VSEBINA – CONTENTS
PREGLEDNI ^LANEK – REVIEW ARTICLE
An overview of the influence of stainless-steel surface properties on bacterial adhesion
Vpliv lastnosti povr{ine nerjavnega jekla na adhezijo bakterij
M. Ho~evar, M. Jenko, M. Godec, D. Drobne . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 609
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Adsorption of hexavalent chromium from an aqueous solution of steel-making slag
Adsorpcija heksavalentnega kroma iz vodne raztopine jeklarske `lindre
A. [trkalj, Z. Glava{, G. Matija{i} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 619
Structural, mechanical and cytotoxicity characterization of as-cast biodegradable Zn-xMg (x = 0.8–8.3 %) alloys
Strukturne, mehanske in citotoksi~ne lastnosti biorazgradljive Zn-xMg (x = 0,8–8,3 %) zlitine v litem stanju
J. Kubásek, I. Pospí{ilová, D. Vojtìch, E. Jablonská, T. Ruml . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 623
Structural characterization of a bulk and nanostructured Al-Fe system
Karakterizacija strukture osnove in nanostrukture sistema Al-Fe
A. Fekrache, M. Yacine Debili, S. Lallouche . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 631
Wear behavior of Al/SiC/graphite and Al/FeB/graphite hybrid composites
Vedenje hibridnih kompozitov Al/SiC/grafit in Al/FeB/grafit pri obrabi
S. ªahýn, N. Yüksel, H. Durmuº, S. Gençalp Ýrýzalp. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 639
Mathematical model for an Al-coil temperature calculation during heat treatment
Matemati~ni model za izra~un temperature v Al-kolobarju med toplotno obdelavo
F. Vode, F. Tehovnik, J. Burja, B. Arh, B. Podgornik, D. Steiner Petrovi~, M. Malen{ek, L. Ko~evar, M. La`eta . . . . . . . . . . . . . . . . . . . 647
Investigating the effects of cutting parameters on the hole quality in drilling the Ti-6Al-4V alloy
Preiskava vpliva parametrov rezanja na kvaliteto izvrtine, izvrtane v zlitino Ti-6Al-4V
Y. Hýþman Çelik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 653
Influence of graphite on the hardness and wear behavior of AA6061–B4C composite
Vpliv grafita na trdoto in vedenje kompozita AA6061–B4C pri obrabi
S. Prabagaran, Govindarajulu Chandramohan, P. Shanmughasundaram . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 661
Tribology of CrAg7N coatings deposited on Vanadis 6 ledeburitic tool steel
Tribologija prevlek CrAg7N na ledeburitnem orodnem jeklu Vanadis 6
P. Bílek, P. Jur~i, M. Hudáková, ¼. ^aplovi~, M. Novák . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 669
Morphology and magnetic properties of Fe3O4-alginic acid nanocomposites
Morfologija in magnetne lastnosti nanokompozitov Fe3O4-alginska kislina
M. KaŸmierczak, K. Pogorzelec-Glaser, A. Hilczer, S. Jurga, £. Majchrzycki, M. Nowicki, R. Czajka, F. Matelski, R. Pankiewicz,
B. £êska, L. Kêpiñski, B. Andrzejewski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 675
Microstructural comparison of the thermomechanically treated and cold deformed Nb-microalloyed trip steel
Primerjava mikrostruktur termomehansko obdelanega in hladno deformiranega, z Nb-mikrolegiranega trip-jekla
A. Grajcar, K. Radwañski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 679
Control of the metallurgical processing of ICDP cast irons
Kontrola metalur{ke obdelave litega `eleza ICDP
J. Hampl, T. Válek, P. Lichý, T. Elbel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 685
Synthesis comparison and characterization of chitosan-coated magnetic nanoparticles prepared with different methods
Primerjava postopkov in karakterizacija magnetnih nanodelcev, prevle~enih s hitozanom
G. Hojnik Podrep{ek, @. Knez, M. Leitgeb . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 689
Study of phase transformations in Cr-V tool steel
[tudij faznih premen v orodnem jeklu Cr-V
M. Pa{ák, R. ^i~ka, P. Bílek, P. Jur~i, ¼. ^aplovi~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 693
Model antimicrobial polymer system based on poly(vinyl chloride) and crystal violet
Model protimikrobnega polimernega sistema na osnovi polivinilklorida in kristal violeta
J. Klofá~, I. Kuøitka, P. Ba`ant, K. Jedli~ková, J. Sedlák . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 697
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 48(5)607–800(2014)
MATER.
TEHNOL.
LETNIK
VOLUME 48
[TEV.
NO. 5
STR.
P. 607–800
LJUBLJANA
SLOVENIJA
SEP.–OCT.
2014
Effect of the tool geometry and welding parameters on the macrostructure, fracture mode and weld strength of friction-stir
spot-welded polypropylene sheets
Vpliv geometrije orodja in parametrov varjenja na makrostrukturo, vrsto preloma in trdnost zvara pri torno-vrtilnem to~kastem varjenju
polipropilenskih plo{~
M. Kemal Bilici, A. Ýrfan Yükler, A. Kastan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 705
Microstructural analysis of CuAlNiMn shape-memory alloy before and after the tensile testing
Analiza mikrostrukture zlitine CuAlNiMn z oblikovnim spominom pred nateznim preizkusom in po njem
I. Ivani}, M. Goji}, S. Ko`uh, B. Kosec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 713
Effect of the initial microstructure on the properties of low-alloyed steel upon mini-thixoforming
Vpliv za~etne mikrostrukture na lastnosti malolegiranega jekla po predelavi mini-thixoforming
B. Ma{ek, D. Ai{man, H. Jirková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 719
Experimental investigation of the surface properties obtained by cutting Brass-353 (+) with an abrasive water jet and other
cutting methods
Preiskava lastnosti povr{ine medenine 353 (+) po rezanju z abrazijskim vodnim curkom in drugimi metodami rezanja
A. Akkurt . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 725
Influence of the working technology on the development of alloys H13-w(Cu) 87.5 %
Vpliv tehnologije izdelave na razvoj zlitine H13-w(Cu) 87,5 %
U. Arti~ek, M. Bojinovi~, M. Brun~ko, I. An`el . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 735
Microstructure characteristics of the Al-w(Cu) 4.5 % model alloy
Mikrostrukturne zna~ilnosti modelne zlitine Al-w(Cu) 4,5 %
B. [u{tar{i~, M. Jenko, B. [arler . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 743
Chromite spinel formation in steelmaking slags
Nastanek kromitnih spinelov v jeklarskih `lindrah
J. Burja, F. Tehovnik, J. Medved, M. Godec, M. Knap . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 753
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Recycling of jute wastes using pulpzyme enzyme
Recikliranje odpadkov jute z uporabo encima pulpzima
K. Mohajershojaei, F. Dadashian . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 757
Influences of the heat input on a 2205 duplex stainless steel weld
Vpliv vnosa toplote v zvar dupleksnega nerjavnega jekla 2205
B. Rajkumar Gnanasundaram, M. Natarajan. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 761
Tribological behavior and characterization of borided cold-work tool steel
Tribolo{ko vedenje in karakterizacija boriranega orodnega jekla za delo v hladnem
I. Gunes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 765
Microstructures of the Al-Fe-Cu-X alloys prepared at various solidification rates
Mikrostruktura zlitin Al-Fe-Cu-X po razli~nih hitrostih strjevanja
M. Vodìrová, P. Novák, F. Prù{a, D Vojtìch . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 771
Assessment of the post-impact damage propagation in a carbon-fibre composite under cyclic loading
Ocena napredovanja po{kodbe po udarcu pri ponavljajo~ih se obremenitvah kompozita z ogljikovimi vlakni
D. Kytýø, T. Fíla, J. [leichrt, T. Doktor, M. [perl. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 777
Possibilities for increasing the purity of steel in production using secondary-metallurgy equipment
Mo`nosti pove~anja ~istosti jekla pri proizvodnji z uporabo opreme za sekundarno metalurgijo
M. Korbá{, L. ^amek, M. Raclavský . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 781
Effect of different surface-heat-treatment methods on the surface properties of AISI 4140 steel
Vpliv razli~nih toplotnih obdelav povr{ine na lastnosti povr{ine jekla AISI 4140
A. Ayday, M. Durman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 787
Preparation and application of polymer inclusion membranes (PIMs) including alamine 336 for the extraction of metals from
an aqueous solution
Priprava in uporaba membrane iz polimera (PIM) in alamina 336 za lo~enje kovin iz vodnih raztopin
Y. Yildiz, A. Manzak, B. Aydýn, O. Tutkun . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 791
Accelerated carbide spheroidisation and refinement (ASR) of the C45 steel during controlled rolling
Pospe{ena sferoidizacija in udrobnjenje karbidov (ASR) pri kontroliranem valjanju jekla C45
D. Hauserova, J. Dlouhy, Z. Novy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 797
M. HO^EVAR et al.: AN OVERVIEW OF THE INFLUENCE OF STAINLESS-STEEL SURFACE PROPERTIES ...
AN OVERVIEW OF THE INFLUENCE OF
STAINLESS-STEEL SURFACE PROPERTIES ON
BACTERIAL ADHESION
VPLIV LASTNOSTI POVR[INE NERJAVNEGA JEKLA NA
ADHEZIJO BAKTERIJ
Matej Ho~evar1,2, Monika Jenko1, Matja` Godec1, Damjana Drobne3
1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2Jo`ef Stefan International Postgraduate School, Jamova 39, 1000 Ljubljana, Slovenia
3Department of Biology, Biotechnical Faculty, University of Ljubljana, Ve~na pot 111, 1000 Ljubljana, Slovenia
matej.hocevar@imt.si
Prejem rokopisa – received: 2014-02-06; sprejem za objavo – accepted for publication: 2014-06-02
The adhesion and growth of bacteria on the surface of stainless steel promotes corrosion of the material, microbiological
contamination, healthcare problems and results in economic losses. There are numerous factors influencing the adhesion of
bacteria to stainless steel, and material properties are one of the most important ones. In particular, surface roughness,
topography, chemistry and surface energy can promote or inhibit the adhesion and growth of bacteria. Surface roughness and
topography are generally accepted as crucial parameters, especially when the surface features are comparable to the size of the
bacteria. The roughening of the surface increases the area available for adhesion and protects the bacteria from environmental
factors, like liquid shear stress, mechanical forces and disinfectants. The surface chemistry and surface energy of the material
can also affect microbial attachment and survival. The surface chemistry of stainless steel is significantly affected by the
formation of an ultra-thin passive chromium-rich oxide film on the surface in the presence of an oxidative environment. Surface
energy is also an important factor in the initial adhesion and it is commonly known that the minimal relative adhesion to
surfaces occurs at surface energies ranging between 20 mN/m and 30 mN/m (Baier curve). Materials with a high surface energy,
such as stainless steel, are mainly hydrophilic, frequently negatively charged and susceptible to contamination, and thus are
rarely clean. This paper presents an overview of the mechanism and theories of bacterial adhesion on surfaces in general,
together with a comprehensive overview of stainless-steel surface properties that may influence the adhesion of bacteria. Here
we give a literature review and discuss how to manage the stainless-steel surfaces in food processing, medicine and other
industries in order to reduce the adhesion of bacteria.
Keywords: stainless steel, surface properties, adhesion, bacteria
Adhezija in rast bakterij na povr{ini nerjavnega jekla pospe{uje korozijo materiala, povzro~a mikrobiolo{ko kontaminacijo,
zdravstvene te`ave in posledi~no gospodarsko {kodo. Na adhezijo bakterij vplivajo {tevilni dejavniki, med katerimi so lastnosti
materiala med pomembnej{imi. Sem spadajo predvsem hrapavost, topografija, kemijska sestava in povr{inska energija.
Hrapavost in topografija povr{ine sta splo{no sprejeta kot klju~na dejavnika, ki vplivata na adhezijo bakterij na povr{ino,
predvsem kadar so topografske zna~ilnosti na povr{ini primerljive velikosti z bakterijo. Pove~ana hrapavost po eni strani pomeni
ve~jo povr{ino za pritrjevanje, po drugi strani pa {~iti bakterije pred okoljskimi dejavniki, kot so stri`ne napetosti v teko~ih
medijih, mehanske sile in razku`ila. Kemijska sestava povr{ine in povr{inska energija materiala lahko tudi vplivata na adhezijo
in pre`ivetje bakterij. Na kemijske lastnosti povr{ine izrazito vpliva oksidna plast, ki se ustvari na povr{ini nerjavnega jekla.
Povr{inska energija je prav tako pomemben dejavnik pri za~etni fazi adhezije. Splo{no znano je razmerje med adhezijo in
povr{insko energijo, ki ga opisuje Baierjeva krivulja, kjer minimalno adhezijo na povr{ini opazimo pri povr{inski energiji med
20 mN/m in 30 mN/m. Materiali z visoko povr{insko energijo, kot je na primer nerjavno jeklo, so pogosto hidrofilni, negativno
nabiti in dovzetni za kontaminacijo, tako da so te povr{ine redko ~iste. ^lanek daje splo{en pregled mehanizmov in teorije
adhezije bakterij na povr{ine. Podrobno je podan pregled povr{inskih lastnosti nerjavnega jekla, ki lahko vplivajo na adhezijo
bakterij. V ~lanku razpravljamo o tem, kako zmanj{ati adhezijo bakterij v `ivilskopredelovalni industriji ter pri medicinski
uporabi. Za uspe{no zmanj{evanje ne`elene adhezije bakterij je potrebno poglobljeno znanje o dejavnikih, ki v najve~ji meri
vplivajo na adhezijo bakterij.
Klju~ne besede: nerjavno jeklo, lastnosti povr{ine, adhezija, bakterije
1 INTRODUCTION
Bacteria generally exist freely or as a population
attached to surfaces.1 When available, bacteria prefer to
grow on surfaces2 forming aggregates known as biofilms.
Bacterial adhesion and the subsequent biofilm formation
is a complex physico-chemical process consisting of
several stages, including the development of a surface-
conditioning film, the approach of bacteria to the sur-
face, adhesion (initial reversible and subsequent irrever-
sible adhesion), the growth and division of organisms
and finally detachment and dispersal of cells.1,3–5
The adhesion of bacteria to surfaces depends on the
properties of the material (surface topography, rough-
ness, surface chemistry and surface energy), the bacteria
(surface charge, surface hydrophobicity and appendages)
and the surrounding environment (type of medium, tem-
perature, pH, period of exposure and bacterial concen-
tration).6,7 Among them, the material surface properties
are one of the most important.
Stainless steels are commonly used in industrial,
medical and food-processing applications,8,9 and the
adhesion of bacteria to stainless steel represents a chro-
nic source of microbial contamination that leads to the
Materiali in tehnologije / Materials and technology 48 (2014) 5, 609–617 609
UDK 669.14.018.8:620.179.11 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 48(5)609(2014)
deterioration of food, healthcare problems, the enhanced
corrosion of stainless steel and reduces the performance
of plants, pipelines, cooling towers and heat exchan-
gers.1,3,10–12
Stainless steels can be produced in various grades
and finishes, and additional surface treatments can affect
surface physico-chemical properties.13–15 The same type
of stainless steel may have distinctly different surface
properties, including topography, roughness, molecular
composition, electrochemistry and physico-chemistry.15
Additionally, an ultra-thin oxide film composed of chro-
mium and iron oxides forms on the stainless-steel sur-
face, which makes the steel resistant to corrosion.14,15
The surface properties of stainless steel depend on the
stainless-steel grade, the surface finish applied and the
cleaning process used.15 The passive oxide layer is also
very susceptible to contamination from the environment
(dissolve solutes and molecules from air)16,17 and conta-
mination can alter the surface properties and influence
the adhesion.
The influence of stainless-steel surface properties on
the adhesion and retention of bacteria has been exten-
sively investigated. There have been numerous studies on
different stainless steels, including AISI 30218, AISI
3048,13,19–32, AISI 31613,23–26,32–37 and AISI 43013,25 using
different bacteria. However, it is still not clear as to
which characteristics of stainless steels are favourable
for bacterial adhesion as they are often interrelated.15
To reduce microbial adhesion and retention on stain-
less-steel surfaces it is necessary to understand the
factors governing microbial adhesion through the syste-
matic research of the various surface properties involved.
This review summarizes the influence of the surface
properties of stainless steel, especially surface rough-
ness, topography, chemistry and energy on the adhesion
and retention behaviour of bacteria. The aim of this
review is to summarize the available literature data on
the material surface characteristics that are responsible
for bacterial adhesion. We will emphasize the stainless-
steel-bacterial adhesion in order to provide information
about how to produce and maintain the surfaces in order
to reduce bacterial contamination.
2 THEORY AND MECHANISM OF BACTERIAL
ADHESION TO A SURFACE
The adhesion of bacteria to a substrate surface is
governed by the physico-chemical properties of both the
substrate and the bacterium, and also the environmental
conditions.6,7,38 Bacteria may adhere to the surface either
directly to the bare material (nonspecific adhesion) or
indirectly to the conditioning film (specific adhesion) on
the surface. Usually, nonspecific adhesion is investigated
and these results are the closest to the predictions of
theoretical models.39 However, in natural environments
the first step of the adhesion process is the formation of a
conditioning layer4,5 of organic and inorganic molecules
that may alter the physico-chemical properties of the
surface, provide a nutrient source for bacteria or inhibit
the adhesion of certain bacteria.5
2.1 Theoretical background of bacterial adhesion
2.1.1 Physico-chemical Models of Bacterial Adhesion
Concepts developed in colloidal research are a
common approach to predicting bacterial adhesion to
surfaces.2,10 If bacteria are treated as colloids in
suspension, it is possible to model the bacterial adhesion
to surfaces as the sum of the chemical and physical
properties of bacteria and the material surface.39 Three
colloidal models are commonly applied when studying
bacterial adhesion to surfaces: the thermodynamic the-
ory, the Deryaguin–Landau–Verwey–Overbeek (DLVO)
theory and the extended-DLVO (XDLVO) theory.10,39,40
2.1.1.1 Thermodynamic theory
The thermodynamic approach is based on the total
change in the potential Gibbs free energy (energy
available in a closed system) when a bacterium attaches
to a surface and is calculated from the Lifshitz-van der
Waals forces and Lewis acid-base interactions39:
GADH = GLW + GAB (1)
GADH is the total change of the Gibbs free energy of
adhesion, GLW is the Gibbs free energy change of the
Lifshitz-van der Waals forces and GAB is the Gibbs
free energy change of the Lewis acid-base forces.
