b enaka ni~.
b) Eksperimentalne vrednosti trdnosti i uredimo po
nara{~ajo~ih vrednostih, kjer gre indeks i od 1 do N =
5100. Vsaki {tevilki i (in s tem ustrezni vrednosti
trdnosti) priredimo verjetnost Pi za zlom, pri tem pa
si pomagamo z enim od pogosto uporabljenih ocen iz
literature: 22
P
i
Ni
=
− 05.
(3)
Vrednost Pi ustreza kumulativni verjetnostni funkciji
P() za izbrano teoreti~no porazdelitev. Namesto z oceno
(3) smo poskusili s {e drugimi ocenami za Pi iz literature
in ugotovili, da ta izbira ni~ ne vpliva na kon~ne
rezultate, vsaj za tolik{no {tevilo merjenih vzorcev.
c) Za vsako vrednost Pi, dobljeno iz ena~be (3), smo z
obratom teoreti~ne funkcije P(a, b; ) z `e prej izra-
~unanima parametroma a in b (gl. to~ko a) izra~unali
ustrezno teoreti~no pri~akovano trdnost i,th.
~) Uporabili smo poenostavljeno razli~ico grafi~ne
ponazoritve Q – Q za primerjavo teoreti~nih trdnosti
i,th z eksperimentalnimi i. Pare vrednosti (i,th, i) v
diagramu smo prikazali tako, da teoreti~na kompo-
nenta ustreza vodoravni osi, eksperimentalna pa
navpi~ni. ^e se torej teoreti~na statisti~na napoved
odli~no ujema z eksperimentom, le`ijo to~ke v dia-
gramu vse blizu premice z naklonom 45° proti
osema.
d) Kvantitativno lahko opredelimo ujemanje teoreti~nih
napovedi z eksperimentalnimi podatki s faktorjem R2:
R
i
i
N
i th
i
i
N
i
2 1
2
1
2
1= −
−
− < >
=
=
∑
∑
( )
( )
, 
 
(4)
kjer je <i> izra~unana povpre~na vrednost eksperi-
mentalnih trdnosti. Vrednost R2 blizu 1 pomeni dobro
ujemanje.
3.2 Primerjane teoreti~ne porazdelitve
Drugi parameter bomo zaradi analogije pri vseh treh
porazdelitvah ozna~ili s podobnim simbolom: 0W, 0N in
0LN. To je umeritveni parameter, ki ima enoto trdnosti.
Prvi parameter pa ima pri treh porazdelitvah povsem
druga~en pomen. Podali bomo le porazdelitveno funkcijo
p za vse tri porazdelitve. Porazdelitvene funkcije za
Weibullovo, normalno in log-normalno porazdelitev
ozna~imo po vrsti s pW, pN in pLN:
p
m
W
W W
m
W
m
( ) exp





=
⎛
⎝
⎜
⎞
⎠
⎟ ⋅ −
⎛
⎝
⎜
⎞
⎠
⎟
⎛
⎝
⎜⎜
⎞
⎠
⎟⎟
−
0 0
1
0
(5a)
pN
W
m
N
( ) exp
 


 

=
⎛
⎝
⎜
⎞
⎠
⎟ ⋅ −
⎛
⎝
⎜
⎞
⎠
⎟
⎛
⎝
⎜
⎞
⎠
⎟
−
1 1
20
1
0
2
π
(5b)
p
w wLN
( ) exp
ln ln

 
  = ⋅ −
⎛
⎝
⎜
⎞
⎠
⎟
⎛
⎝
⎜
⎞
⎠
⎟1 1
2
π
(5c)
Pri Weibullovi porazdelitvi imenujemo brezdimen-
zijski parameter m Weibullov modul, 0W pa umeritveni
parameter. Pri normalni porazdelitvi imata parametra
enostaven pomen: 0N je povpre~na trdnost,  pa njena
standardna deviacija. Pri log-normalni porazdelitvi so
normalno porazdeljeni logaritmi trdnosti.