Thermodynamic theory assumes that adhesion is always
reversible and distance independent. The theory does
not include the effects of surface charge and the
electrolyte concentration of the surrounding media. This
theory is the most accurate with uncharged surfaces or
in the presence of large quantities of ions.39
2.1.1.2 DLVO theory
The DLVO theory like the thermodynamic approach
also assumes that adhesion is the sum of interfacial ener-
gies. However, the DLVO theory considers electrostatic
forces instead of acid-base interactions39:
UDLVO = ULW + UEL (2)
UDLVO is the total interaction energy, ULW is the Lifshitz-
van der Waals interactions energy and UEL is the electro-
static interaction energy. DLVO theory assumes that
adhesion can be reversible and distance dependent. The
theory is most accurate when electrostatic forces are
predominant; however, it is limited due to the dis-
regarded effects of the polar interactions.39
2.1.1.3 XDLVO theory
In an attempt to more accurately model bacterial
adhesion the XDLVO theory combined the features of
the thermodynamic approach and DLVO theory. The
XDLVO model assumes that adhesion is the sum of the
Lifshitz-van der Waals, electrostatic and Lewis acid-base
interactions39:
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610 Materiali in tehnologije / Materials and technology 48 (2014) 5, 609–617
UDLVO = ULW + UEL + UAB (3)
UDLVO is the total interaction energy, ULW is the Lifshitz-
van der Waals interaction energy, UEL is the electrostatic
interaction energy and UAB is the Lewis acid-base
interaction energy. Like with the DLVO, also XDLVO
theory assumes that adhesion can be reversible and
distance dependent.39
All three models favour bacterial adhesion when the
product of the equations’ theories is negative. An
increase or decrease in bacterial adhesion for one set of
parameters compared to a different set of parameters is
calculated. These three theoretic models that predict the
bacterial adhesion to surfaces were developed for ideal
systems; however, the actual bacterial adhesion is
complex and can behave completely differently from the
prediction of the developed models.39
2.2 Mechanism of bacterial adhesion
Actual bacterial adhesion frequently deviates from
the above-described adhesion models.10 Solid materials
exposed to various environments adsorb organic and
inorganic material, thus forming a conditioning layer to
which microorganisms attach.10 The conditioning layer
changes the physico-chemical properties of the surface
and thus plays an important role in the bacterial attach-
ment process.5,10
The adhesion of bacteria to solid surfaces is a
two-phase process composed of an initial reversible
(physical) followed by an irreversible (molecular and
cellular) phase.2,5,41 The adhesion of bacteria to the
surface may be passive or active and this depends on the
motility of the bacteria and the transportation of cells by
gravity, the diffusion of bacteria and the fluid dynamic
forces.3,5 Initial adhesion also depends on the physico-
chemical properties of bacteria cells, their growth phase
and the availability of nutrients.3
The adhesion of bacteria to surfaces occurs rapidly,
within seconds.42,43 Planktonic microbial cells are trans-
ported from the suspension to the conditioned surface
either by bacterial appendages or by physical forces, thus
enabling an initial reversible adhesion.1,3,41 The long-
range physical forces, including van der Waals forces,
steric and electrostatic interactions, influence the initial
reversible adhesion. During this initial stage the bacteria
still show Brownian motion and can be easily removed
from the surface.3,41
After an initial reversible adhesion a number of cells
adhere irreversibly. In this phase molecular reactions
between bacterial surface structures and substrate surfa-
ces become predominant. In contrast to reversible adhe-
sion, various short-range forces such as dipole-dipole
interactions, hydrogen, ionic and covalent bonding and
hydrophobic interactions are involved.1–3,41 Once the bac-
teria attach, irreversible strong physical or chemical
forces are required to remove them from the surface.3
The irreversibly attached bacterial cells start to grow,
divide and form microcolonies, the basic structural unit
of the biofilm.1,3,44 The production of additional extra-
cellular polymeric substances (EPS) helps to strongly
bind the cells to the surface and stabilize the micro-
colonies from the environmental fluctuations1,3 and the
presence of nutrients in the conditioning film and the
surrounding environment determines the rapid growth
and division of cells.1 The biofilm not only enables the
strong attachment of the cells to the surface, but also
helps collect diffuse nutrients, acts as a protection
against environmental stress, antibiotics and disinfec-
tants and enables intercellular communications. As the
biofilm ages the attached bacteria detach and disperse
from the biofilm and colonize new niches.1,3
3 FACTORS AFFECTING BACTERIAL
ADHESION TO SURFACES
Bacterial adhesion is a complex process affected by
the characteristics of the bacteria, environmental
properties and the physical and chemical properties of a
material surface.6,7,38 Environmental factors including the
type of medium, temperature, pH, shear stress of the
flowing medium, bacterial concentration, chemical
treatment and the presence of antibiotics may influence
bacterial adhesion by either changing the surface
characteristics of the bacteria and material or influencing
the interactions in a reversible phase of adhesion. Fur-
thermore, different bacterial species and strains adhere
differently for a given material. This is due to the diffe-
rences in the physic-chemical characteristics of the
bacteria, including surface hydrophobicity, surface
charge, appendages and EPS production.6,7 The physical
and chemical properties of a material surface that can
influence bacterial adhesion to the material surface
include surface roughness, topography or physical confi-
guration, chemical composition, surface energy and
hydrophobicity.6,7 However, the surface characteristics
can be quickly altered by the adsorption of organic and
inorganic compounds forming a conditioning layer.3,5,10
This review will focus on the properties of stainless steel
that can influence bacterial adhesion.
3.1 Stainless-steel surface properties affecting bacte-
rial adhesion
Stainless steels are iron-based alloys containing at
least 10.5 % Cr with numerous alloying elements that
improve the mechanical and corrosion properties.9
Stainless steel is the material of choice in the food-
processing industry14,15,45, mainly because it is inert,
resistant to corrosion, stable at various temperatures and
hygienic.45,46
Stainless steel can be produced in various grades
(AISI 302, AISI 304, AISI 316, AISI 420) and finishes
(2B, 2R, 2D, number 4 finish), thus having different
surface properties (chemistry, topography, roughness,
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Materiali in tehnologije / Materials and technology 48 (2014) 5, 609–617 611
energy).13,14 Furthermore, additional surface treatments
such as mechanical or electro polishing can be applied to
modify the surface topography and roughness and
achieve functionally and aesthetically improved surfa-
ces.13,45 When the commonly used 2B surface finish is
additionally grinded/polished with SiC papers and
diamond paste, different surface patterns are obtained
(Figure 1).47 Stainless steel forms an ultra-thin oxide
film on the surface composed of chromium and iron
oxides that protects the steel from corrosion. The
composition of the oxide film depends on the metal
substrate, the surface finish and the surrounding
environment.15
3.1.1 Effect of surface roughness and topography on
bacterial adhesion
Stainless steel produced in different surface finishes
is designated by a system of standardized numbers: No.
1, 2D, 2B, and 2BA for unpolished finishes; and No. 3,
4, 6, 7, and 8 for polished finishes.9 The 2B pickling
finish, the 2R bright annealed finish and finish 4 are the
most often used.14,15 During production, stainless steel
goes through annealing and pickling processes where the
stainless steel is softened and descaled. These processes
clean the surface of the material prior to processing to a
given finish.14 After cold rolling, which reduces the
thickness of the steel, final annealing (in oxidising
atmosphere) and pickling follows and the surface finish
obtained is designated as a 2D surface finish. When the
2D surface finish is finally light passed on polished rolls
2B or pickling finish is obtained.9,15 To achieve a bright
finish or a 2R finish, the stainless steel is annealed in a
protective atmosphere and the final pickling process is
avoided.15 Finish 4 is achieved when 2D or 2R sheets are
further polished with fine-grained polish belts.9,15 The
surface composition, topography and roughness for a
given material may differ considerably according to the
different surface finishes applied.15
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612 Materiali in tehnologije / Materials and technology 48 (2014) 5, 609–617
Figure 1: Secondary-electron (SE) images taken on scanning electron microscope illustrate the surface features of different surface finishes of
AISI 316L stainless steel: a) 2B surface finish, b) 2B surface finish grinded with 100 SiC grit paper, c) 2B surface finish grinded with 800 SiC
grit paper and d) 2B surface finish polished with diamond paste 3 μm and 1 μm to mirror finish. The 2B surface finish (a) has a network of
subsurface crevices between the oxide grain boundaries; mechanically grinded surface finishes (b, c) exhibit scratch patterns with long linear
alternating grooves and ridges; and the mirror surface finish (d) is the smoothest without pronounced topography features.47
Slika 1: Slike sekundarnih elektronov (SE), posnete z vrsti~nim elektronskim mikroskopom, prikazujejo topografske karakteristike razli~no
obdelanih povr{in nerjavnega jekla AISI 316L: a) 2B povr{ina, b) 2B povr{ina, bru{ena z granulacijo papirja 100 SiC, c) 2B povr{ina, bru{ena z
granulacijo papirja 800 SiC in d) 2B povr{ina, polirana z diamantno pasto 3 μm in 1 μm. 2B povr{ina (a) ima mre`o razpok med oksidnimi zrni
na povr{ini; na mehansko poliranih povr{inah (b, c) je opazen vzorec prask z dolgimi izmeni~nimi dolinami; polirana povr{ina (d) na drugi strani
nima izrazite topografije.47
When studying bacterial adhesion to surfaces, a
comprehensive characterization of the surface-roughness
parameters and visualisation of the surface topography is
very important.38,48 Surface roughness is a two-dimen-
sional parameter of a material surface and is usually
described as the arithmetic average roughness (Ra) and
the root mean square roughness (Rq)6,38, whereas the
topography is a-three dimensional parameter and
describes the shape of the surface features.6 The Ra and
Rq, are commonly reported surface-roughness parameters
when investigating bacterial adhesion; however, they are
measures of the height variation without information
about the topography (surface features).38,48 Therefore, it
is important to measure the spatial or amplitude para-
meters that give information about the spatial variation
and to visualize and describe the morphological features
of the surface.38
Stainless-steel grades AISI 302, AISI 304 and AISI
316 are most often used in adhesion studies due to their
application in the food industry and medicine.14,15,33,49 In
the literature regarding adhesion and retention, bacteria
genus Escherichia, Staphylococcus, Listeria, Pseudo-
monas, Streptococcus and Salmonella are the most often
studied.14 Although a number of studies have investi-
gated the influence of the surface topography and rough-
ness of different stainless steels on the adhesion of
different bacteria the conclusions from these studies are
not consistent.38,48
Several researchers including Jullien et al.13, Ortega
et al.19, Whitehead and Verran23, Flint et al.26, Peterman
et al.27, Hilbert et al.50 and observed no direct correlation
between the surface roughness of the AISI 304/316
stainless Ra ranging between 0.01 μm and 3.3 μm and the
adhesion of bacteria or spores. Arnold et al.8,30 and
Ortega et al.,29 on the other hand, reported a positive
correlation between the adhesion of bacteria and the
surface roughness of the AISI 304 stainless steel. How-
ever, Goulter-Thorsen et al.20 reported that E. coli
attached in greater numbers to significantly smoother
AISI 304 stainless steels. Also interesting are the find-
ings of Medilanski et al.51 who reported that minimal
adhesion occurs at Ra = 0.16 μm and attachment to both
rougher and smother surfaces was significantly higher.
The increased adhesion of bacteria on rougher surfaces
may be explained due to the increase in the surface area
available for adhesion6,7 and the roughening of the sur-
face might also facilitate a firmer attachment by provid-
ing more contact points.52 The opposing observations
reported between the different studies are probably due
to the various experimental conditions, different bacterial
species tested, the material studied and methods used for
bacteria detection.26,42
Besides surface roughness, also the topography of the
surface or surface features such as pits, crevices,
scratches, grooves and ridges, play an important part in
the adhesion process.15,38,43,45 Many researchers con-
cluded that if the surface features are comparable to the
size of the bacteria they can promote bacterial attach-
ment and increase the subsequent microbial reten-
tion.6,7,43,48
The bacteria attach differently to surfaces with
different surface topographies or special surface features
and often the pattern of adhesion reflects the surface
topography (Figure 2).47 Medilanski et al.51 studied the
influence of an AISI 304 stainless-steel surface topo-
graphy (three surface finishes with scratches and two
without observable scratches) on the adhesion of four
bacteria strains (Desulfovibrio desulfuricans, Pseudo-
monas aeruginosa, Pseudomonas putida and Rhodo-
coccus sp.) and found that bacterial cells attach into
scratches in the longitudinal orientation when the width
of the scratches corresponds to the width of the bacterial
cells. Rougher surfaces with wider scratches exhibit a
higher fraction of bacteria adhered in other orientations
and the smoothest surfaces exhibit a random cell orien-
tation.51 Flint et al.26 observed that surface flaws
(scrapes, scratches and pitting) on AISI 304 stainless
steel did not always affect the number of adhered
bacteria; however, bacteria often aligned with the lines
created by the surface flaws. Similar observations were
made by Whitehead and Verran23 on AISI 304 and AISI
316 stainless steel. Barnes et al.21 compared the 2B and 8
mirror finish of AISI 304 stainless steel and reported that
Staphylococcus aureus attach in greater numbers to a
rougher 2B surface finish, whereas little differences
between the 2B and number 8 mirror finish were
observed for Listeria monocytogenes. Furthermore,
scanning electron microscopy revealed that bacteria cells
did not orient exclusively along polishing lines.21 Using
microbial retention assays with a range of differently
sized, unrelated microorganisms on engineered surfaces
(silicon wafers) with controlled topographical features
Whitehead et al.53,54 and Verran et al.55 demonstrated that
the size of the surface features is important with respect
to the size of the bacteria, and its subsequent retention.
The wear of the surfaces may change the adhesion
and retention of bacteria14,45 with the introduction of new
random features (i.e., scratches) different dimensions45,
especially on smooth polished surfaces. Studies of
simulated worn surfaces demonstrated that the hygienic
status of stainless steel was not affected in terms of
microbial retention; however, the cleanability was
affected in terms of the reduced removal of organic
soil.45 Holah and Thorpe46 observed the increased
retention of bacterial cells on abraded sinks compared to
unused ones, this is due to the fact that rougher surfaces
have an increase in the number of attachment sites, a
larger surface contact area and topographical features
that reduce the cleaning shear forces. Verran et al.32
simulated wear on AISI 304 and AISI 316 stainless steel
and studied the retention of Pseudomonas aeruginosa
and Staphylococcus aureus. The results showed that
wear corresponding to Ra < 0.8 μm did not significantly
affect the retention of microorganisms, but the pattern of
M. HO^EVAR et al.: AN OVERVIEW OF THE INFLUENCE OF STAINLESS-STEEL SURFACE PROPERTIES ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 609–617 613
the attachment was highly affected by the surface
topography.32 Linear surface features will be more easily
cleaned along rather than across the features and
presumably also more easily than surfaces with random
linear features across the surface. Furthermore, an in-
crease in the surface roughness may cause the entrap-
ment of microorganisms within the surface features and
reduce the cleanability; however, if the surface features
are significantly larger than the microbial cells, then they
are relatively easily removed from the surface.14 There-
fore, it is important to visualise the surface features as
well as measure the roughness parameters as the wear of
food contact surfaces can affect the topography without
any observable change in roughness.45
3.1.2 Effect of surface chemistry, hydrophobicity and
energy
Stainless steels, produced in various grades and
finishes, also vary in surface properties like chemistry,
topography, roughness and surface energy.14,15 Stainless
steel forms an invisible oxide film (passivation) on the
surface composed of chromium and iron oxides that
protects the steel from corrosion.14,15,45 The composition
of the oxide film depends on the metal substrate, the
surface finish and the surrounding medium.15 When
scratched from surface the oxide layer forms within
seconds and due to the speed of re-passivation it is
difficult to determine the exact chemical composition of
the surface.45 The passive film on a stainless steel is not
static but it changes (grows, dissolves and may adsorb or
incorporate anions) according to the environment.56
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614 Materiali in tehnologije / Materials and technology 48 (2014) 5, 609–617
Figure 2: SE images of attachment patterns of Escherichia coli cells to surfaces with different surface finishes of AISI 316L stainless steel: a) 2B
surface finish, b) 2B surface finish grinded with 100 SiC grit paper, c) 2B surface finish grinded with 800 SiC grit paper and d) 2B surface finish
polished with diamond paste 3 μm and 1 μm to mirror finish. On 2B surface finish (a) microorganisms attach to the crevices between oxide grain
boundaries, whereas on mechanically polished surface finishes (b, c) bacteria align often along longitudinal scratches (when comparable to the
size of the bacteria). On the other hand, mirror finish (d) exhibited a less pronounced topography and microorganisms were observed to be
distributed across the surfaces more randomly.47
Slika 2: SE-slike razporeditve celic bakterije Escherichia coli na razli~no obdelanih povr{inah nerjavnega AISI jekla 316L: a) 2B povr{ina, b) 2B
povr{ina, bru{ena z granulacijo papirja 100 SiC, c) 2B povr{ina, bru{ena z granulacijo papirja 800 SiC in d) 2B povr{ina, polirana z diamantno
pasto 3 μm in 1 μm. Na povr{ini 2B (a) se bakterije pritrjujejo v razpoke med oksidnimi zrni na povr{ini, medtem ko se na mehansko bru{enih
vzorcih (b, c) bakterije pogosto orientirajo vzdol` prask (kadar so primerljivih velikosti z bakterijami). Po drugi strani poliran vzorec nima
izrazite topografije in razporeditev bakterij na povr{ini je naklju~na.47
From a physico-chemical standpoint, the energy charac-
teristics of stainless steel depend on the surface finish
and on the cleaning process used15 and a high- or
low-energy surface can be obtained depending on the
cleaning treatment.15,28
Surface energy is inversely proportional to the
thickness of the contaminating carbon layer that is not
eliminated by cleaning. The cleaning also affects the
surface charge of the steel. For a given surface finish,
and with a pH above the isoelectric point, a more or less
negatively charged surface can be obtained.15 In the food
industry, the electrostatic interactions are repulsive
because stainless-steel surfaces are generally negatively
charged at neutral or alkaline pH and microorganisms are
also negatively charged at these pH values in low-con-
centration aqueous solutions. In weakly charged liquids
such as water, repulsive electrostatic interactions are
significant, whereas in high electrolyte concentrations
(milk, wine) the effect of surface charge is obscured.15
Metals compared to polymers have a high surface ener-
gy, they are mainly hydrophilic, frequently negatively
charged6,15,17 and when exposed adsorb dissolved solutes
or atmospheric contaminants, thus being rarely clean.17
On the other hand, metal oxides provide positively-
charged surfaces that can significantly increase the
adhesion of negatively-charged bacteria to surfaces, pri-
marily due to their positive charge and hydrophobicity.40
It is thought that hydrophobic materials are more sus-
ceptible to bacterial adhesion in contrast to hydro-
philic.6,7 The adhesion of vegetative cells, bacterial
spores and freshwater bacteria has been shown to
increase with increasing surface hydrophobicity. The cell
attachment to hydrophobic plastic occurs very quickly
compared to hydrophilic surfaces (metals oxides, metal
and glass) where longer exposure times are needed.43
Marine Pseudomonas sp. attach in large numbers to
hydrophobic plastics with little or no surface charge,
moderate to hydrophilic metals with a positive or neutral
surface charge and few to hydrophilic, negatively
charged materials such as glass and oxidized plastics.6
Teixeira et al.57 reported that hydrophobic and hydro-
philic bacteria attach in greater numbers to relatively
hydrophobic surfaces with a low surface energy like
AISI 316 and AISI 304 stainless steel compared to
polymethylmethacrylate (PMMA) and glass which are
more hydrophilic. Sinde and Carballo58 studied the adhe-
sion of Salmonella spp. and Listeria monocytogens
strains to AISI 304 stainless steel, rubber and polytetra-
fluorethylene (PTFE). The attachment results showed
that in general Salmonella and Listeria monocytogens
strains adhered in greater numbers to more hydrophobic
material (rubber and PTFE), with stainless steel being
the least hydrophobic.58 Boulangé-Petermann et al.28
studied the wettability of AISI 304 stainless steel with
2B and 2RB surface finishes with respect to the cleaning
process. The cleaning process affected the wettability of
a solid stainless steel surface; however, the results
obtained regarding bacterial adhesion showed no direct
correlation between the wettability or surface energy and
the adhesion of Streptococcus thermophiles.28 Flint et
al.26 studied the adhesion of thermoresistant streptococci
(Streptococcus thermophilus and Streptococcus waiu) to
different substrates (stainless steel, aluminium, zinc,
cooper and glass) and different grades of stainless steel
(AISI 304L and AISI 316L). The influence of substrate
hydrophobicity, charge and a thin oxide film on stainless
steel surfaces was also investigated with respect to the
adhesion of thermo-resistant streptococci.26 The results
showed that bacteria preferentially attach to stainless
steel and zinc compared to copper, aluminium and glass.