M. AMBRO@I^, L. GORJAN: UPOGIBNA TRDNOST KORUNDNE KERAMIKE: PRIMERJAVA RAZLI^NIH ...
420 Materiali in tehnologije / Materials and technology 46 (2012) 4, 419–422
4 REZULTATI IN DISKUSIJA
Vseh 5100 eksperimentalnih trdnosti prikazuje slika
1, kjer je vsaka vrednost zapisana v kronolo{kem
zaporedju (glede na ~as meritve), njihova aritmeti~na
povpre~na vrednost pa je <> = 289,557 MPa. V tabeli 1
so podani z ML-metodo ocenjeni parametri treh porazde-
litev, skupaj z R2 faktorjem, ki je najve~ji pri Weibullovi
porazdelitvi. Slika 2 prikazuje ustrezni Q – Q-diagram
za to porazdelitev, druga dva pa sta temu kvalitativno
podobna, le ujemanje to~k s premico z nagibom 45° je
nekoliko slab{e.
Nazadnje smo naredili {e primerjavo treh teoreti~nih
porazdelitev glede ujemanja teoreti~ne pri~akovane
trdnosti <>, njene standardne deviacije  in "kubi~ne
deviacije" 3σ z ustreznimi eksperimentalnimi vrednost-
mi. Kubi~no deviacijo smo definirali takole:
   3
33= < − < > >( )
kjer trikotni oklepaji ozna~ujejo statisti~no povpre~je.
Rezultati so prikazani v tabeli 2. V prvi vrstici v vsa-
kem od treh parov so eksperimentalne vrednosti teh
parametrov, v drugi pa teoreti~ne, in sicer v levem
stolpcu (oznaka *) izra~unane direktno iz prostih para-
metrov vsake porazdelitve, v desnem stolpcu (oznaka
**) pa so izra~unani s statisti~nim povpre~jem "simuli-
ranih teoreti~nih" trdnosti i,th. Zaradi velikega {tevila
podatkov lahko pri~akujemo, da se bodo vrednosti,
ozna~ene z **, zelo dobro ujemale s tistimi, ozna~enimi
z *, in to tabela 2 tudi potrjuje. Medtem ko se pri~ako-
vana vrednost trdnosti in njena standardna deviacija pri
vseh treh porazdelitvah dokaj dobro ujemata z
ustreznimi eksperimentalnimi parametri, pa daje vsaj
pribli`no pravilen rezultat za kubi~no deviacijo samo
Weibullova porazdelitev.
5 SKLEP
Primerjava treh teoreti~nih porazdelitev s 5100 eks-
perimentalnimi trdnostmi keramike Al2O3 je pokazala, da
je najustreznej{i statisti~ni opis z Weibullovo 2-para-
M. AMBRO@I^, L. GORJAN: UPOGIBNA TRDNOST KORUNDNE KERAMIKE: PRIMERJAVA RAZLI^NIH ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 419–422 421
Tabela 1: ML-parametri in R2 faktor za tri porazdelitve
Table 1: ML-parameters and the R2 factor for the three distributions
Porazdelitev 1. parameter 2. parameter R2
Weibullova m = 9,048 0W = 305,54 MPa 0,9984
Normalna  = 37,49 MPa 0N = 289,56 MPa 0,9855
Log-normalna w = 0,1372 0LN = 286,86 MPa 0,9468
Tabela 2: Eksperimentalni in teoreti~ni statisti~ni parametri porazdelitev
Table 2: Experimental and theoretical statistical parameters of the distributions
<>  3
Eksperiment 289,56 37,49 29,58
Weibullova *289,41 ** 289,41 *38,26 **38,26 *32,15 *32,12
Normalna *289,56 ** 289,56 *37,49 **37,48 *0 *1,19
Log-normalna *289,57 ** 289,56 *39,92 **39,91 *29,80 *29,75
Slika 2: Q – Q-diagram za Weibullovo porazdelitev
Figure 2: Q – Q-diagram for the Weibull distribution
Slika 1: 5100 vrednosti upogibne trdnosti; {tevilke na vodoravni osi
ustrezajo ~asovnemu vrstnemu redu meritev
Figure 1: 5100 bending-strength values; the numbers on the horizon-
tal axis correspond to the time sequence of the measurements
metri~no porazdelitveno funkcijo, kar med drugim
potrjuje najvi{ji faktor R2: 99,84 %, razlike v vizualni
primerjavi statistik s Q – Q-diagrami pa so manj o~itne.