Bacteria adhere in higher numbers to AISI 316L
stainless steel with a 2B surface finish compared to AISI
304L stainless steel with the same surface finish, indi-
cating the role of the chemical composition on the
adhesion.26 On the other hand, Percival et al.59,60 reported
a greater number of adhering viable cells on AISI 304
stainless steel compared to grade AISI 316 in a water
piping system over a period of a few months. Flint et al.26
also observed that negatively charged surfaces attracted
more bacteria than positively charged surfaces and the
unpassivated stainless-steel surface (without oxide layer)
reduced the adhesion of thermo-resistant streptococci,
thus suggesting that a stainless-steel surface oxide film
enhances adhesion. However, when stainless steel is
exposed to air repassivation occurs and the ability to
attract bacteria is restored.26
The influence of surface energy on adhesion has been
studied extensively and the surface energy of the
material is an important factor influencing adhesion. The
Baier curve demonstrates the relationship between the
relative adhesion of organisms and the energy of the
surface where minimal adhesion occurs between 20–30
mN/m.61 The surface energy of the substrate also
depends on the conditioning layer (defined by the sur-
rounding environment) and surface structure with surface
irregularities. In an aqueous environment a conditioning
film forms immediately after exposure of the surface and
changes the substrate properties and affects microbial
retention.43
More comprehensive investigations on different
stainless steels with well-known properties (roughness,
topography and surface chemistry) are necessary in order
to determine the effect of surface chemistry on bacterial
adhesion. However, surface physico-chemical properties
are interrelated, therefore, it is difficult to draw the
conclusions of the effects on adhesion due solely to one
of them.15
4 CONCLUSIONS
Bacterial adhesion is governed by properties of the
material surface, bacterial surface characteristics and the
surrounding environment, therefore a comprehensive and
multidisciplinary approach is necessary in order to
improve the understanding of factors contributing to the
adhesion and retention of bacteria to surfaces.
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Materiali in tehnologije / Materials and technology 48 (2014) 5, 609–617 615
The adhesion of bacteria to stainless steel and
retention on surfaces can enhance the corrosion of steel,
present the source of contamination in the food-pro-
cessing industry, cause healthcare problems in medicine
and decreases the performance of equipment in other
industries, thus causing economic losses. Therefore, it is
important to control and reduce the adhesion process.
Stainless-steel surface properties including rough-
ness, topography, chemistry, surface energy and hydro-
phobicity affect the adhesion of bacteria. These factors
are interdependent.
The surface topography and roughness play a crucial
role, especially when they are comparable to the size of
the bacteria and can promote the adhesion and retention
while reducing the cleanability of the surface. On the
other hand, hydrophobicity and surface energy also play
an important role in the adhesion process as hydrophobic
surfaces are more susceptible to adhesion in comparison
to hydrophilic ones and a low surface energy is better
than a high surface energy.
The physico-chemical properties of the substrate are
important in initial cell adhesion; however, once a bio-
film is formed the effect of surface properties on adhe-
sion diminishes, but the effect on retention and cleanabi-
lity is still observable.
In order to reduce or manage the adhesion to stain-
less-steel surfaces in food processing, medical appli-
cation and other industries, knowledge of factors that
govern bacterial adhesion is necessary for each material
being used. It is important to take into account the grade
of the steel, the surface finish applied, the surface
roughness, the cleaning procedures used and the age of
the steel.
A surface-modification approach should concentrate
on a reduction of the initial bacterial adhesion process
and, on the other hand, cleaning protocols used should
be improved, to increase the removal of bacteria. With
wear these protocols should be adjusted (intensified).
Stainless steel is hard, inert, hygienic and has good
wear resistance compared to plastic and ceramics, and
when using smooth surfaces with effective cleaning and
disinfection procedures this is the best approach to
reducing adhesion in food processing, medicine and
industry.
However, we also have to take into account the pro-
perties of different bacteria and surrounding environment
where stainless steel is exposed. Therefore, the adhesion
of a particular group of bacteria that are expected to
contaminate the surfaces should be tested.
Acknowledgment
The work was conducted at the Institute of Metals
and Technology in the frame of research programme
Surface physics and chemistry of metallic materials
(P2-0132).
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Materiali in tehnologije / Materials and technology 48 (2014) 5, 609–617 617

A. [TRKALJ et al.: ADSORPTION OF HEXAVALENT CHROMIUM FROM AN AQUEOUS SOLUTION ...
ADSORPTION OF HEXAVALENT CHROMIUM FROM AN
AQUEOUS SOLUTION OF STEEL-MAKING SLAG
ADSORPCIJA HEKSAVALENTNEGA KROMA IZ VODNE
RAZTOPINE JEKLARSKE @LINDRE
Anita [trkalj1, Zoran Glava{1, Gordana Matija{i}2
1University of Zagreb, Faculty of Metallurgy, Aleja narodnih heroja 3, 44 000 Sisak, Croatia
2University of Zagreb, Faculty of Chemical Engineering and Technology, Maruli}ev trg 19, 10 000 Zagreb, Croatia
strkalj@simet.hr
Prejem rokopisa – received: 2013-01-15; sprejem za objavo – accepted for publication: 2013-11-18
A batch removal of Cr(VI) ions from an aqueous solution under different experimental conditions using steel-making slag as a
low-cost adsorbent is presented in this paper. The obtained results showed that the steel-making slag is an effective adsorbent for
the removal of Cr(VI) ions from aqueous solutions. The adsorption of Cr(VI) with the steel-making slag follows the Langmuir
isotherm equation. Among the tested kinetics models in this study (pseudo-first-order, pseudo-second-order, Elovich and
intraparticle-diffusion models), the pseudo-second-order equation successfully predicted the adsorption. The thermodynamic
parameters for the adsorption process were determined and discussed.
Keywords: adsorption, Cr(VI) ions, steel-making slag, isotherms, kinetic, thermodynamic
V ~lanku je predstavljena raziskava odstranjevanja ionov Cr(VI) iz vodne raztopine z jeklarsko `lindro kot poceni adsorbenta pri
razli~nih eksperimentalnih razmerah. Dobljeni rezultati so pokazali, da je jeklarska `lindra u~inkovit adsorbent za odstranje-
vanje Cr(VI) ionov iz vodnih raztopin. Adsorpcija ionov Cr(VI) z jeklarsko `lindro se sklada z Langmuirovo izotermno ena~bo.
V tej {tudiji je od preizku{enih kineti~nih modelov (psevdo prvega reda, psevdo drugega reda, Elovichev model in model
difuzije med delci) ena~ba psevdo drugega reda uspe{no napovedala adsorpcijo. Dolo~eni in obravnavani so termodinami~ni
parametri za proces adsorpcije.
Klju~ne besede: adsorpcija, ioni Cr(VI), jeklarska `lindra, izoterme, kinetika, termodinamika
1 INTRODUCTION
Over the last few decades, due to their increased use
in the treatment of metals and ceramic, glass production,
mining operations and the production of batteries,
various heavy metals were released into terrestrial and
aquatic ecosystems.1 The increase in the environmental
contamination with heavy metals is a big concern for
ecological systems and human health due to their
toxicity, accumulation in food and persistence in nature.2
The main focus in water and wastewater treatment is
given to hexavalent chromium due to its carcinogenic
properties.3 Cr(VI) ions are considered as one of the top
16 toxic pollutants and due to their carcinogenic effect
on humans, they have become a serious health problem.4
Cr(VI) ions can be released into the environment from
various industrial operations such as the treatment of
metals, production of iron and steel and the production
of inorganic chemicals.5
Heavy metals can be removed from aqueous solu-
tions using various techniques such as ion exchange,6
precipitation7 and adsorption.8 Adsorption has been
successfully used to remove heavy metals.9 For several
decades, activated carbon has been used as an adsorbent
for the purification of industrial wastewater.10 Although
it is a most appropriate adsorbent for removing heavy
metals, its widespread use was limited due to its high
cost. Steel-making slag as an alternative adsorbent has
been used to remove heavy metals in the environmental
field. Due to its unique properties, steel-making slag is
used as an alternative adsorbent for removing heavy
metals from aqueous solutions. Steel-making slag is a
by-product of the steel production and a waste material
that is widely used due to its useful properties.11
2 EXPERIMENTAL WORK
Adsorption experiments were performed using the
batch-equilibration technique. The initial concentrations
of Cr(VI) ions were prepared in the range of 50–300
mg/L of dissolving K2Cr2O7 in distilled water. A series of
Erlenmeyer flasks containing steel-making-slag samples
1 g and solutions 50 mL were sealed until equilibrium
was obtained. Then the adsorbent was removed by fil-
tration. The concentration of Cr(VI) ions was determined
using an atomic absorption spectrometer with a graphite
furnace, equipped with Zeeman background correction.
A Cr hollow-cathode lamp operating at the current of
4 mA was used as a line source. The measurements were
performed at 357.9 nm, with the slit fixed at 0.8 nm. The
atomization was carried out with the following
parameters: T = 2400 °C, the ramp rate of 1200 °C/s and
the dwell time of 6 s.12 Each analysis was performed in
triplicate and the average value was taken as the result.
The chemical composition of the steel-making slag was
determined with the standard chemical analysis13 and the
result in mass fractions (w) is: CaO – 36 %, MgO –
Materiali in tehnologije / Materials and technology 48 (2014) 5, 619–622 619
UDK 544.726:544.3:544.4 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)619(2014)
0.3 %, MnO – 18 %, SiO2 – 17 %, Al2O3 – 1.4 %, FeO –
27 %.
3 RESULTS AND DISCUSSION
3.1 Adsorption isotherms
Figure 1 shows the adsorption isotherms of Cr(VI)
ions on the steel-making slag. The adsorption of Cr(VI)
ions on the steel-making slag increased with an increase
in the ion concentration. This finding is similar to the
other studies.11,14
The adsorption-equilibrium data were processed
using two isotherms: the Freundlich and Langmuir iso-
therms. The Langmuir isotherm (Equation 1) and the
Freundlich isotherm (Equation 2) can be expressed as:7,8
q
q K c
K ce
L e
L e
=
+
max
1
(1)
q K c ne F e=
1 / (2)
where ce is the equilibrium concentration of Cr(VI) ions,
qmax (mg/g) and qe (mg/g) are the maximum adsorption
capacity and the adsorbed amount of Cr(VI) ions. KF
(mg/g) and n are the Freundlich constants. KF (L/mg) is
the adsorption capacity of the adsorbent and n is the
intensity of adsorption. KL is the Langmuir constant
related to the energy of adsorption. The parameters of
the two isotherms calculated with Equations (1) and (2)
are presented in Table 1 and the fitting curves for the
experimental data are shown in Figures 2 and 3.
Table 1: Parameters of the Langmuir and Freundlich adsorption-iso-
therm models
Tabela 1: Parametri Langmuirovega in Freundlichovega modela
adsorpcijskih izoterm
Langmuir Freundlich
T
K
qmax
mg/g
KL
L/mg
r2 n KF
mg/g
r2
293
313
333
3.14
3.05
3.59
0.020
0.034
0.029
0.9580
0.9765
0.9912
2.24
2.51
2.80
0.251
0.366
0.930
0.8678
0.9327
0.9172
A comparison of the correlation coefficients and the
fitting curves obtained using the two models shows that
the Langmuir model was more suitable for the adsorp-
tion of Cr(VI) ions from aqueous solutions by the
steel-making slag. The Langmuir theory considers the
adsorption onto the materials with homogeneous specific
surfaces. Therefore, the maximum adsorption capacity
(qmax) can be obtained from the fittings of the adsorption
isotherm.9
The main feature of the Langmuir adsorption iso-
therm is a dimensionless constant called the separation
factor or equilibrium parameter (RL), presented by the
following equation:15,16
R
K cL
=
+
1
1 L 0
(3)
where c0 is the initial concentration of Cr(VI) ions
(mg/L). The RL value indicates the shape of the isotherm
to be irreversible (RL = 0), favorable (0 < RL < 1), linear
(RL = 1) or unfavorable (RL > 1).17,18 The RL value for the
studied Cr(VI) ions/steel-making slag system varied
from 0.17 to 0.63. It means that the steel-making slag is
A. [TRKALJ et al.: ADSORPTION OF HEXAVALENT CHROMIUM FROM AN AQUEOUS SOLUTION ...
620 Materiali in tehnologije / Materials and technology 48 (2014) 5, 619–622
Figure 3: Freundlich isotherms for the adsorption of Cr(VI) ions by
steel-making slag
Slika 3: Freundlichove izoterme za adsorpcijo ionov Cr(VI) na
jeklarsko `lindro
Figure 2: Langmuir isotherms for the adsorption of Cr(VI) ions by
steel-making slag
Slika 2: Langmuirove izoterme adsorpcije ionov Cr(VI) na jeklarsko
`lindro
Figure 1: Adsorption isotherms of system steel-making slag – Cr(VI)
ions
Slika 1: Adsorpcijske izoterme sistema jeklarska `lindra – ioni Cr(VI)
favorable for the adsorption of Cr(VI) ions from
aqueous solutions under the conditions used in this
study.
For the Freundlich isotherm, as shown in Table 1, n
is equal to 2.24. In many cases n > 1; this may be a result
of the distribution of surface sites or another factor
causing a decrease in the adsorbent/adsorbate interaction
due to an increase in the surface density.19 The values of
n in the range from 2 to 10 indicate a good adsorption.20
3.2 Adsorption dynamics
The study of adsorption dynamics describes the rate
of solute removal. This rate controls the residence time
of the adsorbate uptake at the solid/solution interface.
The kinetics of the Cr(VI) ion adsorption on steel-
making slag was analyzed using pseudo-first-order,
pseudo-second-order, Elovich and intraparticle-diffu-
sion-kinetic models21–24 (diagram not shown here). The
correlation coefficient (r2) indicates the fitting experi-
mental data and the values predicted by the model. A
high value of r2 (close or equal to 1) indicates that the
model successfully describes the kinetics of the
adsorption of Cr(VI) ions.
The pseudo-first-order equation is generally ex-
pressed as follows:15
d
d e t
q
t
k q q
t = −1 ( ) (4)
where qe and qt are the adsorption capacities at equili-
brium and at time t, respectively (mg/g), and k1 is the
rate constant of the pseudo-first-order adsorption
(L/min).
The pseudo-second-order adsorption-kinetic-rate
equation is expressed as:15
d
d e t
q
t
k q q
t = −2
2( ) (5)
where k2 is the rate constant of the pseudo-second-order
adsorption (g/(mg min)).
The Elovich model equation is generally expressed
as:15
d
d t
q
t
q
t = − exp( ) (6)
where  is the initial adsorption rate (mg/(g min)) and 
is the desorption constant (g/mg) during the experiment.
The intraparticle-diffusion model is expressed as:15
R k t a= id ( ) (7)
where R is the percent of the Cr(VI) ions adsorbed, a is
the gradient of linear plots and kid is the intraparti-
cle-diffusion-rate constant (h–1).
Figure 4 shows the results of the adsorption capacity
of Cr(VI) ions on the steel-making slag versus time at
293 K. The parameters of adsorption-kinetic models are
presented in Table 2.
The kinetic adsorption of Cr(VI) ions on the steel-
making slag followed the pseudo-second-order model
(the correlation coefficient (r2) was the highest –
0.9999). The r2 values indicate that the intraparticle-
diffusion process is the rate-limiting step. Higher values
of kid indicate a better mechanism of adsorption and an
increase in the adsorption rate, which is related to an
improved bonding between Cr(VI) ions and the adsor-
bent particles.24
3.3 Thermodynamic studies
It is necessary to perform thermodynamic studies of
the adsorption process to conclude whether the process is
spontaneous or not. The Gibbs free energy change (G°)
is an important parameter determining the spontaneity of
the chemical reaction. In order to determine the Gibbs
free energy of the process it is necessary to know the
entropy and enthalpy. The reaction occurs spontaneously
at a given temperature if G° has a negative value. The
Gibbs free energy change (G°), the enthalpy change
(H°) and the entropy change (S°) were calculated
from a variation of the thermodynamic equilibrium
Langmuir constant, KL. The thermodynamic parameters
were calculated using the following equations:25
ΔG RT K° = − ln L (8)
ln K
S
R
H
RTL
=
°
−
°Δ Δ
(9)
A. [TRKALJ et al.: ADSORPTION OF HEXAVALENT CHROMIUM FROM AN AQUEOUS SOLUTION ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 619–622 621
Figure 4: Adsorption capacity of Cr(VI) ions on steel-making slag
versus time at 293 K
Slika 4: ^asovna odvisnost kapacitete adsorpcije ionov Cr(VI) na
jeklarsko `lindro pri 293 K
Table 2: Constants of the adsorption kinetic models (system Cr(VI) ions – steel-making slag)
Tabela 2: Konstante kineti~nih modelov adsorpcije (sistem Cr(VI) ioni – jeklarska `lindra)
Pseudo first order Pseudo second order Elovich model Intraparticle-diffusion model
k1
L/min
r2 k2
g/(mg min)
r2 
g/mg

mg/(g min)
r2 kid
h–1
a r2
0.121 0.9204 2.777 0.9999 1.560 2.615 0.9061 17.227 0.340 0.8200
where H° and S° were determined from the slope
and the intercept of the plot of lnKL versus 1/T (diagram
not shown here). Thermodynamic parameters are given
in Table 3.
Table 3: Thermodynamic parameters of the adsorption of Cr(VI) ions
on steel-making slag
Tabela 3: Termodinami~ni parametri adsorpcije ionov Cr(VI) na
jeklarsko `lindro
T
K
G°,
J mol–1
H°,
kJ mol–1
S°,
J mol–1 K–1
293
313
333
–9.529 –3.397 –2.159
A negative value of G° (Table 3) indicates that the
adsorption is highly favorable and spontaneous. A nega-
tive value of H° indicates that the adsorption is exo-
thermic. The adsorption in the solid/liquid system
consists of two processes: the adsorption of the adsorbate
(solute) and desorption of the solvent (water) molecules
that have been previously adsorbed. In an endothermic
process, to be adsorbed, the adsorbate particles have to
displace more than one water molecule. This results in
the endothermic reaction of the adsorption process. On
the order hand, in an exothermic process, the total energy
consumed for bond breaking is lower than the total
energy released during the formation of the bond
between an adsorbate and an adsorbent. This results in
the release of the extra energy in the form of heat.
Therefore, H° will be negative. The value of H° also
indicates the type of adsorption. The heat produced
during the physical adsorption is the same as the heats of
condensation, i.e., 2.1–20.9 kJ mol–1.26 On the order
hand, the heat of chemisorption generally falls into the
range of 80–200 kJ mol–1.26 Accordingly, the data in
Table 3 show that the adsorption of Cr(VI) ions can be
attributed to the physical-adsorption process. A negative
value of S° indicates that the adsorption process is
enthalpy driven. A negative value of the entropy change
(S°) also indicates a decreased disorder at the solid/
liquid interface during the adsorption process causing the
adsorbate ions/molecules to escape from the solid phase
to the liquid phase.26 It this case, there is a decrease in
the amount of the adsorbate that can be adsorbed.27
4 CONCLUSIONS
The present study shows that steel-making slag is an
effective adsorbent for a Cr(VI) ion removal from an
aqueous solution. The equilibrium studies confirmed that
the Langmuir model was better in describing the adsorp-
tion of Cr(VI) ions on steel-making slag. The dimension-
less-separation factor (RL) showed that steel-making slag
could be used for the removal of Cr(VI) ions from an
aqueous solution. The amount of the adsorbed Cr(VI)
ions increased with an increase in the temperature. The
kinetics of the adsorption of Cr(VI) ions on steel-making
slag followed a pseudo-second-order model. The r2
values indicate that the intraparticle-diffusion process is
the rate-limiting step. The negative value of G° indi-
cates that the adsorption is highly favorable and sponta-
neous. A negative value of H° indicates that the adsorp-
tion is exothermic. A negative value of S° indicates that
the adsorption process is enthalpy driven.
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A. [TRKALJ et al.: ADSORPTION OF HEXAVALENT CHROMIUM FROM AN AQUEOUS SOLUTION ...
622 Materiali in tehnologije / Materials and technology 48 (2014) 5, 619–622
J. KUBÁSEK et al.: STRUCTURAL, MECHANICAL AND CYTOTOXICITY CHARACTERIZATION ...
STRUCTURAL, MECHANICAL AND CYTOTOXICITY
CHARACTERIZATION OF AS-CAST BIODEGRADABLE
Zn-xMg (x = 0.8–8.3 %) ALLOYS
STRUKTURNE, MEHANSKE IN CITOTOKSI^NE LASTNOSTI
BIORAZGRADLJIVE Zn-xMg (x = 0,8–8,3 %) ZLITINE V LITEM
STANJU
Jiøí Kubásek1, Iva Pospí{ilová1, Dalibor Vojtìch1, Eva Jablonská2, Tomá{ Ruml2
1Department of Metals and Corrosion Engineering, Institute of Chemical Technology, Prague, Technická 5, 166 28 Prague 6, Czech Republic
2Department of Microbiology, Institute of Chemical Technology, Prague, Technická 5, 166 28 Prague 6, Czech Republic
kubasek.jiri@gmail.com
Prejem rokopisa – received: 2013-04-05; sprejem za objavo – accepted for publication: 2013-11-04
In the present work, the structural, tensile, compressive and bending mechanical properties as well as the cytotoxicity of the
Zn-Mg biodegradable alloys containing mass fractions from 0 % up to 8.3 % Mg were studied. It was found that the maximum
tensile and compressive strengths of 170 MPa and 320 MPa, respectively, were obtained for the Zn-0.8Mg alloy. This alloy also
showed the highest tensile elongation of 2 %. Mechanical properties were discussed in relation to the various structural features
of the alloys. The structure of the strongest Zn-0.8Mg alloy was composed of a fine mixture of -Zn dendrites and -Zn +
Mg2Zn11 eutectics. The cytotoxicity was evaluated with an indirect contact assay using human osteosarcoma cells (U-2 OS). The
cytotoxicity of the Zn-0.8Mg alloy extract was low and only slightly higher than in the case of the pure-Mg extract.
Keywords: biodegradable material, zinc, mechanical properties, structure, cytotoxicity
V tem delu so bile preu~evane struktura, natezne, tla~ne in upogibne mehanske lastnosti biorazgradljive zlitine Zn-Mg z masnim
dele`em od 0 % do 8,3 % Mg. Ugotovljeno je bilo, da sta bili pri zlitini Zn-0,8Mg dose`eni najve~ja natezna in tla~na trdnost
170 MPa oziroma 320 MPa. Ta zlitina je tudi pokazala najve~ji raztezek 2 % pri natezni obremenitvi. Mehanske lastnosti so
prikazane v odvisnosti od mikrostrukturnih zna~ilnosti zlitin. Mikrostruktura najmo~nej{e zlitine Zn-0,8Mg je bila sestavljena iz
-Zn-dendritov in -Zn + Mg2Zn11-evtektika. Citotoksi~nost je bila ocenjena s posrednim preizkusom kontakta s celicami
~love{kega osteosarkoma (U-2 OS). Citotoksi~nost ekstrakta zlitine Zn-0,8Mg je bila majhna in samo malo ve~ja v primerjavi z
ekstraktom ~istega Mg.
Klju~ne besede: biorazgradljiv material, cink, mehanske lastnosti, struktura, citotoksi~nost
1 INTRODUCTION
Biodegradable implant materials progressively de-
grade in the human body after the implantation, produc-
ing relatively non-toxic compounds and, simultaneously,
they are being replaced by the growing tissue. Biodegra-
dable polymeric materials have been known and used for
a long time. However, polymeric materials are not suit-
able for the load-bearing applications such as screws or
plates for fractured bones, due to their low mechanical
strength.1 Among biodegradable metallic materials,
magnesium alloys have attracted the greatest interest
since the beginning of the 20th century.2 The reason for
this is that magnesium is relatively non-toxic to the
human body and excessive amounts of it can be readily
excreted by the kidneys. Moreover, it is very important
for many biological functions of the human body. The
main disadvantage of most magnesium alloys explored
so far is that they corrode too rapidly in physiological
environments, producing excessive amounts of hydrogen
and increasing the alkalinity close to the implant.2–6 Both
factors retard the healing process. Therefore, large
efforts have been devoted to finding magnesium-based
alloys degrading at acceptably low rates and many kinds
of Mg-based alloys, like AZ, AM, LAE, WE, Mg-Zn,
Mg-Zn-Ca, Mg-Zn-Mn-Ca, Mg-Zn-Y, Mg-Gd, Mg-Zn-
Si and others, have been studied until now.7–16
In Mg-based biodegradable alloys, zinc is often one
of the major constituents. Zinc is known to improve the
strength and corrosion resistance of magnesium and,
from the biological point of view, it is generally consi-
dered relatively non-toxic.17 In the majority of the Mg
biodegradable alloys given above, the concentrations of
zinc do not exceed several mass fractions w/%. But a
deep eutectic in the Mg-Zn binary-phase diagram18 with
about 51 % Zn supports the glass-forming ability (GFA)
of Mg-Zn-based alloys with high concentrations of zinc.
Amorphous ternary Mg-Zn-Ca alloys whose composi-
tions are close to the eutectic point were already pre-
pared19,20 and shown to be promising candidates for
biodegradable implants. Due to the unordered atomic
structures and high Zn contents, amorphous alloys have
an excellent strength, high corrosion resistance, low
hydrogen evolution rate and a good biocompatibility in
animals. Until now, bulk amorphous Mg-Zn-Ca samples
of only a few millimeters in size have been prepared.
The fact that zinc is a biologically tolerable element,
even when its content in Mg-based alloys approaches
Materiali in tehnologije / Materials and technology 48 (2014) 5, 623–629 623
UDK 669.55:620.17 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)623(2014)
w = 50 %,21 indicates that Zn-based alloys may also be
promising candidates for biodegradable implants. This
has motivated our research of Zn-based biodegradable
alloys. The main advantage of these materials over
Mg-Zn metallic glasses is that they are much easier to
prepare using the classical routes like gravity or die cast-
ing, hot rolling, hot extrusion, ECAP, etc. In our previous
work22 we provided information on the basic mechanical
and corrosion properties of three binary Zn-Mg alloys
containing w(Mg) = 1–3 %. We have shown that these
alloys are significantly more corrosion resistant in a
simulated body fluid (SBF) than Mg alloys. Possible
zinc doses and toxicity were estimated and found to be
negligible when compared to the tolerable biological
daily limit of zinc. Magnesium was shown to strengthen
the alloys and it is also known to support the healing
process of the hard tissue.22
To properly design a load-bearing implant, the me-
chanical properties of an implant material are of a great
importance. For this reason, this work is focused on the
mechanical characterization including the hardness,
tensile, compressive and bending properties of a series of
the Zn-Mg binary alloys containing w(Mg) = 0–8 %.
These limits were selected to be around the eutectic point
in the Zn-Mg system (approximately w(Mg) = 3 %).18
Moreover, the cytotoxicity of zinc was also assessed and
compared with magnesium and other elements.
2 EXPERIMENTAL WORK
In this study, pure zinc and six binary Zn-Mg alloys
containing w = 0.8–8.3 % of Mg were studied. The
designations and chemical compositions of the studied
alloys are given in Table 1.
Zn-based alloys were prepared by melting pure zinc
(99.95 %) and magnesium (99.90 %) in a resistance
furnace in air. To prevent an excessive evaporation of
volatile zinc, the melting temperature did not exceed
500 °C, and the homogenization was ensured with an
intense mechanical stirring using a graphite rod. After a
sufficient homogenization, the melts were poured into a
cast-iron mold to prepare cylindrical ingots of 20 mm in
diameter and 130 mm in length. The chemical compo-
sitions were verified at both ends of the ingots with
X-ray fluorescence spectrometry (XRF), as shown in
Table 1.
The structures of the alloys were studied using light
(LM) and scanning electron microscopy (SEM, Tescan
Vega 3) with energy dispersive X-ray spectrometry (EDS,
Oxford Instruments Inca 350). The phase composition
was also confirmed with an X-ray diffraction analysis
(XRD, X Pert Pro).
The mechanical properties of the as-cast alloys were
characterized with hardness, tensile, compressive and
bending tests. The samples for these tests were cut di-
rectly from the as-cast ingots. A loading of 5 kg was
used for Vickers-hardness measurements. The rod
samples for tensile tests had a diameter and length of 10
mm and 120 mm, respectively. Compressive tests were
realized with rectangular samples of 10 mm × 10 mm ×
15 mm in size. The compressive loading direction was
parallel to the longest dimension. For three-point bend-
ing tests, rod samples of 4 mm in diameter and 40 mm in
length were used. All the mechanical tests used a defor-
mation rate of 1 mm/min. Fracture surfaces were exa-
mined after the tensile and bending tests using SEM.
The cytotoxicity of Zn-Mg alloys was assessed by
using human osteosarcoma cells (U-2 OS). The investi-
gated Zn-0.8Mg alloy and pure magnesium were used in
these tests. Pure magnesium was used for comparison
because it is generally considered to have a good bio-
compatibility and it is, thus, the basis of the most exten-
sively studied metallic biomaterials. Before the extrac-
tion, cells were cultured in Dulbecco’s modified Eagle’s
medium (DMEM) with a 10 % fetal bovine serum
(FBS), 100 U/mL penicillin, 1 mg/mL streptomycin and
250 ng/mL amphotericin B at 37 °C in a humidified
atmosphere of 5 % CO2. The cytotoxicity was evaluated
with an indirect-contact assay. Extracts were prepared by
immersing the alloys in a DMEM medium containing a
5 % FBS and antibiotics at 37 °C for 7 d. The ratio of the
surface area of the alloy samples to the extraction me-
dium was 1 cm2/mL. The extracts were then withdrawn
and diluted twice. The concentrations of Zn and Mg
released from Zn-0.8Mg and Mg, respectively, were
determined using inductively coupled plasma mass
spectrometry (ICP MS). The cells were seeded at a
density of cells 2.5 × 104 mL–1 and incubated in 96-well
cell culture plates for 24 h to allow attachment. The
medium was then replaced with 100 μL of the extracts.
The controls for a comparison of the cell viability
involved both pure DMEM medium (100 % viability)
and 0.64 % phenol in the DMEM medium as the toxic
control. After the incubation of the cells in a humidified
atmosphere of 5 % CO2 at 37 °C for 1 d the extracts were
removed. The cells were then washed twice using
phosphate buffered saline (PBS) and overlaid with a
phenol red-free DMEM medium containing 5 μL of the
WST-1 reagent (Roche) per well. The plates were
incubated with WST-1 for 4 h at 37 °C. The assay is
based on the reduction of tetrazolium salt to soluble
formazan due to mitochondrial enzymes of the viable
J. KUBÁSEK et al.: STRUCTURAL, MECHANICAL AND CYTOTOXICITY CHARACTERIZATION ...
624 Materiali in tehnologije / Materials and technology 48 (2014) 5, 623–629
Table 1: Chemical compositions of the investigated alloys in mass fractions (w/%)
Tabela 1: Kemijske sestave preiskovanih zlitin v masnih dele`ih (w/%)
Alloy designation Zn Zn-0.8Mg Zn-1.6Mg Zn-2.5Mg Zn-3.5Mg Zn-5.4Mg Zn-8.3Mg
Mg concentration <0.01 0.79 ± 0.05 1.57 ± 0.04 2.51 ± 0.03 3.47 ± 0.03 5.36 ± 0.12 8.32 ± 0.04
cells. The absorbance of the samples characterizing the
cell viability was measured using a microplate reader at
450 nm with a reference wavelength of 630 nm. The
higher the absorbance, the higher is the viability of the
cells.
3 RESULTS AND DISCUSSION
3.1 Structures
The light micrographs of the investigated alloys are
shown in Figure 1. It is seen in Figure 1a that the pure
zinc is composed of almost equi-axed grains of approxi-
mately 20 μm in size. The structures of the alloys from
Zn-0.8Mg to Zn-2.5Mg (Figures 1b to 1d) are hypo-
eutectic, i.e., they are composed of the primary -Zn
dendrites (light) and the -Zn + Mg2Zn11 eutectic
mixture (dark) in interdendritic regions, dominated by
the lamellar and rod morphologies of the eutectic phases
(see also a detailed view inserted in Figure 1d). The
presence of the two phases of -Zn and Mg2Zn11 in the
alloys was also confirmed with XRD (not shown). The
average thickness of the dendritic branches in these
alloys is approximately 30 μm and the volume fraction of
the eutectic mixture increases with the increasing Mg
concentration. In all the hypoeutectic alloys the average
eutectic interparticle spacing is approximately 2 μm
(Figure 1d). The composition of the Zn-3.5Mg alloy is
very close to the eutectic point in the binary Zn-Mg
phase diagram.18 Its structure (Figure 1e) is, thus, domi-
nated by a very fine rod-and-lamellar -Zn + Mg2Zn11
eutectic mixture, in which the average eutectic interpar-
ticle spacing is close to that in the previous alloys (2
μm). The eutectic mixture creates the colonies of 50–200
μm in size that differ in the spatial orientation of the rods
and lamellae. The Zn-5.4Mg and Zn-8.3Mg alloys (Fig-
ures 1f and 1g, respectively) are hypereutectic, contain-
ing sharp-edged Mg2Zn11 intermetallic phases (light) and
the -Zn+Mg2Zn11 eutectic mixture (dark). Rod and
lamellar eutectic particles are observed (Figure 1g). Like
in the hypoeutectic and eutectic alloys, the average
eutectic interparticle spacing is approximately 2 μm. The
volume fraction and dimensions of the primary inter-
J. KUBÁSEK et al.: STRUCTURAL, MECHANICAL AND CYTOTOXICITY CHARACTERIZATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 623–629 625
Figure 1: Light and detailed SEM micrographs of the alloys: a) Zn, b) Zn-0.8Mg, c) Zn-1.6Mg, d) Zn-2.5Mg, e) Zn-3.5Mg, f) Zn-5.4Mg, g)
Zn-8.3Mg
Slika 1: Svetlobni in podrobni SEM-posnetki mikrostrukture zlitin: a) Zn, b) Zn-0,8Mg, c) Zn-1,6Mg, d) Zn-2,5Mg, e) Zn-3,5Mg, f) Zn-5,4Mg,
g) Zn-8,3Mg
metallic phases increase with the increasing Mg contents
in the alloys. The nature of the Mg2Zn11 intermetallic
phases was verified with XRD (not shown) and EDS,
determining mole fractions x = 15.1 % Mg and x = 84.9
% Zn in these particles. Due to enrichment in
magnesium, the particles are often surrounded by a thin
layer of the -Zn phase.
3.2 Mechanical properties
Figure 2 shows various mechanical properties of the
Zn-Mg alloys as functions of the Mg content. One can
see that there is a direct relationship between the Mg
content and mechanical properties. The hardness of the
Zn-Mg alloys increases with the Mg concentration from
approximately 37 HV5 for the pure zinc up to 226 HV5
for the Zn-8.3Mg alloy (Figure 2a). This behavior can
be attributed to the increasing volume fraction of the
hard Mg2Zn11 intermetallic phase due to magnesium
(Figure 1).
The compressive mechanical properties summarized
in Figure 2b show a similar trend, i.e., the compressive
yield strength increases with the Mg concentration from
80 MPa (the pure zinc) up to 625 MPa (the Zn-3.5Mg
alloy). The ultimate compressive strength of the alloys
containing 0–3.5 % Mg was not measured because they
were not broken during the loading, suggesting a good
compressive plasticity of hypoeutectic and eutectic
alloys. In the case of hypoeutectic alloys this plasticity is
attributable to a relatively large volume fraction of the
soft -Zn phase (Figures 1a to 1d). The compressive
plasticity of the eutectic Zn-3.5Mg alloy is a little sur-
prising because this alloy contains approximately
volume fraction  = 50 % of the brittle Mg2Zn11 eutectic
phase (Figure 1e). A detailed insert in Figure 1e shows
that the eutectic mixture is very fine and that the average
diameter of eutectic rods and the thickness of eutectic
lamellae do not significantly exceed 1 μm. During the
compressive loading, hard eutectic particles act as stress
concentrators. The larger are the particles, the higher is
the local-stress increase around them. In the case of fine
eutectic particles, the local-stress increase in their vici-
nity is probably small; therefore, the alloy shows an
unlimited plastic deformation in compression. In con-
trast, at higher Mg concentrations the fracture took place
before the onset of the plastic deformation, therefore,
only the values of the ultimate compressive strength of
the Zn-5.4Mg and Zn-8.3Mg alloys are shown in Figure
2b. Both these alloys contain sharp-edged primary
Mg2Zn11 intermetallic phases (Figures 1f to 1g). Their
sizes, shapes and brittle nature indicate that the stress
concentration around them is high. Fracture cracks, thus,
grow at a low nominal compressive stress of slightly
above 200 MPa, i.e., significantly below the onset of the
plastic deformation.
Bending mechanical properties are illustrated in
Figure 2c. It is observed that a bending strength first
J. KUBÁSEK et al.: STRUCTURAL, MECHANICAL AND CYTOTOXICITY CHARACTERIZATION ...
626 Materiali in tehnologije / Materials and technology 48 (2014) 5, 623–629
Figure 2: Mechanical properties of Zn-Mg alloys versus Mg con-
centration: a) Vickers hardness HV5, b) compressive tests (UCS –
ultimate compressive strength, CYS – compressive yield strength), c)
bending tests (UBS – ultimate bending strength, BYS – bending yield
strength), d) tensile tests (UTS – ultimate tensile strength, TYS –
tensile yield strength, E – elongation)
Slika 2: Mehanske lastnosti zlitin Zn-Mg glede na koncentracijo Mg:
a) trdota po Vickersu HV5, b) tla~ni preizkus (UCS – kon~na tla~na
trdnost, CYS – meja plasti~nosti pri tla~nem preizkusu), c) upogibni
preizkus (UBS – kon~na upogibna trdnost, BYS – meja plasti~nosti pri
upogibnem preizkusu), d) natezni preizkus (UTS – kon~na natezna
trdnost, TYS – meja plasti~nosti pri nateznem preizkusu, E – raztezek)
increases with the increasing Mg concentration and
reaches the maximum of 320 MPa at 0.8 % Mg. At
higher Mg concentrations, the bending strength progres-
sively reduces to 110–120 MPa at 3.5–5.4 % Mg. The
bending yield strength also increases with the increasing
Mg concentration, i.e., with the increasing volume
fraction of the eutectic mixture, and reaches 253 MPa at
1.6 % Mg. The alloys with higher Mg amounts fracture
before the onset of the macroscopic plastic deformation.
The results illustrated in Figure 2c also show that the
plasticity of Zn-Mg alloys during bending is lower com-
pared to that during compressive loading. In the com-
pressive tests, the Zn-3.5Mg alloy can withstand a
significant plastic deformation, whereas in the bending
tests this alloy is macroscopically brittle. The reason for
this difference is in the nature of the local stresses in the
alloys during loading. During macroscopic compressive
loading, the compressive component of the local stresses
predominates. On the other hand, bending induces the
local tensile stresses supporting the formation and
growth of defects and macroscopic fracture cracks.
Tensile mechanical properties are summarized in
Figure 2d. As in the previous case, the ultimate tensile
strength increases up to 170 MPa at 0.8 % Mg. Higher
Mg concentrations lead to a decrease in the ultimate
tensile strength to 73 MPa for the Zn-2.5Mg alloy. The
tensile yield strength reaches the maximum of 124 MPa
at 0.8 % Mg. At this concentration an elongation of the
alloy is 2 %. The alloys with the Mg concentrations
above this limit fractured before the macroscopic plastic
deformation due to an increased fraction of the brittle
Mg2Zn11 phase. These results confirm the above finding
that the plasticity of Zn-Mg alloys decreases with the
increasing tensile component during loading. In the
tensile testing, the most detrimental tensile-stress com-
ponent is the most pronounced supporting an easy nucle-
ation and growth of fracture cracks.