Omenimo lahko {e, da daje ML-metoda za normalno
porazdelitev natan~ne eksperimentalne vrednosti prvih
dveh statisti~nih parametrov v tabeli 2 (pri~akovana
vrednost trdnosti in standardna deviacija); v tem je torej
normalna porazdelitev nekaj bolj{a od Weibullove,
vendar pa daje popolnoma napa~en rezultat za kubi~no
deviacijo. Ve~ podrobnosti lahko bralec najde drugje25.
Zahvala
To raziskavo sta podprla Ministrstvo za {olstvo in
{port ter Evropski socialni sklad. Zahvaljujemo se
podjetju Hidria AET za eksperimentalne podatke.
6 LITERATURA
1 W. Weibull, A statistical distribution function of wide applicability, J.
Appl. Mech., 18 (1951), 293–297
2 R. V. Curtis, A. S. Juszczyk, Analysis of strength data using two- and
three-parameter Weibull models, J. Mater. Sci., 33 (1998),
1151–1157
3 H. Peterlik, N. Orlovskaja, W. Steinkellner, K. Kromp, Prediction of
strength of recrystallized siliconcarbide from pore size measurement
– Part I – The bimodality of the distribution, J. Mater. Sci., 35
(2000), 699–707
4 Q. S. Li, J. Q. Fang, D. K. Liu, J. Tang, Failure probability prediction
of concrete components, Cem. Concr. Res., 33 (2003), 1631–1636
5 D. Wu, J. Zhou, Y. Li, Methods for estimating Weibull parameters
for brittle materials, J. Mater. Sci., 41 (2006), 5630–5638
6 D. Wu, J. Zhou, Y. Li, Unbiased estimation of Weibull parameters
with the linear regression method, J. Eur. Ceram. Soc., 26 (2006),
1099–1105
7 R. Danzer, T. Lube, P. Supancic, Monte Carlo simulations of strength
distributions of brittle materials – Type of distribution, specimen and
sample size, Z. Metallkunde, 92 (2001), 773–783
8 B. Bergman, On the estimation of the Weibull modulus, J. Mater. Sci.
Lett., 3 (1984), 689–692
9 A. Khalili, K. Kromp, Statistical properties of Weibull estimators, J.
Mater. Sci., 26 (1991), 6741–6752
10 J. Gong, A new probability index for estimation Weibull modulus for
ceramics with the least-square method, J. Mater. Sci. Lett., 19
(2000), 827–829
11 E. Barbero, J. Fernandez-Saez, C. Navarro, Statistical distribution of
the estimator of Weibull modulus, J. Mater. Sci. Lett., 20 (2001),
847–849
12 I. J. Davies, Empirical correction factor for the best estimate of
Weibull modulus obtained using linear least squares analysis, J.
Mater. Sci. Lett., 20 (2001), 997–999
13 D. F. Wu, Y. D. Li, J. P. Zhang, L. Chang, D. H. Wu, Z. P. Fang, Y.
H. Shi, Effects of the number of testing specimens and the estimation
methods on the Weibull parameters of solid catalysts, Chem. Eng.
Sci., 56 (2001), 7035–7044
14 L. Song, D. Wu, Y. Li, Optimal probability estimators for deter-
mining Weibull parameters, J. Mater. Sci. Lett., 22 (2003),
1651–1653
15 J. A. Griggs, Y. Zhang, Determining the confidence intervals of
Weibull parameters estimated using a more precise probability
estimators, J. Mater. Sci. Lett., 22 (2003), 1771–1773
16 I. J. Davies, Best estimate of Weibull modulus obtained using linear
least squares analysis: An improved empirical correction factor, J.
Mater. Sci., 39 (2004), 1441–1444
17 T. Tanaka, H. Nakayama, A. Sakaida, T. Imamichi, Evaluation of
Weibull parameters for static strengths of ceramics by Monte Carlo
simulation, Mater. Sci. Res. International, 1 (1995), 51– 8
18 B. Faucher, W. R. Tyson, On the determination of Weibull parame-
ters, J. Mater. Sci. Lett., 7 (1988), 1199–1203
19 R. Langlois, Estimation of Weibull parameters, J. Mater. Sci. Lett.,
10 (1991), 1049–1051
20 M. Cacciari, G. Mazzanti, G. C. Montanari, Comparison of maxi-
mum likelihood unbiasing methods for the estimation of the Weibull
parameters, IEEE Trans. On El. Ins., 3 (1996), 18–27
21 ASTM C 1239 – 95: Standard practice for Reporting Uniaxial
Strength Data and Estimating Weibull Distribution Parameters for
Advanced Ceramics, American Society for Testing and Materials,
Philadelphia, 1995
22 C. Lu, R. Danzer, F. D. Fischer, Fracture statistics of brittle
materials: Weibull or normal distribution, Phys. Rev. E, 65 (2002),
067102
23 J. D. Gibbons, S. Chakraborti, Nonparametric statistical inference,
CRC Press 2003
24 R. Gnanadesikan, M. B. Wilk, Probability plotting methods for the
analysis of data, Biometrika, 55 (1968), 1–17
25 L. Gorjan, M. Ambro`i~, Bend strength of alumina ceramics: A
comparison of Weibull statistics with other statistics based on very
large experimental data set, J. Eur. Ceram. Soc., 32 (2012) 6,
1221–1227
M. AMBRO@I^, L. GORJAN: UPOGIBNA TRDNOST KORUNDNE KERAMIKE: PRIMERJAVA RAZLI^NIH ...
422 Materiali in tehnologije / Materials and technology 46 (2012) 4, 419–422
M. TORKAR et al.: CREVICE CORROSION OF STAINLESS-STEEL FASTENING COMPONENTS ...
CREVICE CORROSION OF STAINLESS-STEEL
FASTENING COMPONENTS IN AN INDOOR
MARINE-WATER BASIN
[PRANJSKA KOROZIJA PRITRDILNIH KOMPONENT IZ
NERJAVNEGA JEKLA V NOTRANJEM BAZENU Z MORSKO
VODO
Matja` Torkar, Franc Tehovnik, Matja` Godec
Institute of Metals and Technology, Lepi pot 11, SI-1000 Ljubljana, Slovenia
matjaz.torkar@imt.si
Prejem rokopisa – received: 2012-02-07; sprejem za objavo – accepted for publication: 2012-02-16
Equipment made from austenitic stainless steel corroded already after six months of the operation of an indoor marine-water
basin. Two super chlorinations were performed during this period. Corroded stainless-steel components and different fastening
components were investigated to detect the corrosion causes and improve the bathers’ safety. Pitting corrosion was observed on
flat surfaces, while crevice corrosion prevailed on the nuts and spring washers of bolted joints. The main reasons for the
corrosion damages were: a high concentration of chlorides, a deficient control of sacrificial anodes, a lower corrosion resistance
of spring washers and an inaccurate montage of some fastening components. The corrosion processes due to chloride ions can
be reduced with frequent washing of all stainless-steel components with clear water and timely replacements of the sacrificing
anodes.
Keywords: stainless steel, SEM, cathodic protection, chlorination, pitting corrosion, crevice corrosion
V notranjem bazenu z morsko vodo je vsa oprema, izdelana iz avstenitnega nerjavnega jekla, po {estih mesecih obratovanja
za~ela rjaveti. V tem obdobju sta bili izvr{eni tudi dve superkloriranji. Povr{ine komponent in razli~ni pritrdilni elementi iz
nerjavnega jekla, prizadeti s korozijo, so bili preiskani zaradi ugotavljanja vzroka korozije in tudi zaradi varnosti kopalcev. Na
ravnih povr{inah je prevladovala jami~asta korozija, {pranjska korozija pa je prevladovala pri maticah, vzmetnih podlo`kah in
pri vija~enih spojih. Glavni razlogi za korozijske po{kodbe so bili kombinacija velike koncentracije kloridov, pomanjkljiva
kontrola `rtvenih elektrod, slab{a korozijska obstojnost vzmetnih podlo`k in nenatan~na monta`a drugih pritrdilnih komponent.