The fracture surfaces of the selected alloys after
tensile testing are shown in Figure 3. One can observe
that zinc shows a brittle and intercrystalline fracture
(Figure 3a) without any plastic deformation. Individual
grains are clearly visible in this figure. Figure 3b shows
the fracture surface of the Zn-0.8Mg alloy with the
highest bending and tensile strengths. In contrast to the
pure zinc, this surface has a significantly refined
morphology, in which both the primary zinc and eutectic
(Figure 1b) can be clearly distinguished. The primary
zinc dendrites are characterized by an almost brittle
fracture that corresponds to the flat facets on the fracture
surface. The eutectic mixture that surrounds these facets
exhibits a refined morphology and there is an indication
of a plastic deformation in this area. The refined fracture
morphology with a certain degree of plastic deformation
is associated with improved bending and tensile
strengths of this alloy because both the fine grains and
the hard network represent barriers for the growing
crack. It appears that the Zn-0.8Mg alloy provides the
optimum volume fractions of both structural compo-
nents. The fracture surface of the Zn-2.5Mg alloy in
Figure 3c also includes the flat facets of the primary
-Zn and the regions of refined morphology corres-
ponding to the -Zn + Mg2Zn11 eutectic mixture (Figure
1d). But a high portion of the brittle Mg2Zn11 eutectic
phase negatively affects both the bending and tensile
strengths because this phase acts as a source of defects
during mechanical loading at which fracture cracks
nucleate.
Based on the results shown in Figure 2, it can be
assumed that the Zn-0.8Mg alloy is the most promising
material for load-bearing implants because it reaches the
maximum bending and tensile strengths of 320 MPa and
168 MPa, respectively. The Zn-0.8Mg alloy also shows a
good plasticity in all three loading modes. In compres-
sion this alloy is able to be deformed without a fracture.
When considering an application of Zn-Mg alloys as, for
example, the fixation screws or plates for fractured
bones, the mechanical properties of the alloys should be
compared to those of bones or other biomaterials. Table
2 provides a summary of the tensile, compressive and
bending mechanical properties of bones, the Zn-0.8Mg
J. KUBÁSEK et al.: STRUCTURAL, MECHANICAL AND CYTOTOXICITY CHARACTERIZATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 623–629 627
Figure 3: Fracture surfaces of Zn-Mg alloys after the tensile tests: a) Zn, b) Zn-0.8Mg, c) Zn-2.5Mg
Slika 3: Povr{ina preloma zlitin Zn-Mg po nateznih preizkusih: a) Zn, b) Zn-0,8Mg, c) Zn-2,5 Mg
alloy, biodegradable polymers (polylactic acid, PLA),
hydroxyapatite, inert Ti-, Co- and Fe-based alloys.1,23–28
One can see that the Zn-0.8Mg alloy is characterized by
significantly higher tensile and bending strengths as
compared to the biodegradable PLA and hydroxyapatite.
Moreover, the strength and elastic modulus of this alloy
are much closer to those of the bone, as compared to the
inert Ti-, Co- or Fe-based biomaterials.
3.3 Cytotoxicity
The Zn and Mg concentrations in the extracts pre-
pared using the Zn-0.8Mg alloy and Mg are 4 μg/mL and
43 μg/mL, respectively. The significantly lower concen-
tration of zinc results from a much better corrosion
resistance of the zinc alloy as compared to magnesium,
as shown in our previous work.22 Figure 4 illustrates the
cytotoxic effects of the extracts on the U-2 OS cells,
expressed as the percent absorbance of the DMEM
control. It is observed that pure magnesium is tolerated
well by U-2 OS cells because these cells are fully viable
in the extract containing 43 μg/mL Mg (almost a 100 %
absorbance). This measurement confirms the presump-
tion stated in the experimental section that magnesium
has a good biocompatibility. What is more important in
this study is that the U-2 OS cells exposed to the extract
from the Zn-0.8Mg alloy also show a good viability of
80 %, i.e., only slightly lower than in the case of the Mg
extract. Due to a low corrosion rate of zinc,22 its concen-
trations in the extracts are very low and, therefore, such
extracts are not toxic for the cells. A similar situation can
be expected in the case of Zn implants in the human
body. The zinc present in the body fluids would probably
not cause any toxic effects, despite the lower tolerable
biological limits of zinc as compared to magnesium.
4 CONCLUSIONS
Biodegradable Zn-Mg alloys containing from 0 % to
more than 8 % Mg were investigated in this work. It was
shown that only the alloys containing relatively low con-
centrations of Mg (approximately 1 %) are suitable for
the load-bearing implants in the as-cast state, because
they have high tensile and bending strengths and an
acceptable elongation. The strength of such alloys is
higher than those of biodegradable polymers and hydro-
xyapatite and comparable to that of the bones. Good
mechanical properties result from a relatively fine struc-
ture composed of primary zinc and interdendritic eutec-
tic. It is assumed that an additional improvement in the
strength can be achieved with hot extrusion. At the Mg
concentrations of above 1 %, Zn-Mg alloys become
relatively brittle mainly during tensile loading.
Zn-Mg alloys can be considered as the alternatives to
Mg-based biodegradable alloys. The main advantage of
zinc alloys over magnesium alloys lies in their signifi-
cantly better corrosion resistance in simulated body
fluids. Therefore, the concentrations of the Zn ions
extracted from alloys are low and they do not cause any
significant toxic effects, as demonstrated with the cyto-
toxicity tests involving human osteosarcoma U-2 OS
cells in this study.
Acknowledgements
The research of the biodegradable metallic materials
was financially supported by the Czech Science
Foundation (project no. P108/12/G043).
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J. KUBÁSEK et al.: STRUCTURAL, MECHANICAL AND CYTOTOXICITY CHARACTERIZATION ...
628 Materiali in tehnologije / Materials and technology 48 (2014) 5, 623–629
Figure 4: Cytotoxic effects of diluted extracts on the U-2 OS cells,
expressed as the percent absorbance of the DMEM control
Slika 4: Citotoksi~en vpliv raztopljenih ekstraktov na celice U-2 OS,
izra`en kot dele` absorbance kontrolnega DMEM
Table 2: Basic mechanical properties of various biomaterials and natural bones (PLA-polylactic acid)1,23–28
Tabela 2: Osnovne mehanske lastnosti razli~nih biomaterialov in naravnih kosti (PLA-polimle~na kislina)1,23–28
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J. KUBÁSEK et al.: STRUCTURAL, MECHANICAL AND CYTOTOXICITY CHARACTERIZATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 623–629 629

A. FEKRACHE et al.: STRUCTURAL CHARACTERIZATION OF A BULK AND NANOSTRUCTURED Al-Fe SYSTEM
STRUCTURAL CHARACTERIZATION OF A BULK AND
NANOSTRUCTURED Al-Fe SYSTEM
KARAKTERIZACIJA STRUKTURE OSNOVE IN
NANOSTRUKTURE SISTEMA Al-Fe
Abdelhak Fekrache, Mohamed Yacine Debili, Saliha Lallouche
LM2S, Physics department, Faculty of Science, Badji Mokhtar-Annaba University, 23200 Annaba, Algeria
mydebili@yahoo.fr
Prejem rokopisa – received: 2013-08-06; sprejem za objavo – accepted for publication: 2013-11-12
The main purpose of the present paper is a study of the properties of stable and metastable structures of several binary Al1–xFex
alloys (0 = x = 0.92) made with high-frequency induction fusion and radiofrequency (13.56 MHz) cathodic sputtering from
composite Al-Fe targets, resulting in homogeneous thin films. The study of the lattice parameters and mechanical behaviour was
followed by X-ray diffraction and Vickers microhardness measurements of bulk and sputtered Al-Fe thin films. The
phenomenon of a significant mechanical strengthening of the aluminium by means of iron is essentially due to a combination of
the solid-solution effects and the grain-size refinement. A further decrease in the thin-film grain size can cause a softening of the
solid and then the Hall-Petch relation slope turns from positive to negative at a critical size called the strongest size, which is
coherent with the thin-film dislocation density.
Keywords: aluminium alloys, sputtering, microhardness, thin films, grain size, Hall-Petch
Glavni namen te predstavitve je {tudij lastnosti stabilnih in metastabilnih struktur v ve~ binarnih zlitinah Al1–xFex (x-vrednosti so
v molskih dele`ih 0 = x = 0,92), izdelanih z visokofrekven~nim zlivanjem in radiofrekven~nim (13,56 MHz) katodnim
napr{evanjem iz kompozitnih tar~ Al-Fe, ki omogo~ajo homogene tanke plasti. Po {tudiju mre`nih parametrov in mehanskih
lastnosti je bila izvr{ena rentgenska difrakcija in dolo~ena mikrotrdota po Vickersu osnove in napr{ene tanke plasti Al-Fe. Pojav
ob~utnega pove~anja mehanske trdnosti aluminija z `elezom je zaradi kombinacije med vplivi trdne raztopine in zmanj{anja
velikosti zrn. Nadaljnje zmanj{anje zrn v tanki plasti lahko povzro~i meh~anje in potem se smer razmerja Hall-Petch obrne od
pozitivnega k negativnemu pri kriti~ni velikosti, za katero je zna~ilna najve~ja koherenca z gostoto dislokacij v tanki plasti.
Klju~ne besede: aluminijeve zlitine, napr{evanje, mikrotrdota, tanke plasti, velikost zrn, Hall-Petch
1 INTRODUCTION
The characterization of solidification microstructures
is essential in many applications. However, the com-
position complexity of most technical alloys makes such
an analysis quite difficult. Microcrystalline and nano-
crystalline materials can currently be produced with
several methods, like the rapid solidification (RS) or
physical vapor deposition (PVD), and the resulting metal
has a polycrystalline structure without any preferential
crystallographic grain orientation. Aluminium and its
alloys with their low densities and easy working have a
significant place in the car industry, aeronautics and food
conditioning. The on-glass-slides, sputter-deposited,
aluminium-based, alloy thin films such as Al-Mg,1
Al-Ti,2,3 Al-Cr4 and Al-Fe5–7 exhibit a notable solid
solution of aluminium in the films and microhardness
values higher than those of the corresponding traditional
alloys.
The inverse Hall–Petch effect (IHPE) has been
observed for nanocrystalline materials by a large number
of researchers.8,9 This effect implies that nanocrystalline
materials get softer as the grain size is reduced below its
critical value. In this paper, we report on a study of the
inverse Hall–Petch effect with respect to a practical
question as to whether ductility is increased in high-
strength metals.
The goal of this paper is to highlight the particular
structural behaviours of different Al-Fe alloys prepared
by RF magnetron sputtering on glass substrates in terms
of structure, lattice parameter, grain size, dislocation
density and deviation from the normal Hall-Petch
relation.
2 EXPERIMENTAL DETAILS
The eight bulk samples used in the present work,
shown in Table 1, were quenched from the liquid state
after high-frequency induction fusion. Powder alumi-
nium and iron (99.999 %) were used in the proportions
defined according to the required compositions. The total
mass of the samples to be elaborated was between 8 g
and 10 g. A cold compaction of the mixed powder
(Al-Fe) was achieved to obtain a dense product (60 %),
intended for a high fusion frequency (HF). A sample
densified in this way was then placed in a cylindrical
alumina crucible (height of 3 cm and diameter of 16
mm), introduced into a quartz tube and placed in the coil
prior to the high-frequency fusion. After the primary
vacuum, the heating of the sample was carried out in
steps, with a ten-minute maintenance stage towards 600
°C until the complete fusion of the alloys at a tempera-
ture of about 1140 K, as determined with a pyrometer.
Materiali in tehnologije / Materials and technology 48 (2014) 5, 631–637 631
UDK 669.715:532.6:669.058 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)631(2014)
Light microscopy (using a Philips microscope) was
used for the polished surface observations. The micro-
structure of the alloys was examined on metallographic
microsections. The mechanical polishing technique
involved 600–4000 SiC grinding paper. The samples
were etched for 15 s with Keller’s reagent (5 mL HF + 9
mL HCl + 22 mL HNO3 + 74 mL H2O). X-ray diffrac-
tion analyses were performed using a Philips X-ray
diffractometer working with a copper anticathode ( =
0.154 nm) and covering 180° in 2. The samples were
subjected to heat treatments in primary vacuum media at
500 °C for a period of 1 h.
The twelve targets used in the elaboration of the
aluminium-iron thin films were made from a bulk
aluminium crown of 70 mm of diameter in which is
inserted a bulk copper or iron disc. Using bulk material
minimizes the oxygen in the films. This target shape
enables the easy control of an additional element
composition in the films (Table 1).
The films were on 75 mm × 25 mm × 1 mm glass
slides that were radiofrequency (13.56 MHz) sputter-
deposited under low pressure of 0.7 Pa and a substrate
temperature that does not exceed 400 K. The sub-
strate–target distance was 80 mm. The sputtering is
carried out with a constant power of 200 W, an auto-
polarization voltage of –400 V, that acquired by the
plasma is –30 V, a regulation intensity of 0.5 A and a
argon flow of 30 cm3/min. After 1 h and 30 min, the
deposition velocity is 2.5 μm/h and films of about 3 μm
to 4 μm thickness were obtained.
The chemical analysis of atomic Fe in Al-Fe was
made by X-ray dispersion spectroscopy. The microstruc-
ture of the films was studied by X-ray diffraction (XRD)
using a Philips X-ray diffractometer working with a
cobalt K anticathode ( = 0.179 nm) and covering 120°
in 2, and transmission electronic microscopy (Philips
CM12) operating under an accelerating voltage of 120
kV. The Vickers indentation under low load allows us to
specify by means of the microhardness the mechanical
strengthening of the aluminium by iron addition. The
measures were realized by means of a Matsuzawa MTX
microdurometre. To reach the intrinsic hardness of the
deposit and free itself from the influence of the substrate,
the Bückle law10 must be taken into account. This law
imposes a depth of penetration h that does not exceed a
tenth of the thickness e of the deposit. So, to have h < 0.1
e, it is necessary to respect the condition D < 0.7 e,
where D is the diagonal of the square impression left by
the Vickers indenter (pyramid of angle in the summit
equal to 136 °). We chose to work with a normal load of
0.1 N (10 g). In addition, the deposit had to have a
thickness of at least 10 μm so that the previous condition
was satisfied.
Thin film specimens were sealed in silica ampoules
in an argon atmosphere, after a previous evacuation to a
pressure of 1.33 × 10–6 mbar, and then heat treated at 500
°C for a period of 1 h.
3 RESULTS AND DISCUSSION
The liquid-quenched Al-60 % Fe alloy is characte-
rized by an ordered B2 (FeAl) CsCl-type structure, as
revealed by the X-ray diffraction pattern (Figure 1a),
and described in Table 2. For 74 % Fe we observe a
structural change leading to a DO3-ordered structure
(Figure 1b), while when the iron content reaches 85 %
and as showed by X-ray diffraction pattern of Figure 1c,
the structure changes completely, giving rise a disor-
dered -Fe solid solution (Figure 2 and Table 2).
Table 2: Phase limits in bulk Al-Fe produced by high-frequency
induction fusion
Tabela 2: Fazne meje v osnovi iz Al-Fe, izdelani z visokofrekven~nim
indukcijskim zlivanjem
A. FEKRACHE et al.: STRUCTURAL CHARACTERIZATION OF A BULK AND NANOSTRUCTURED Al-Fe SYSTEM
632 Materiali in tehnologije / Materials and technology 48 (2014) 5, 631–637
Figure 1: X-ray diffraction patterns of various as-solidified Fe-rich
alloys: a) B2, b) DO3 and c) -Fe
Slika 1: Posnetki rentgenske difrakcije razli~nih strjenih z Fe bogatih
zlitin: a) B2, b) DO3, c) -Fe
Table 1: Chemical compositions of the bulk and the sputtered Al-Fe alloys (amount fractions, x/%)
Tabela 1: Kemijska sestava osnove in napr{ene zlitine Al-Fe (mno`inski dele`i, x/%)
x(Fe)/%
Bulk 5.09 10.78 17.16 30 40 60 74 85
Sputtered 5.6 7.8 18 23 27 29 36 39 47 70 71.9 71.9
For sputtered Al-Fe, the lattice parameter of the -Al
phase decreases from 0.405 nm (pulverized pure alumi-
nium deposit) to 0.403 nm (pulverized deposit contain-
ing x(Fe) = 5 %). This decrease in the parameter is not
surprising since the radius of the iron atom (RFe = 0.124
nm) is lower than that of the aluminium atom (RAl =
0.143 nm).
The lattice parameter of the body-centred-cubic
phase or the B2 phase (in the composition field x(Fe) =
45–55 %) decreases in an appreciably linear way bet-
ween x(Fe) = 38 % (x(Al) = 62 %) (a = 0.295 nm) and
pure iron (a = 0.287 nm) while passing the value a =
0.291 nm for the pulverized deposit containing x(Fe) =
70 % (x(Al) = 30 %) (Figure 3). This decrease is again
explained by the difference in size between the alumi-
nium and iron atoms.
With the fraction of Al increasing, the bulk Al-Fe
lattice parameter increases linearly, which indicates that
the Al simply substitutes for Fe on the Fe sublattice.
There is a change of slope that occurs at 20 % Fe, but
when the percentage of Fe is larger than 20 %, the lattice
parameter decreases, which may indicate that these com-
positions are in a two-phase field, where the BCC-to-
FCC transition may occur between 30 % and 40 % Fe,
see the inset in Figure 3. Similar results were obtained
by Pike et al.11 The results clearly indicate that the larger
Fe atom preferentially occupies the anti-structure sites
on the Al sublattice, and only when these are filled do
the Fe atoms begin to occupy the vacancy sublattice.
Between amount fractions 10 % and 20 % Fe, the bulk
alloy microhardness remains almost constant. Beyond
20 % Fe it will begin increasing until a maximum at
40 % Fe (Figure 4). We observed a Gaussian-shaped
curve for the as-solidified alloys. The effect of iron on
the mechanical properties of aluminium alloys has been
reviewed extensively.12,13 The detrimental effect of iron
on the ductility is due to two main reasons: 1) the size
and number density of iron-containing intermetallics like
Al3Fe and Al2Fe increases with iron content, and the
more intermetallics there are, the lower the ductility; 2)
as the iron-level increases, the porosity increases, and
this defect also has an impact on the ductility (Table 2).
A. FEKRACHE et al.: STRUCTURAL CHARACTERIZATION OF A BULK AND NANOSTRUCTURED Al-Fe SYSTEM
Materiali in tehnologije / Materials and technology 48 (2014) 5, 631–637 633
Figure 3: Lattice-parameter variation with iron composition on the
aluminium-rich side and iron-rich side. Inset shows a change of the
slope for 20 % iron.
Slika 3: Spreminjanje mre`nih parametrov s koli~ino `eleza na z alu-
minijem bogati strani in z `elezom bogati strani. Vlo`eni diagram
prikazuje spremembo naklona pri x(Fe) = 20 %
Figure 2: As-quenched microstructures from bulk Al-Fe
Slika 2: Kaljena mikrostruktura osnove iz Al-Fe
Figure 4: Microhardness evolution with iron content for bulk Al-Fe
Slika 4: Spreminjanje mikrotrdote z vsebnostjo `eleza v osnovi iz
Al-Fe
For the Al-Fe deposits the intrinsic microhardness of
the thin films increases according to the content of iron
from 130 HV (pure aluminium) up to a maximum in the
form of plate of 800 HV, located between 45 % and 70 %
Fe, and then follows a decrease to reach that of iron
towards 400 HV (Figure 5).
We have shown in previous work6,14 that in alumi-
nium-based thin films the microhardness is always
related to the structural and sub-structural features via
the influence of the technological physical conditions of
vapour condensation and film growth.