Korozijske procese zaradi kloridnih ionov zmanj{amo s pogostim ~i{~enjem in spiranjem komponent iz nerjavnega jekla s ~isto
vodo in s pravo~asno menjavo `rtvenih elektrod.
Klju~ne besede: nerjavno jeklo, vrsti~na elektronska mikroskopija (SEM), katodna za{~ita, klorinacija, jami~asta korozija,
{pranjska korozija
1 INTRODUCTION
The atmosphere of the indoor basins is one of the
most aggressive in a building environment. Under spe-
cific temperature conditions chlorine containing chemi-
cal species in the vapours of the pool water can condense
onto the stainless-steel components and dry out. After
several repetitions of this process very aggressive con-
centrations of chlorine-containing mixtures may build
up. As the presence of marine water accelerates the chlo-
ride attack1, the use of austenitic stainless steel alloyed
with molybdenum is recommended. The situation is ag-
gravated when the austenitic stainless-steel components
are not regularly cleaned.
Galvanic or sacrificial anodes are often used for the
corrosion protection of the basin equipment. Stain-
less-steel components are connected with wires to the
sacrificing anodes absorbing the corrosion currents to the
anodes as they are designed to have a more negative
electrochemical potential than the equipment to be pro-
tected. A sacrificing anode2 (made of Zn or Mg alloys)
continues to corrode (sacrifice), being consumed until a
replacement becomes necessary. If the anode is not re-
placed in time, the system loses the protective role and
the corrosion processes on the components become more
intensive.
Bathers also introduce contaminants into the water.
Being added to the water, the chlorine hydrolyzes rapidly
and produces hypochlorous acid (HOCl) and hydrochlo-
ric acid (HCl) as described by equation (1):
Cl2 +H2O 
 HOCl + HCl (1)
Depending on the variation of the water pH the con-
centration of hypochlorous acid versus the concentration
of hypochlorite (OCl–) varies as well. Hypochlorous acid
is weak and it ionizes at pH 7.5 and at 25 °C 3 by equa-
tion (2).
HOCl 
 H+ + OCl– (2)
When nitrogen and HOCI combine, chloramines like
monochloramine, dichloramine and nitrogen trichloride
are formed according to the equations (3 to 5):
NH3 + HOCl 
 NH2Cl + H2O (3)
NH2Cl + HOCl 
 NHCl2 + H2O (4)
NHCl2 + HOCl 
 NCl3 + H2O (5)
Materiali in tehnologije / Materials and technology 46 (2012) 4, 423–427 423
UDK 620.19 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 46(4)423(2012)
The reactions (1) to (5) are in equilibrium and occur
in the forward and reverse directions.
The chloramine causes the characteristic chlorine
pool smell4. The nitrogen trichloride is more problematic
in the case of a heated indoor basin because more nitro-
gen trichloride is vaporised5.
Two methods are used to eliminate chloramines6,7:
break-point chlorination and super-chlorination. The
later process is known as shocking because, with an ad-
dition of chlorine, the total chlorine in the basin water
rises up to an amount ten times above the normal chlo-
rine level. This operation is performed in a basin without
any bathers. The basin can be used again when the level
of chlorine is back down to around 5 μg/g. In general, the
super-chlorination process is harmful to stainless-steel
equipment and accelerates pitting- and crevice
corrosions.
Chlorine is 100 % effective when the pH of water is
5.5. Such pH is too acidic for people to comfortably
swim and it also accelerates the corrosion. At a higher
pH, up to 7.0, algae will grow. The pH of 7.2 is the most
comfortable for swimmers; however, having a pH of 7.2
chlorine is only 50-% effective and so a free-chlorine
level should be increased to a minimum of 0.6 μg/g. The
pH range of the pool water should be kept between 6.8
and 8 8.
The latest research on super-chlorination revealed
that the process is effective only for eliminating inor-
ganic, ammonia-based chloramines. It was established
that large doses of free chlorine react with organic con-
taminants and form a variety of disinfectants that are
hazardous to the swimmers’ health9.
A new indoor basin with marine water was in opera-
tion for about six months. After that period traces of rust
appeared around the screws and bolts, on all the surfaces
of the equipment and on the stainless-steel decorative el-
ements. The pillars holding the roof of the hall were dec-
orated with vertical polished stainless-steel strips. The
surface was covered with numerous speckles of rust, eas-
ily removable with a cloth. Corrosion pits were present
on the steel surface under the rust. An in-situ control of
all the basin components revealed rusted sites around the
screws in the plastic inlets for fresh water (Figure 1), in
the frames for the underwater lights and also on the
joints connecting a step with a vertical holder tube (Fig-
ure 2). In fact, all stainless steel parts were more or less
corroded.
The aim of this investigation was to reveal the rea-
sons for general corrosion on all stainless-steel decora-
tive and fastening components around and in the swim-
ming pool. The safety of bathers was endangered also
because of the rusted fastening joints of the vertical
stairs.
2 EXPERIMENTAL PROCESS
After only six months of the operation corrosion
products were observed on all the stainless-steel compo-
nents. The rusted screws on the plastic frame of the
fresh-water inlet, a nut, a spring washer and a bolt from
the vertical stairs were removed and investigated.
An individual stair was fixed with a plastic inlet and a
bolt on the vertical stair-holder tube. The rust was spread
around the bolted connection (Figure 2), while no rust
was observed around the direct contact of a longer bolt
and the step holder. The removal of the step showed that
the rust originated in the joint of the spring washer and
the nut. The rust was concentrated around the spring
washer, below the nut and spread around the bolted
joints. The spring washer was heavily corroded.
A photograph of the rusted nut, the washer and the
cross-head screw was taken (Figure 3) before further
cleaning and the rust was removed from all the investi-
gated samples in an ultrasound bath (Figure 4).
The surfaces damaged with the corrosion were exam-
ined with scanning electron microscopy (SEM) Jeol
M. TORKAR et al.: CREVICE CORROSION OF STAINLESS-STEEL FASTENING COMPONENTS ...
424 Materiali in tehnologije / Materials and technology 46 (2012) 4, 423–427
Figure 2: Rust on the vertical holder around the bolted-step
connection
Slika 2: Rja na vertikalnem nosilcu okrog vija~enega spoja stopnice
Figure 1: Rusted cross-head screws at the bottom of the pool
Slika 1: Rjasti vijaki s kri`no glavo z dna bazena
JSM-6500F having a field-emission gun and analysed
with energy dispersive X-ray spectroscopy (EDS).
3 RESULTS AND DISCUSSION
Inside the screw head we observed dip caverns (Fig-
ure 5 and 6) resulting from crevice corrosion. The basic
conditions for the start of crevice corrosion are a wet en-
vironment with chlorides and a narrow crevice where a
renewal of a damaged protective oxide layer on stainless
steel is not possible due to the restricted oxygen diffu-
sion in the crevice. Chloride ions from the salt water mi-
grate into the crevice and increase the solution acidity
accelerating the corrosion attack on the protective layer
on the stainless-steel surface. The results of crevice cor-
rosion are either deep corrosion caverns, as observed in
M. TORKAR et al.: CREVICE CORROSION OF STAINLESS-STEEL FASTENING COMPONENTS ...
Materiali in tehnologije / Materials and technology 46 (2012) 4, 423–427 425
Figure 4: Corrosion-damaged structural components after cleaning
with ultrasound
Slika 4: S korozijo prizadeti konstrukcijski elementi po ~i{~enju z
ultrazvokom
Figure 7: Corroded screw (SEM)
Slika 7: S korozijo po{kodovan vijak (SEM)
Figure 6: Detail from Figure 5: Crevice corrosion of a screw head
(SEM)
Slika 6: Detajl s slike 5: {pranjska korozija glave vijaka (SEM)
Figure 5: Screw damaged by crevice corrosion (SEM)
Slika 5: Vijak, po{kodovan s {pranjsko korozijo (SEM)
Figure 3: A nut, a spring washer, a structural bolt of the stairs and a
cross-head screw, all cowered with rust
Slika 3: Matica, vzmetna podlo`ka, vijak stopnice in vijak s kri`no
glavo, vsi pokriti z rjo
the screw head, or shallow pits, as observed on the screw
body (Figure 7).