3.1 Grain size
Two methods were used for the quantitative approach
of the grain size. The first is the application of the
Scherer formula.15 This is based on a measure of the
width of the X-ray diffraction field via a measurement of
the angular width (2). The crystallites average dimen-
sion being given by <D> = 0. 9/(2(cos), where  is
the wavelength of the radiation used,  is the angular
position of the diffraction line and  (2) is the width
with half intensity expressed in radians. This method
assumes the exploitation of diagrams obtained in /2
focusing mode with a low divergence of the incidental
beam. In order to limit the errors, diagrams on alumi-
nium and iron with coarse grains (several micrometres)
allowed a free from the instrumental widths of the lines
(111) Al and (110) bcc (body centred cubic) which was
used. The results that come from this method provide a
good estimate of the grain size when the grain is smaller
than 1 μm.
The second method consists of evaluating the grain
size starting from images obtained using transmission
electron microscopy (Figure 6). The evolution of the
-Al grain size in the presence of iron is similar to that
already observed in the presence of chromium or
titanium.3,4 Whereas the grain of a pulverized pure
aluminium deposit has a size of about 1 μm, this falls to
approximately 500 nm for x(Fe) = 5 %. Beyond this
composition, the microstructure in the two-phase field
(-Al + amorphous) becomes increasingly fine with
grains whose dimensions do not exceed 30 nm to 40 nm
(Figure 7). The refinement of the microstructure, in
cathodic sputtering, at the time of the addition of an
alloy element in aluminium, is constant, because this
element is substituting in the solid solution5,14 or insert-
ing in the aluminium.16
Concerning the body-centred-cubic phase and the
ordered B2 simple cubic phase observed for iron con-
centrations higher than x = 38 %,17 the grain size varies
slightly with iron content in the range of the composition
A. FEKRACHE et al.: STRUCTURAL CHARACTERIZATION OF A BULK AND NANOSTRUCTURED Al-Fe SYSTEM
634 Materiali in tehnologije / Materials and technology 48 (2014) 5, 631–637
Figure 7: X-ray diffraction pattern of wholly amorphous Al-81.5 %
Fe deposit and quasi-amorphous Fe-70.5 % Al deposit
Slika 7: Posnetek rentgenske difrakcije popolnoma amorfnega nanosa
Al-81,5 % Fe in kvazi amorfnega nanosa iz Al-70,5 % Fe
Figure 6: Bright-field transmission electron micrographs and associ-
ated selected-area diffraction ring pattern showing a mixture of nano-
crystalline and amorphous phases from an Al-7.5 % Fe deposit
Slika 6: Posnetek mikrostrukture s presevnim elektronskim mikro-
skopom in izbrano podro~je difrakcije, ki prikazuje me{anico
nanokristalini~nih in amorfnih faz v nanosu iz Al-7,5 % Fe
Figure 5: Comparative microhardness variations with iron content for
bulk and sputtered Al-Fe
Slika 5: Primerjava spremembe mikrotrdote z vsebnostjo `eleza v
osnovi in v napr{enem Al-Fe
studied (38 % to 72 % Fe) and lies between 200 nm and
250 nm (Figure 8 and Table 3).
Table 3: Phase limit in deposits
Tabela 3: Meje faz v nanosih
For Al-Fe films containing between amount fractions
30 % and 40 % iron we observe an inverse evolution of
the Hall-Petch relationship (IHPR) (Figure 9).
However, as the crystal is refined from the micro-
metre regime into the nanometre scale, this mechanism
will break down because the grains are unable to support
dislocation pile-ups. Typically, this is expected to occur
for grain sizes below 10 nm for most metals.18
There is a growing body of experimental evidence for
such unusual deformations in the nanometre regime;
however, the underlying atomistic mechanisms for the
IHPR remain poorly understood. The physical origin of
the IHPR transition and the factors dominating the
strongest size are a long-standing puzzle.19
Two main plausible hypotheses have been advanced
to explain the deviation from the Hall-Petch relation.
First in the HPR regime, crystallographic slips in the
grain interiors govern the plastic behaviour of the poly-
crystallite; while in the IHPR regime, grain boundaries
dominate the plastic behaviour. This hypothesis is sup-
ported by recent computer simulations of deformation in
ultrafine-grained material.20 However, it is not clear from
these simulations that grain-boundary sliding could
become dominant at grain sizes as large as 20 nm; a
recent simulation for Cu suggests a transition at 6–7 nm.9
Second, very small grains cannot support distributions of
dislocations, so the pile-up and dislocation-density
mechanisms for Hall-Petch behaviour cease to apply.
Relevant experimental work has recently been published
by Misra et al.21
3.2 Dislocation density
The finer the grains, the larger the area of the grain
boundaries that impedes the dislocation motion. Further-
more, grain-size reduction usually improves toughness as
well. The dislocation density for Al-Fe thin films has
been determined by using the Williamson and Smallman
method.22 Between amount fractions 7.8 % Fe and 36 %
Fe, the dislocation density of heat-treated thin films is
more sensitive to iron than in the as-deposited specimen
(Figure 10). This phenomenon may be explained by the
relatively small grain size of the as-prepared coatings.
From 36 % Fe the dislocation density drops drastically,
probably due to the structural change of the coatings
from a mixture Al (fcc) or Fe (bcc) with an amor-
phous phase to crystalline (bcc) phase. This behaviour is
coherent with the inverse Hall-Petch effect (IHPE)
A. FEKRACHE et al.: STRUCTURAL CHARACTERIZATION OF A BULK AND NANOSTRUCTURED Al-Fe SYSTEM
Materiali in tehnologije / Materials and technology 48 (2014) 5, 631–637 635
Figure 8: Grain size evolution with iron content
Slika 8: Spreminjanje velikosti zrn od vsebnosti `eleza
Figure 10: Dislocation density versus iron content for as-deposited
and annealed Al-Fe thin films
Slika 10: Gostota dislokacij v odvisnosti od vsebnosti `eleza za nane-
sene in `arjene tanke plasti Al-Fe
Figure 9: Variation of microhardness with the inverse square root of
the grain size
Slika 9: Spreminjanje mikrotrdote z inverzno vrednostjo kvadratnega
korena velikosti zrn
observed in the same iron concentration range. However,
as the crystal is refined from the micrometre scale into
the nanometre scale, this mechanism will break down
because the grains are unable to support dislocation
pile-ups.
The dislocation density after a subsequent heat treat-
ment of the coatings decreases in nearly the same
manner as in the as-deposited state until 36 % Fe.
Beyond this composition the variation become very
slight, because the grain has reached its micrometre size
3.3 Micro deformation
There are several methods (we chose that of William-
son and Hall23) that can determine the average grain size
(D) and the average microstrain rate (	).
For aluminium compositions between amount frac-
tions 30 % and 70 %, the microstrain coming from the
tension stress varies smoothly. Beyond 70 % Al the vari-
ation become more pronounced, until 90 % Al, and cor-
responds to the amorphous domain phase. For alumi-
nium compositions higher than 90 %, the microstrain
falls again in the domain of -Al solid solution (Figure
11).
A transition of the microstrain from tensile to com-
pressive can be seen after the subsequent heat treatment
at 500 °C for a period of 1h, for amorphous films with an
aluminium content between mole fractions 68 % and
90 %.
It is well known that the atomic peening of the grow-
ing film by energetic particles is currently believed to
favor both a dense morphology and grain size refine-
ment. The energetic particles are not only the sputtered
metal atoms, but also the high-energy neutral reflected
gas atoms.24,25 Both the flux and the energy of the
high-energy neutral reflected gas atoms are proportional
to the Mt : Mg ratio, where Mt is the atomic mass of the
target material and Mg is that of the gas. As the transient
metal (TM) is always heavier than Al, increasing the TM
insert size on the target is equivalent to increasing the
Mt : Mg ratio, and this leads to an enhancement of the
in-situ bombardment of the growing film.
4 CONCLUSION
The analysis of the main experimental results issued
from the present study leads to an extended solid
solution in sputtered films versus liquid quenched alloys
and significant mechanical strengthening of the alumi-
nium by means of iron, essentially due to a combination
of solid-solution effects and grain-size refinement. The
lattice-parameter change of slope that occurs at 20 % Fe
in a bulk alloy may indicate that the BCC-to-FCC
transition may occur between amount fractions 30 % and
40 % Fe.
The bulk alloy microhardness is related to the
detrimental effect of iron on the ductility. The Gaussian
increase in the hardness of the alloys as the Fe content
increases is explained by the intermetallic Al3Fe and
Al2Fe phase formation. These phases are found in an
eutectic-like structure over a wide composition range,
while for Al-Fe deposits, the intrinsic microhardness of
the thin films increases in a parabolic way according to
the content of iron.
A transition of the microstrain from tensile to com-
pressive can be seen after the heat treatment at 500 °C
for a period of 1 h, for amorphous films with aluminium
contents between mole fractions 68 % and 90 %.
On an other hand, the dislocation density is observed
to exhibit a decreasing trend in both the as-deposited and
heat-treated specimens. From 36 % Fe the dislocation
density drops sharply, probably due to the structural
change of the coatings from a mixture Al (fcc) or Fe
(bcc) with an amorphous phase to a crystalline (bcc)
phase. This behaviour is in line with the inverse Hall-
Petch effect (IHPE) observed for the same concentration
range of iron.
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A. FEKRACHE et al.: STRUCTURAL CHARACTERIZATION OF A BULK AND NANOSTRUCTURED Al-Fe SYSTEM
Materiali in tehnologije / Materials and technology 48 (2014) 5, 631–637 637

S. ªAHÝN et al.: WEAR BEHAVIOR OF Al/SiC/GRAPHITE AND Al/FeB/GRAPHITE HYBRID COMPOSITES
WEAR BEHAVIOR OF Al/SiC/GRAPHITE AND
Al/FeB/GRAPHITE HYBRID COMPOSITES
VEDENJE HIBRIDNIH KOMPOZITOV Al/SiC/GRAFIT IN
Al/FeB/GRAFIT PRI OBRABI
Salim ªahýn, Nilay Yüksel, Hülya Durmuº, Simge Gençalp Ýrýzalp
Celal Bayar University, Faculty of Engineering, Materials Engineering Department, Manisa, Turkey
nilay.yuksel@cbu.edu.tr
Prejem rokopisa – received: 2013-08-12; sprejem za objavo – accepted for publication: 2013-11-26
Silicon carbide is often the preferred reinforcement in the production of aluminium-powder composites. In this study, alumi-
nium composites were produced with 10 % and 20 % silicon-carbide and ferroboron reinforcements and (0, 0.5, 1 and 1.5) %
graphite additions using powder metallurgy. The effects of the reinforcement type, the amount and the graphite content on the
wear resistance were investigated. When compared with the unreinforced aluminium sample, it was clear that the increasing
reinforcement increased the wear resistance. It was determined that the increasing graphite content negatively affects the wear
resistance. The sample including 20 % ferroboron and 0 % graphite showed the minimum wear rate.
Keywords: aluminum hybrid composite, ferroboron, silicon carbide, wear
Silicijev karbid se pogosto uporablja za oja~anje pri proizvodnji aluminijevih kompozitov iz prahov. V tej {tudiji so bili alumi-
nijevi kompoziti izdelani po metalurgiji prahov, utrjeni z 10 %, 20 % silicijevega karbida, fero-bora in dodatkom (0, 0,5, 1 in
1,5) % grafita. Raziskani so bili vplivi vrste oja~itve, koli~ine oja~itve in koli~ine grafita na odpornost proti obrabi. V primerjavi
z neutrjenimi vzorci aluminija je razvidno, da se odpornost proti obrabi pove~uje z ve~anjem utrjevanja. Ugotovljeno je, da
pove~anje koli~ine grafita negativno vpliva na odpornost proti obrabi. Najmanj{o stopnjo obrabe je pokazal vzorec, ki je
vseboval 20 % fero-bora in bil brez grafita.
Klju~ne besede: aluminijev hibridni kompozit, fero-bor, silicijev karbid, obraba
1 INTRODUCTION
Powder metallurgy is considered as a good technique
for producing metal-matrix composites1 and has a wide
range of applications ranging from automotive to ad-
vanced aerospace components2. P/M components are an
established economic alternative to the components
made with other manufacturing processes as well as the
only means to produce those components that cannot be
made with other methods2. One of the best properties of
the composites fabricated with powder metallurgy is ob-
tained when the reinforcement is homogeneously dis-
persed in the matrix, as proven with both experimental
and theoretical studies3. Another advantage of the
powder-metallurgy technique is the fact that it allows us
to manufacture near-net-shape products at low costs1.
Al-based particulate-reinforced metal-matrix compo-
sites have attracted much interest due to their potential
use and desirable properties4. Aluminum is one of the
best materials for the matrix because of its low density,
high conductivity and high toughness. The other
advantage of using Al for the matrices of MMCs is its
corrosion resistance which is very important when using
composites in different environments3,5. Aluminum P/M
alloys are used in the automobile industry for cylinder
liners, cylinder blocks and drive shafts, replacing more
traditional ferrous alloys6,7. Moreover, Al composites are
used for helicopter parts in aeronautics, such as the parts
of the body, the support for rotor plates and rotor vanes
in compressors7. Their use is a part of the trend toward
the materials that can reduce the weight of a vehicle6.
However, a low wear resistance of pure aluminum is a
serious drawback in using it in many applications1. An
addition of a non-metallic second phase such as oxides,
carbides, nitrides and borides to aluminum alloys can
dramatically improve the mechanical properties and
wear resistance of the materials8. Particle reinforcements
are more favorable than the fiber type as they allow a
better control of the microstructure and mechanical pro-
perties obtained by varying the size and the volume frac-
tion of the reinforcement1.
Aluminium-matrix composites (AMCs) reinforced
with SiCP are recognized as important advanced structu-
ral materials due to their desirable properties, including a
high specific stiffness, a high specific strength, a high
temperature resistance and an improved wear resistance9.
So, many studies were carried out on the preparation and
wear properties of the Al-matrix composites with the
SiC-particle reinforcements10–13. However, there are not
many works on the wear behavior of the Al-matrix com-
posites with the FeB-particle reinforcements. Generally,
the effect of the FeB reinforcement on the wear pro-
perties was investigated for iron-based P/M alloys14–17.
In addition to the contact configuration for different
types of the wear tester, the wear behavior of materials is
related to several factors such as the properties of the
materials sliding against each other and experimental
Materiali in tehnologije / Materials and technology 48 (2014) 5, 639–646 639
UDK 621.762:669.018.9 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)639(2014)
conditions including environmental conditions, load,
speed, etc.8 The wear rate and friction coefficient of an
Al-matrix composite strongly depends on the reinforce-
ment particles, the amounts of the reinforcement and
graphite12,13,18,19. It was found that the wear rate of a
hard-particle composite is substantially lower than that
of the base material7,20. The reinforcement of an alumi-
num matrix with SiC or FeB was generally found to
improve the wear resistance7,14,21.
In this study, an aluminum-matrix composite with the
reinforcement of SiC or FeB and an addition of graphite
(Gr) particulates was explored for tribological properties.
In many works, it was reported that graphite particulates
form a solid lubricant on a tribosurface19,22,23. Generally,
an addition of graphite to aluminum alloys is known to
decrease the strength, fracture energy, ductility and hard-
ness of the materials. A graphite particulate has a brittle
structure; therefore, the tendency of crack initiation and
propagation increases at the graphite-metal interface19,22.
Also, the final properties of the metal-matrix composites
depend on the matrix and ceramic reinforcements, the
bonding of the ceramic reinforcements, the size and dis-
tribution of ceramic reinforcements and the graphite
particulates in an aluminum matrix. According to the
literature studies, little information about the effect of
the FeB reinforcement and graphite particulates on the
wear properties is available. In this work, the effect of
the FeB reinforcement, the SiC reinforcement and
graphite on the wear properties of P/M composites was
investigated.
2 MATERIALS AND METHODS
2.1 Production of the composites
In this study Al/SiC/Gr hybrid composites with the
size of Ø 20 mm × 10 mm were produced using the pow-
der-metallurgy method. The chemical composition of the
aluminium powder was 98.23 % Al–0.0056 % Fe–1.52 %
MgO. The sizes of aluminium, FeB and SiC powders
were below 53 μm.
Al/Gr composites were reinforced with SiC and FeB
particles. The composites were produced with the addi-
tions of w = (0, 0.5, 1, 1.5) % graphite and w = (10, 20) %
FeB or SiC (Table 1). The powder mixtures were
pressed under a 400 MPa load, with the size of Ø 20 mm
× 10 mm. The green samples were sintered at 620 °C for
1 h.
2.2 Microstructural investigation
Microstructural examinations were carried out with
SEM to investigate the porosities and particle clusters.
2.3 Density
The densities of the composite samples were
measured according to Archimedes’ principle and their
porosity ratios were calculated.
2.4 Wear tests
A CSM instruments ball-on-disc wear-test unit was
employed in the present work for a tribological analysis
under dry-sliding conditions with the Al-SiCp + Al-FeBp
composites against a 100 Cr6 stainless-steel ball. The
stainless-steel counter-face material had a diameter of
6 mm. All the experiments were carried out at a load of
3 N at room temperature. Each test was performed with a
sliding speed of 20 cm/s and the track diameter was 8
mm. The speed, temperature and sliding-distance
conditions were all kept constant in each test. This
procedure was executed for each sample along the total
sliding distance of 250 m. The coefficient of friction was
recorded during the wear testing by a transducer on the
load arm of the tribometer. The quantitative value of the
wear was obtained by measuring the cross-sectional area
of the wear track and then the wear rate was calculated
by the TRIBOX 2.10.C program.
3 RESULTS AND DISCUSSION
3.1 Microstructural investigation
It can be seen that SiC particles dispersed uniformly
in the aluminium matrix (Figure 1). But in the
FeB-reinforced samples, some clusters of FeB particles
S. ªAHÝN et al.: WEAR BEHAVIOR OF Al/SiC/GRAPHITE AND Al/FeB/GRAPHITE HYBRID COMPOSITES
640 Materiali in tehnologije / Materials and technology 48 (2014) 5, 639–646
Figure 1: Microstructure of 10 % SiC + 1.5 % Gr reinforced com-
posite
Slika 1: Mikrostruktura kompozita, oja~anega z 10 % SiC + 1,5 %
grafita
Table 1: Powder compositions of composite samples (mass fractions,
w/%)
Tabela 1: Sestava prahov kompozitnih vzorcev (masni dele`i, w/%)
Sample Al SiC Gr Sample Al FeB Gr
1 90 10 0 9 90 10 0
2 89.5 10 0.5 10 89.5 10 0.5
3 89 10 1 11 89 10 1
4 88.5 10 1.5 12 88.5 10 1.5
5 80 20 0 13 80 20 0
6 79.5 20 0.5 14 79.5 20 0.5
7 79 20 1 15 79 20 1
8 78.5 20 1.5 16 78.5 20 1.5
were noticed (Figure 2a). The sample including 20 %
FeB + 1.5 % Gr, exhibiting the maximum porosity, had
pores around the reinforcements and also between the
grain boundaries due to insufficient sintering (Figure
2b). In contrast, there were no pores around the rein-
forcement particles of the sample including 10 % SiC +
1.5 Gr which exhibited the minimum porosity (Figure 1);
moreover, no insufficient sintering was confirmed. The
maximum porosity was revealed in the sample with the
maximum reinforcement. It is thought that the pores
originating from insufficient sintering can result from the
cold-pressing pressure. Due to the negative influence of
an increase in the reinforcement on the compressibility,
the sintering process for the sample reinforced with 20 %
FeB + 1.5 % Gr was not adequately maintained. More-
over, the pores arising from the clusters must be consi-
dered.
3.2 Density
The compressibility of composite powders is notice-
ably lower than that of the unreinforced matrix, often
producing a low green density and an insufficient
strength to support secondary processing like sintering,
machining or extrusion4. Reinforcement particles have a
tendency to associate themselves with porosity and give
rise to particle-porosity clusters24. The presence of non-
metallics such as unreduced oxides reduces the com-
pressibility because of their hardness and low specific
gravity2. Rahimian et al.25 reported that as the amount of
alumina increases, the relative density declines. The
reason for this, in comparison with pure aluminum, is the
decline in the pressing capacity of the samples with the
increase in the amount of alumina. This is due to a
higher hardness of alumina. Therefore, these composites
have a lower compressibility resulting in a lower relative
density25. In Figure 3 it can be seen that the increasing
amount of reinforcement has increased the porosity.
SEM images of the microstructures show the increase in
the porosity with the increasing reinforcement (Figures
1 and 2). Also, Tekmen et al.24 found that the increasing
reinforcement volume fraction increases the porosity
content. In the present study the porosity ratio of the
unreinforced aluminium composite was calculated as
3.77 %. So, the minimum porosity ratio compared to the
reinforced samples was shown. It was observed that the
composites with an addition of FeB have a higher poro-
sity compared to the SiC samples. During the micro-
structural investigation, agglomerations of FeB particles
were noticed particularly in larger reinforcement
amounts (Figure 2a). The SiC-reinforced composites
showed a more homogeneous distribution in the alumi-
nium matrix. So, it was concluded that a particle agglo-
meration causes an increase in the porosity.
In general, the porosity due to the increasing graphite
amount exhibited a downward trend except for the com-
posites reinforced with 20 % FeB (Figure 3). Due to the
FeB clusters (Figure 2a), the graphite could not properly
show its lubricant effect during cold pressing.
S. ªAHÝN et al.: WEAR BEHAVIOR OF Al/SiC/GRAPHITE AND Al/FeB/GRAPHITE HYBRID COMPOSITES
Materiali in tehnologije / Materials and technology 48 (2014) 5, 639–646 641
Figure 2: Microstructure of 20 % FeB + 1.5 % Gr reinforced
composite: a) FeB cluster in the structure, b) porosities due to
insufficient sintering
Slika 2: Mikrostruktura kompozita, oja~anega z 20 % FeB + 1,5 %
grafita: a) skupek FeB v strukturi, b) poroznosti zaradi nezadostnega
sintranja
Figure 3: Porosity values of the composites (w/%)
Slika 3: Dele` poroznosti kompozitov (w/%)
3.3 Wear tests
From Figure 4, it can be concluded that the maxi-
mum wear rate was obtained for the sample consisting of
10 % FeB + 1.5 % Gr and the minimum wear rate was
obtained for the sample without graphite and reinforced
with 20 % FeB. The wear-test parameters generated a
very large wear scar on the unreinforced aluminium, so
that the profilometer could not scan the whole section of
the wear scar. For this reason, the wear rate of the un-
reinforced aluminium composite could not be calculated;
however, it can be said that the unreinforced aluminium
shows a higher wear rate than the reinforced ones.
The wear resistance increased as the reinforcement
proportion increased from 10 % to 20 % since the hard
SiC and FeB particles in the matrix resisted the coun-
terpart. The increased reinforcement reduced the contact
area between the counterpart and the relatively soft
matrix, thus the abrasion (wear) was reduced. Figure 5
explains this phenomenon. Here, the hard SiC particles
outcropped and formed an impediment for the contact
between the composite surface and the counterpart (Fig-
ure 5a). The outcropped particles abraded the counter-
part. Figure 6 also shows the abraded areas of the
counterparts of the 10 % FeB + 1.5 % Gr and 20 % SiC
+ 0 % Gr composites. The counterpart of the 20 % SiC +
0 % Gr composite exhibited a larger abrasion area in
comparison with the counterpart of 10 % FeB + 1.5 %
Gr due to the increased reinforcement amount. Ravin-
dran et al.26 reported that the dispersion of silicon
carbide, the hard phase in the soft aluminium matrix,
tends to reduce the wear loss of hybrid composites.
In the literature some researchers maintain that the
graphite amount in aluminium composites enhances the
tribological properties and wear resistance13,27,28.
However, Vencl et al.29 reported that an addition of gra-
phite particles (w = 1 %) to a composite with SiC parti-
S. ªAHÝN et al.: WEAR BEHAVIOR OF Al/SiC/GRAPHITE AND Al/FeB/GRAPHITE HYBRID COMPOSITES
642 Materiali in tehnologije / Materials and technology 48 (2014) 5, 639–646
Figure 6: Counter parts of the composites: a) 10 % FeB + 1.5 % Gr,
b) 20 % SiC + 0 % Gr
Slika 6: Sti~na povr{ina kompozitov: a) 10 % FeB + 1,5 % grafita, b)
20 % SiC + 0 % grafita
Figure 5: Wear surface of 20 % SiC + 0 % Gr reinforced composite:
a) wear tracks, b) EDX analysis
Slika 5: Obrabljena povr{ina kompozita, oja~anega z 20 % SiC + 0 %
grafita: a) sledi obrabe, b) EDX-analiza
Figure 4: Wear rate
Slika 4: Hitrost obrabe
cles further reduced the wear rate and the coefficient of
friction, but this influence of such a relatively small
amount of graphite was not clear enough and should be
considered only as a trend of behavior. According to the
results of the present study, the increased Gr amount
increased the wear rate. Actually, this was an expected
result because the main reason for the graphite addition
to aluminium composites reinforced with hard particles
was to simplify the machining. The presence of the hard,
brittle and abrasive SiC reinforcement makes the mate-
rial difficult to form or machine using traditional manu-
facturing processes. In order to improve the machina-
bility of the SiCp/Al composites, graphite was added to
the composites30.
As an alternative approach, an explanation for the
decreasing wear rate with the increasing reinforcement
might be the pores filled up with the wear debris. A high
reinforcement amount causes high porosity, so the wear
debris could easily fill these numerous pores instead of
being pushed out of the tribological system. Actually, a
negative effect of the increasing pores on the wear resi-
stance was expected. In this case, the pores presumably
filled with the wear debris gave rise to a decrease in the
wear rate calculated as the volumetric wear loss. As can
be seen from Figure 7, a pore on the worn surface of the
20 % FeB + 0 % Gr hybrid composite was filled with the
wear debris. According to the results of the EDX ana-
lysis it can be concluded that wear debris is composed of
the aluminium matrix, the stainless-steel counterpart and
FeB-reinforcement particles (Table 2).
The pores could have been formed due to the insuffi-
cient sintering (Figure 2) and they could have also been
created by the mechanical force during the wear test.
Figure 8 shows the pores caused by the wear test, as the
reinforcement particles were pulled out. These kinds of
pores could be large and easily filled with the wear
debris. The reason for the particle pullout was found
with the EDX analysis carried out along a line. The EDX
graph of graphite shows a significant increase of the
pores arrowed as "reinforcement particle pullouts" in
Figure 8. By means of the EDX analysis (in Figure 8,
the EDX graph on the right-hand side shows graphite) it
was realized that the wall of pores was smeared by
graphite, so the wettability between the matrix and the
reinforcement was reduced. In this case, the particle
pullout occurred easily and due to the low bonding of the
aluminium matrix and the reinforcement. Thus, a detri-
mental effect of graphite on the wear resistance was
confirmed.
The samples that exhibit the maximum wear rate in
both groups, the SiC- and FeB-reinforced composites,
exhibit both abrasive and adhesive wear tracks. For both
reinforcement groups, if the wear surfaces of the most
and the least wear-resistant samples were compared,
adhesion wear was observed to be more intense on the
least wear-resistant samples.
The wear surface of the sample reinforced with 10 %
SiC + 1.5 % Gr exhibited both large craters arising from
adhesion and abrasion scratches (Figure 9a). Similarly,
the wear surface of the sample including 10 % FeB + 1.5
% Gr shows adhesion craters and also flaky wear debris
subjected to delamination (Figure 10). According to the
EDX analysis (Table 3) the aluminium amount of the
delaminated area was larger than the other analyzed area;
consequently, it is maintained that the delaminated part
had a smaller reinforcement amount and was subjected
to a large plastic deformation. The plastically deformed
part became strain hardened and brittle. Ravindran et al.26
S. ªAHÝN et al.: WEAR BEHAVIOR OF Al/SiC/GRAPHITE AND Al/FeB/GRAPHITE HYBRID COMPOSITES
Materiali in tehnologije / Materials and technology 48 (2014) 5, 639–646 643
Figure 7: Pore and wear debris on the worn surface of 20 % FeB + 0
% Gr hybrid composite
Slika 7: Praznine in delci obrabe na obrabljeni povr{ini hibridnega
kompozita z 20 % FeB + 0 % grafita
Table 2: EDX analysis of the pore on the worn surface of 20 % FeB +
0 % Gr hybrid composite (w/%)
Tabela 2: EDX-analiza povr{ine pore na obrabljeni povr{ini
hibridnega kompozita z 20 % FeB + 0 % grafita (w/%)
Element Al Fe Cr Si Mn
1 81.799 13.687 3.542 0.203 0.769
2 76.015 18.834 4.319 0.172 0.661
Figure 8: EDX analysis of 10 % SiC + 1.5 % Gr hybrid composite
along the line; left: Si, right: C (graphite)
Slika 8: Linijska EDX-analiza hibridnega kompozita z 10 % SiC + 1,5
% grafita; levo: Si, desno: C (grafit)
also observed severe plastic deformation of the Al 2024
matrix and a brittle fracture on the wear surface.
Table 3: EDX analysis of 10 % FeB + 1.5 % Gr reinforced composite
(w/%)
Tabela 3: EDX-analiza kompozitov, oja~anih z 10 % FeB + 1,5 %
grafita (w/%)
Element Al Fe Cr C Si Mn
1 93.197 4.631 1.259 0.352 0.188 0.374
2 88.810 6.680 1.628 1.184 1.069 0.630
If the wear surfaces in Figures 9 and 11 are com-
pared, it can be observed that the adhesive-wear tracks
reduced and the abrasive scratches became shallower
with the increase in the reinforcement proportion from
10 % to 20 %. Similar observations were made for the
wear surfaces of the FeB-reinforced samples (Figures 10
and 12).
On the wear surfaces of the most wear-resistant
samples including a 20 % reinforcement of (FeB or SiC)
+ 0 % Gr, it was observed that the reinforcement parti-
cles were pulled out but not taken away from the tribo-
logical system (Figure 9). It is maintained that the
detached hard reinforcement particles might have
become embedded a little into the soft matrix due to the
applied load during the wear test and decreased the wear
rate by reducing the real contact area between the surface
and the counterpart. As can be seen from the EDX analy-
sis (Table 4), the particle marked as "1" is the FeB
reinforcement covered with some of the Al matrix (Fig-
ure 12). Similarly, the pulled-out SiC particle marked as
"1" and the presumably oxidized aluminium particle
marked as "2" were detected with EDX as presented in
Table 5 and Figure 9.
The shape of the reinforcement particles caused a
difference in the abrasive-wear tracks, particularly for
the 10 % reinforced samples. The sharp and angularly
shaped SiC particles caused narrow but deep grooves
(Figure 9), while the relatively round-shaped FeB
particles formed shallower and wider grooves (Figure
11).
S. ªAHÝN et al.: WEAR BEHAVIOR OF Al/SiC/GRAPHITE AND Al/FeB/GRAPHITE HYBRID COMPOSITES
644 Materiali in tehnologije / Materials and technology 48 (2014) 5, 639–646
Figure 11: Wide and shallow wear tracks of 10 % FeB + 1.5 % Gr
reinforced composite
Slika 11: [iroke in plitve sledi obrabe na kompozitu, oja~anem z 10 %
FeB + 1,5 % grafita
Figure 9: Wear surface of 10 % SiC + 1.5 % Gr reinforced composite:
a) abrasion and adhesion marks, b) deep abrasion grooves
Slika 9: Obrabljena povr{ina kompozita, oja~anega z 10 % SiC + 1,5
% grafita: a) obraba in lepljenje, b) globoki abrazijski utori
Figure 10: Wear surface of 10 % FeB + 1.5 % Gr reinforced com-
posite
Slika 10: Obrabljena povr{ina kompozita, oja~anega z 10 % FeB + 1,5 %
grafita
As the two surfaces are brought in contact, the
nanoscale asperities are the first to come into contact,
instantly plastically deforming and merging to form
micro- or macro-contacts, which, upon further applica-
tion of the load, may deform due to elastic or elastic-
plastic deformation31. A further plastic deformation leads
to a brittle fracture. In this study, the sample including
only 20 % SiC exhibited brittle flakes during the wear
test. Here, the plastically deformed aluminium matrix
underwent strain hardening. The brittle layer consisted of
the strain-hardened aluminium and the hard SiC particles
were smeared on its surface. In Figure 13, the cracks
and debris separated from the brittle layer can be seen. It
is assumed that the existence of SiC particles contributes
to the strain-hardened aluminium and so the hardness of
the strain-hardened matrix increases further.
Analyzing the EDX results of the 10 % FeB + 1.5 %
Gr and 20 % FeB + 0 % Gr reinforced samples (Tables 3
and 4), some elements such as Cr and Mn belonging to
the stainless-steel counterpart were observed. In contrast,
the EDX result for the 20 % SiC + 0 % Gr reinforced
composite did not show any element of the stainless-
steel counterpart (Table 5).
4 CONCLUSIONS
• The minimum wear rate was obtained for the sample
without graphite and reinforced with 20 % FeB.
• The 20 % FeB-reinforced samples showed the maxi-
mum porosity and insufficient sintering due to the in-
creased reinforcement. In a further study, pressure
can be raised during cold pressing.
• The maximum wear rate was obtained for the sample
consisting of 10 % FeB + 1.5 % Gr.
• The increased graphite amount increased the wear
rate of all the Al-powder composites.
• The increased reinforcement increased the porosity.
• The wear rate decreased with the increasing rein-
forcement.
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Materiali in tehnologije / Materials and technology 48 (2014) 5, 639–646 645
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Element Al Fe Cr C Si Mn
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2 79.865 17.242 1.994 - 0.459 0.441
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Element Al Fe C Si O
1 0.860 0.220 26.665 71.914 0.341
2 91.952 0.202 3.615 1.020 3.210
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646 Materiali in tehnologije / Materials and technology 48 (2014) 5, 639–646
F. VODE et al.: MATHEMATICAL MODEL FOR AN Al-COIL TEMPERATURE CALCULATION ...
MATHEMATICAL MODEL FOR AN Al-COIL
TEMPERATURE CALCULATION DURING HEAT
TREATMENT
MATEMATI^NI MODEL ZA IZRA^UN TEMPERATURE V
Al-KOLOBARJU MED TOPLOTNO OBDELAVO
Franci Vode1, Franc Tehovnik1, Jaka Burja1, Bo{tjan Arh1, Bojan Podgornik1,
Darja Steiner Petrovi~1, Matja` Malen{ek2, Leonida Ko~evar2, Marjana La`eta2
1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2IMPOL d.d., Partizanska cesta 38, 2310 Slovenska Bistrica, Slovenia
franci.vode@imt.si
Prejem rokopisa – received: 2013-09-09; sprejem za objavo – accepted for publication: 2014-05-28
This paper presents the implementation of a mathematical model for an Al-coil’s temperature evolution during heat treatment in
a forced-circulation furnace. The coil’s spatial-temperature evolution is calculated using the finite-difference method in the axial
and radial directions, i.e., 2 dimensional. The model is verified by measuring the coil’s temperature during reheating using 10
pre-installed thermocouples arranged in two lines of 5 thermocouples, each at a different coil radius. The sensitivities of the
furnace-air temperature and the initial-coil temperature on the time to switch are determined, i.e., –52.5 s/°C and –55.5 s/°C,
respectively. The mathematical model proved to be a multipurpose tool for different simulation-based reheating studies, while
its industrial real-time application offers an even wider range of uses and capabilities: automatic coil-temperature control and
observing reheating conditions on an industrial level.
Keywords: Al-coil, temperature, finite difference, mathematical model
^lanek obravnava uporabo matemati~nega modela razvoja temperature v Al-kolobarju med toplotno obdelavo v pe~i s prisiljeno
cirkulacijo. Razvoj prostorske temperature v kolobarju je izra~unan z uporabo metode kon~nih diferenc v osni in radialni smeri.
Model je bil preizku{en z merjenjem temperature v kolobarju med ogrevanjem s predhodno vgrajenimi 10 termoelementi,
razporejenimi v dveh vrstah po 5 na dveh razli~nih premerih kolobarja. Ugotovljena je bila ob~utljivost temperature atmosfere v
pe~i in za~etne temperature v kolobarju v odvisnosti od pe~i: –52,5 s/°C oziroma –55,5 s/°C. Matemati~ni model se je izkazal
kot ve~namensko orodje za razli~ne simulacije ogrevanja, medtem ko industrijska uporaba v realnem ~asu ponuja {e {ir{e
podro~je uporabe in zmogljivosti: avtomatsko kontrolo temperature kolobarja, opazovanje razmer pri ogrevanju na
industrijskem nivoju.
Klju~ne besede: Al-kolobar, temperatura, kon~ne diference, matemati~ni model
1 INTRODUCTION
The description of physical processes using appro-
priate mathematical models is nowadays achieving
practical usefulness in different fields.1,2 One of these
fields are certainly models for describing heat transfer,
especially in metallic materials, where the metal’s tem-
perature drives, determines or triggers most mate-
rial-transformation processes. Coil temperature is one of
the most important reheating parameters for Al-coils.
The precision of the reheating process is therefore
limited by the precision of the coil’s temperature data.
The coil’s temperature data can be obtained by measur-
ing or from a calculation using mathematical models.
Nowadays, furnace control systems optionally provide a
control method known as the šair-to-work ratio control
system’.3 This method uses two pairs of thermocouples
to measure the šwork–coil’ and šair’ temperatures and
calculates the furnace-temperature set-point as a scaled
difference of both. The method is efficient, but requires a
proper measuring coil temperature. For this method, a
hole is drilled into the coil and a thermocouple is
mounted into it. If the location of the hole is correct and
representative for the whole-coil temperature, the
method is accurate. But the drilled hole represents a
potential defect in the final strip. In any case, measuring
the temperature of every coil in industry is not practical.
A much more practical and non-destructive technique is
to employ a mathematical model (MM) for the calcula-
tion and prediction of coil temperatures. For accurate
MM coil-temperature predictions, (1) the measured air
temperature in the furnace (which all furnaces already
measure and use), (2) the coil-alloy and (3) the coil-
dimension data are required. The rest of the required
information is integrated into the MM during the deve-
lopment and calibration procedure. Thus the MM can be
used instead of the measuring coil’s temperature, but it
should be capable of running in real-time. Another
advantage of the MM for the coil’s temperature is that
the model can be used for optimizations and various
other services, which are impossible without a MM. The
fastest reheating of Al coils in reheating furnaces is
achieved by an overshoot of the furnace-air temperature
above the coil temperature.3 The time from beginning of
the reheating process until the time of decreasing the
furnace temperature from Tot is denoted as the time to
switch, i.e., ts (Figure 1).
Materiali in tehnologije / Materials and technology 48 (2014) 5, 647–651 647
UDK 519.61/.64 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)647(2014)
The aim of the work was to quantify the influence of
the furnace over-temperature and the initial temperature
of the coil on the time to switch (ts) using a mathematical
model.
2 MATHEMATICAL MODEL OF THE COIL
TEMPERATURE
2.1 2D model of the Al-coil temperature using the
finite-difference method
Heat transfer in solid state is mathematically ex-
pressed using the diffusion equation:
∂
∂
T
t
a T= ∇ 2 ; a
c
=


p
(1)
where T is the temperature, t is the time,  is the thermal
conductivity, cp is the specific heat and 
 is the density.
The temperature field is calculated in two dimensions,
i.e., the radial r and the axial x directions.