Corrosion caverns were also observed in the spring
washer (Figure 8, 9), taken from the connection step of
the vertical holder. A test of the spring washer with a
magnet revealed slight magnetism in the cold-formed
spring washer, probably due to the presence of
deformation-induced martensite in the washer. This
martensite in the austenitic stainless steel is magnetic
and has a lower resistance to corrosion9. Deformation-
induced martensite can be transformed back to austenite
by heating at the temperature of 1 050 °C and by water
rapid cooling. In the examined case, the spring washer
was in the cold deformed state to keep the spring
properties and for this reason the spring washer was the
most sensitive component to corrosion. In general, there
are two possible causes for the corrosion of a spring
washer in a chloride environment: either crevice
corrosion10,11 due to geometrical reasons (the presence of
a crevice) or stress corrosion12 due to internal stresses in
deformation-induced martensite of cold deformed steel.
In the examined case it is supposed that the first cause
prevailed because of the presence of corrosion caverns in
the spring washer.
The nut was from austenitic stainless steel, it was
non-magnetic and after the removal of the rust the shal-
low pits were observed on the surface only. Comparing
the corrosion damages on the nut and the spring washer,
it looks that most of the rust originated from the spring
washer.
The rust from the spring washer was spread also on
the bolt and the bolt surface was coloured by the rust. No
corrosion damages were observed on the bolt surface af-
ter being cleaned in the ultrasound bath.
An EDS analysis confirmed that the nut and the
spring washer were made of austenitic stainless steel
AISI 316 with 2 % to 3 % of molybdenum added to in-
crease the corrosion resistance of stainless steel in a
chloride environment.
The cross-head screw from the bottom of the basin
did not contain molybdenum and was, thus, made of
AISI 304 austenitic stainless steel that is less corrosion
resistant in a chloride environment.
Stainless-steel components are usually corrosion pro-
tected with sacrificing electrodes made of Mg or Zn al-
loys. An in-situ control revealed that the sacrificing elec-
trodes that were dissolved had neither been periodically
checked nor replaced. For this reason, the austenitic
stainless-steel stairs and fences were more serious at-
tacked by corrosion.
4 CONCLUSIONS
Examinations revealed that after six months of the
operation, severe corrosion damages appeared on all the
stainless-steel components in the hall and in the basin
with marine water.
Decorative bands, fences, the steps of vertical stairs
and other austenitic stainless-steel components were
damaged either by pitting or by crevice corrosion.
The spring washers and nuts were the most sensitive
structural elements to corrode in the examined chloride
environment.
The super-chlorination performed twice over a short
period additionally increased the corrosion processes
within the basin and the basin hall.
The corrosion processes due to chloride ions can be
reduced with frequent cleaning of all stainless-steel com-
ponents with clear water and timely replacements of the
sacrificial anodes.
For the bathers’ safety sake, all the vital joint connec-
tions in the vertical stairs need to be checked periodically
and, if necessary, replaced in time.
M. TORKAR et al.: CREVICE CORROSION OF STAINLESS-STEEL FASTENING COMPONENTS ...
426 Materiali in tehnologije / Materials and technology 46 (2012) 4, 423–427
Figure 8: Crevice corrosion of a spring washer (SEM)
Slika 8: [pranjska korozija vzmetne podlo`ke (SEM)
Figure 9: Remains of NaCl around a crevice-corrosion cavern (SEM)
Slika 9: Ostanki NaCl okrog izjede, povzro~ene s {pranjsko korozijo
(SEM)
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