Equation 1 for the selected r and x coordinate direc-
tions in a cylindrical coordinate system can be written as:


  
c T
T
t
x
T
T
x r
T
T
r
p ( )
( ) ( ) (
∂
∂
∂
∂
∂
∂
∂
∂
∂
∂
=
= ⎛⎝
⎜ ⎞
⎠
⎟ + ⎛⎝
⎜ ⎞
⎠
⎟ + T
r
T
r
)
1 ∂
∂
(2)
where the heating conditions in the angle  direction are
assumed to be symmetric and thus the second derivative
∂ ∂2 2 0T/  = is zero. In equation 2,  is a function of
temperature, while the temperature field as well as the
convective boundary conditions are not constant. In fact,
the boundary conditions for equation (2) are a function
of the measured values of the furnace-atmosphere-gas
temperature and the equation is thus analytically
unsolvable. A handy way for numerically solving eq. 2,
bearing in mind that the boundary conditions (furnace
temperatures) for real-time operation will be available
(measured) one per calculation period, is the explicit
method of finite difference.
For the used finite-difference method the calculation
space is discretized in both time and space:
{ }Δ Δ Δt x r
x r rz n
(sec), ( ), ( ) , ,m m =
−
⎧
⎨
⎩
⎫
⎬
⎭
1
21 21
,
so that the stability condition is satisfied1 within the
proposed coil dimensions and sample time. For each
point on the grid, eq. 2 is rewritten as
( ) ( )
 

c
x r
t
T T
r
x
T T
r
i j
t t
i j
t x
i j
t
i j
t
x
p
Δ Δ
Δ
Δ
Δ
Δ
Δ
Δ
, , , ,
+
−− = − +
+
1
( ) ( )
x
T T
x
r
T T
x
r
T
i j
t
i j
t r
i j
t
i j
t
r
i j
t
+ −
+
− + − +
+ −
1 1
1
, , , ,
,


Δ
Δ
Δ
Δ ( ) ( )T
x
r r
T Ti j
t r
i j
t
i j
t
, , ,+ −+
 Δ
Δ
1
1
(3)
According to the literature data,4 the thermal conduc-
tivity of the Al-coils differs in the r and x directions due
to the air-gap between the strip wraps and the thin
oxide-layer of the strip surface. Thus, the thermal con-
ductivities r and x are considered different, where x
equals the alloy data, while r is lower then x due to the
air gap and the oil films between the coil wraps, which
decreases the thermal conductivity in the radial direction.
From eq. 3, the temperature at time t + 1 for each ele-
ment on the grid (i, j) is expressed as:
T
t
c
T T T
xi j
t t x i j
t
i j
t
i j
t
,
, , ,+ − +=
+ −⎛
⎝
⎜
⎞
⎠
⎟ +
+
Δ Δ
Δ




p
1 1
2
2
r i j
t
i j
t
i j
t
r i
t
c
T T T
r
t
c r
T
Δ
Δ
Δ
p
p





, , ,− ++ −⎛
⎝
⎜
⎞
⎠
⎟ +
+
1 1
2
2
1 , ,
,
j
t
i j
t
i j
t
T
r
T
+ −⎛
⎝
⎜
⎞
⎠
⎟ +
1
Δ
(4)
F. VODE et al.: MATHEMATICAL MODEL FOR AN Al-COIL TEMPERATURE CALCULATION ...
648 Materiali in tehnologije / Materials and technology 48 (2014) 5, 647–651
Figure 2: Coil dimensions, discretization grid, boundary lines and
indexation of x–r cross-section
Slika 2: Dimenzije kolobarja, diskretizacijska mre`a, robne ploskve in
indeksiranje x–r prereza
Figure 1: Furnace over-temperature; time to switch ts
Slika 1: Vi{ja temperatura zraka v pe~i; ~as preklopa ts
Eq. 4 is recursive and is used for inner points of the
calculation grid, while for the boundary points the
equation is modified to consider the convective boundary
conditions.3
2.2 Model calibration
The thermo-physical data of around 40 alloys were
obtained using JMatPro software ver.4.0: 
, cp and  in
the desired temperature range. The MM was calibrated to
the measured temperature values in 10 points, i.e.,
5 points close to the outer radius rz and the remaining
5 on the radius, which is in the middle of the coil’s
length. The air temperature in the furnace (boundary
condition) was measured above the measured coil. The
temperature of the heat treatment for Al-alloys rarely
exceeds 550 °C, and thus convective heat transfer domi-
nates the heat-transfer mode. The conductive heat-trans-
fer mode is not present due to the stage design, while the
radiative heat-transfer thermal mode is estimated to be
sufficiently low compared to the conductive mode. The
maximum ratio of the radiative-to-convective heat flux
  ( ) / ( )q /q A F T T Ah T Tr c c= − −	  air coil air coil
4 4 is estimated
to be  0.031. For a furnace air temperature that is 5 K
above the coil temperature of 550 K (close to the end of
the reheating) this ratio is estimated to be  0.025. For
the estimation, the following assumptions are consi-
dered: 	 = 0.09, F = 1, hc = 150, Tair = 823 K, Tcoil = 293
K. The radiation heat transfer is thus neglected. Note that
a forced-air-circulation furnace is studied. The cali-
bration is made by changing the convective heat-transfer
coefficient hc on the boundary surfaces of the calculated
cross-section area – see the lines of A, B, C, D, F and G
in Figure 2, until a calculated and measured temperature
match is achieved. When the calculated and measured
temperature profiles match at this point, it can be con-
cluded that the model is capable of describing the heat
transfer for the considered conditions (Tair = f(t), hc =
f(T)). To check the possible variability of the process,
additional coil-temperature measurements were per-
formed at the same position in the furnace and with the
same coil, but with a modified Tair = f(t) profile. Unfortu-
nately, the reheating conditions for the measured coil
were slightly modified (closed steel-coil on both sides)
due to studies for improving the homogeneity conditions.
The boundary conditions on the inner coil surface –D
were therefore modified: hc drops from 22 to 5. The mo-
del’s validation is therefore performed with a modified hc
on boundary D and this is compared to the temperature
measurement obtained with the closed steel-coil. A com-
parison is shown in Figure 3 and it can be seen that the
maximum absolute difference in the coil temperature is
around ± 10 °C. However, for a strictly correct model
validation, the process condition should remain intact, so
that hc would remain intact during the validation as well.
3 USE OF THE MATHEMATICAL MODEL
3.1 Off-line calculations of the re-heating procedures
The MM was firstly applied for the prediction of the
Al coil’s temperature evolution for modified reheating
conditions, e.g., modification of the desired end-tempe-
rature of the Al coil and the task was to modify the
furnace time and temperature in such a way as to obtain
a specific temperature homogeneity of the coil at the end
of the treatment. To perform such a task, the boundary
conditions – furnace air temperatures – need to be pre-
dicted (Figure 4).
This is achieved by employing another model, which
predicted the furnace-air-temperature, closed-loop res-
ponse, where the inputs are the furnace-temperature
set-points. A first-order system5 with a different time
constant for increasing/decreasing the furnace tempera-
tures is used to mathematically express the relation. The
furnace temperature model is verified on 10 furnace-tem-
perature profiles for various charging data (coil total
mass from 22–78 t, 9 alloy grades, etc.). Due to the
uncertainty of the model-obtained furnace air tempera-
ture (boundary conditions), the calculated coil tempera-
F. VODE et al.: MATHEMATICAL MODEL FOR AN Al-COIL TEMPERATURE CALCULATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 647–651 649
Figure 4: Furnace air temperature predictions for off-line calculations
substitute measured temperatures, available in real-time calculations
Slika 4: Izra~unane temperature zraka v pe~i v "off-line"-na~inu
nadome{~ajo merjene temperature zraka, ki so v "real-time" na~inu na
voljo
Figure 3: Comparison of measured and calculated coil temperatures.
Note that the validation temperature measurements were obtained for
slightly modified conditions on surface D.
Slika 3: Primerjava merjenih in izra~unanih temperatur kolobarjev.
Validacijske meritve temperature so bile merjene pri spremenjenih
robnih pogojih na povr{ini D.
ture for these boundary conditions decreases the accu-
racy of the coil temperatures compared to the measured
boundary conditions. Note that the prediction-accuracy
of the furnace-air temperature is ± 10 °C. The crucial
benefit of the model-obtained furnace-temperatures
(boundary condition) combined with the coil-tempera-
ture model is the capability for off-line simulations
involving various conditions and parameter sets (coil
dimensions, alloy data, coil initial temperature, furnace
temperature profile during reheating, etc.), but at the cost
of a slightly decreased coil-temperature accuracy. For the
off-line calculations a Graphical User Interface (GUI) is
provided (Figure 5).
3.2 Quantification of the furnace-air temperature and
the initial coil temperature for the coil-temperature
evolution
A frequent task in the Al-coil reheating process is the
modification of the reheating parameters (furnace time
and temperature) according to a specific variation of the
parameters, e.g., variation of the initial coil temperature
of the Al coils or a higher furnace over-temperature Tot.
The coil with an increased initial temperature reaches the
desired temperature faster. A straightforward solution to
such a problem is to employ the developed mathematical
model, including a furnace-temperature model (Figure
4) and change the reheating parameters until the
conditions are met (trial-and-error method). Another way
is to quantify in advance the response of a certain para-
meter change on the reheating parameters. For industrial
production, the quantification of such reheating shorten-
ing is beneficial. For each furnace over-temperature Tot =
(400, 420, 440, 460, 480) °C in set and independently for
each initial coil temperature Tcoil,0 = (0, 20, 40, 60, 80,
100, 120, 140, 160) °C, a simulation is performed and
the resulting time to switch ts is determined from the
model results (Figure 6).
The ts is actually determined by the model during the
simulation, since the model also calculates the Tair,t curve
(Figure 3). The criterion to start decreasing the air
temperature in the simulation model is when the Al-coil
temperature in the node Tcoil,t (10, 10) is Td = 30 °C under
the desired final temperature of the coil, in this case
(280–30) °C = 250 °C. Note that, different Td values lead
to different reheating profiles. Therefore, Td is adjustable
through the GUI. The node (10, 10) is in the middle of
the coil (Figure 2). The other simulation parameters are
constant: Tcoil,0 = 20 °C, Tair,0 = 20 °C, coil dimensions rz
= 900 mm, rN = 280 mm, x = 1250 mm, alloy EN 8079
AlFe1Si. The simulation is repeated for each Tot. The
obtained values are presented in Figure 6. The relation
between Tot and ts is almost linear with a coefficient of
–52.5 s/°C around the working point Tot = 440 °C. The
result shows that the increment of the furnace over-
temperature for a single °C means ts is shorter by 52.5 s.
And conversely, the reduction of the furnace temperature
by a single °C means a longer ts for 52.5 s.
In the same way we estimated the influence of the
initial coil temperature Tcoil,0 on ts. The relation between
the initial coil temperature Tcoil,0 and ts is presented in
Figure 7. The time to switch ts is fairly linear, with a
F. VODE et al.: MATHEMATICAL MODEL FOR AN Al-COIL TEMPERATURE CALCULATION ...
650 Materiali in tehnologije / Materials and technology 48 (2014) 5, 647–651
Figure 7: Influence of initial coil temperature Tcoil,0 = (0 : 20 : 160)
°C on the time to switch ts
Slika 7: Vpliv za~etne temperature kolobarja Tcoil,0 = (0 : 20 : 160) °C
na ~as preklopa ts
Figure 6: Influence of furnace over-temperature on the time to switch
ts
Slika 6: Vpliv vi{je temperature v pe~i na ~as preklopa ts
Figure 5: GUI for offline calculations of reheating procedures
Slika 5: Grafi~ni uporabni{ki vmesnik za izra~une ogrevanja kolobar-
jev v pe~i
coefficient of –55.5 s/°C of the initial coil temperature.
The simulation results show that an increment of +1 °C
for the initial coil temperature means a shorter time 55.5
s to switch and, conversely, a reduction of –1 °C of the
initial coil temperature means a longer time 55.5 s to
switch.
3.3 Potential use of the model modified for a real-time
calculation of the coil temperature
To run the developed model in real-time, accurate
coil charging data, accurate air temperature in the fur-
nace combined with the presented mathematical model
are needed and thus provide the coil-temperature data
without measuring the temperature in the coil. Note that
the used explicit finite-difference method for the conduc-
tion calculation in the coil (eq. 4) is suitable for a real-
time calculation, the only modification of the method is
that the real-time simulation is interrupted after every
iteration until the next measured furnace-air temperature
is delivered to the model.
Unequal coil alloys, dimensions, initial temperature,
malfunctions in the furnace and burners all lead to more
or less unknown coil temperatures in the situation with-
out either the coil temperature measurements or the
real-time coil-temperature calculation. The MM obtained
coil temperatures in such a situation are crucial for
proper operator decisions and can be upgraded in the
automatic coil-temperature control system. Furthermore,
the stored real-time-calculated coil-temperatures can be
used as documentation for coil-customer claims,
research and development purposes. Real-time operation
of the model is underway.
4 CONCLUSION
A mathematical model for an Al-coil temperature
evolution during heat treatment proved to be an efficient
and multi-purpose backbone tool for advanced planning,
control and documentation of the Al-coil heat-treatment
process. Employing the mathematical model, which pro-
vides coil-temperature knowledge, offers a simple, cost-
effective and punctual tool for the determination of the
proper furnace-temperature time settings during modi-
fied Al-coil reheating conditions.
5 REFERENCES
1 A. Jakli~, T. Kolenko, B. Glogovac, B. Ke~ek, J. Jamer, Simulacijski
model ogrevanja gredic v pe~i allino, Mater. Tehnol., 38 (2004) 1–2,
33–38
2 F. Vode, A. Jakli~, R. Robi~, A. Ko{ir, F. Perko, J. Novak, Deter-
mination of the deformational energy during slab-width rolling on an
edger mill, Mater. Tehnol., 40 (2006) 2, 61–64
3 www.secowarwick.com/assets/Documents/Brochures/AP-Alumi-
nium-Annealing-Furnaces2.pdf
4 A. von Starck, A. Mühlbauer, C. Kramer (eds.), Handbook of ther-
moprocessing technologies: Fundamentals, Processes, Components,
Safety, Vulkan Verlag GmbH, Essen 2005, 557
5 R. Burns, Advanced Control Engineering, 1st ed., Butterworth-
Heinemann, 2001, 145
F. VODE et al.: MATHEMATICAL MODEL FOR AN Al-COIL TEMPERATURE CALCULATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 647–651 651

Y. H. ÇELIK: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE HOLE QUALITY ...
INVESTIGATING THE EFFECTS OF CUTTING
PARAMETERS ON THE HOLE QUALITY IN DRILLING
THE Ti-6Al-4V ALLOY
PREISKAVA VPLIVA PARAMETROV REZANJA NA KVALITETO
IZVRTINE, IZVRTANE V ZLITINO Ti-6Al-4V
Yahya Hýþman Çelik
Mechanical Engineering Department, Faculty of Engineering-Architecture, Batman University, Batman, Turkey
yahyahisman.celik@batman.edu.tr, yahyahisman@gmail.com
Prejem rokopisa – received: 2013-09-20; sprejem za objavo – accepted for publication: 2013-11-19
In this study, the effects of cutting parameters on the surface roughness, burr height, hole-diameter deviation, cutting
temperature and structure of a chip formation were investigated during the drilling of the Ti-6Al-4V alloy. For this purpose, the
Ti-6Al-4V alloy was drilled at different cutting parameters, longitudinally in the 10 mm depth with Ø = 10 mm high-speed-steel
(HSS) drills, having 90°, 118°, 130° and 140° point angles on the CNC vertical machining centre. Experiments were carried out
at the (12.5, 18.75 and 25) m/min cutting speeds and the (0.05, 0.1 and 0.15) mm/r feed rates without using the cutting fluid. As
a result, as the feed rate and the drill-point angle were increased, the surface roughness increased as well; however, as the
cutting speed increased, the surface roughness decreased. When the feed rate and drill-point angle increased, the burr height
decreased. On the other hand, an increase in the cutting speed increased the burr height. In general, an increase in the feed rate
and drill-point angle increased the hole diameters, and the hole diameters obtained were close to the nominal size when the
cutting speed was increased.
Keywords: Ti-6Al-4V, surface roughness, chip, drilling, burr height, hole diameter
V tej {tudiji so bili preiskovani vplivi parametrov rezanja na hrapavost povr{ine, vi{ino zarobka, odmike premera izvrtine,
temperaturo rezanja in strukturo nastalega izvrtka pri vrtanju zlitine Ti-6Al-4V. V ta namen je bila zlitina Ti-6Al-4V vrtana na
vertikalnem CNC obdelovalnem stroju pri razli~nih parametrih rezanja vzdol`no v globino 10 mm s svedri iz hitroreznega jekla
(HSS) premera 10 mm, ki so imeli vr{ni kot 90°, 118°, 130° in 140°. Preizkusi so bili izvr{eni pri hitrostih rezanja (12,5, 18,75
in 25) m/min in podajanju (0,05, 0,1 in 0,15) mm/r, brez uporabe hladilne teko~ine. Rezultati so pokazali, da pri pove~evanju
podajanja in ve~anju vr{nega kota svedra nara{~a tudi hrapavost povr{ine, vendar pa se pri pove~evanju hitrosti rezanja
zmanj{uje hrapavost povr{ine. ^e se pove~ujeta podajanje in velikost vr{nega kota, se zmanj{uje vi{ina zarobka. Po drugi strani
pove~anje hitrosti rezanja pove~a vi{ino zarobka. Na splo{no velja, da pove~anje podajanja in pove~anje vr{nega kota svedra
pove~ujeta premer izvrtine. Premeri izvrtin so bili blizu nazivni velikosti, ~e se je pove~ala hitrost rezanja.
Klju~ne besede: Ti-6Al-4V, hrapavost povr{ine, izvrtek, vrtanje, vi{ina zarobka, premer izvrtine
1 INTRODUCTION
Titanium and its alloys are widely used in many
fields, especially in aircraft engines and automotive parts
because they are light, exhibiting quite a good perfor-
mance, high resistance to corrosion and high strength
and being appropriate for high-temperature applica-
tions.1,2 To give the final shape to these materials pro-
duced by casting, forging and powder metallurgy,
various manufacturing methods are utilized.3 The drilling
process is one of the main manufacturing methods for
obtaining the final shape of a material.4 Nowadays, the
drilling of Ti and its alloys within the desired tolerance is
required. However, the drilling of these materials is very
difficult because of their superior features.5 During
drilling, the chip welds to the cutting tools. In addition,
the temperature of the tools and materials increase due to
a low thermal conductivity of the HSS versus the
Ti-Al-V alloys.6 These factors negatively affect the tool
wear, chip type, burr formation, surface roughness and
geometric quality.7 A correct selection of the cutting
parameters and cutting conditions is important to mini-
mize these problems and determine the ideal cutting con-
ditions.8,9 When literature studies are examined, it can be
seen that there are not many studies related to this topic.
Some studies performed on titanium alloys are given
below.
Rahim and Sasahara10 measured the tool wear, the
temperature of the workpiece and the cutting forces
occurring at various cutting parameters and different
cooling types at high-speed drilling of Ti-6Al-4V. They
found that the maximum temperature of the workpiece
occurred when it was cooled with an air blower, being
lower when using minimum quantity lubrication palm oil
(MQLPO) and minimum quantity lubrication synthetic
ester (MQLSE) and the lowest when using water. How-
ever, they reported that MQLSE led to a higher thrust
force than MQLPO affecting the material during the
cutting process. Park et al.11 measured the thrust force
and investigated the wear of tungsten carbide (WC) and
polycrystalline diamond (PCD) during the drilling of
titanium- and carbon-fibre-reinforced plastics. They
determined that PCD drills showed a lower titanium
adhesion when compared to WC drills. They also
observed a higher spindle speed, causing a significant
Materiali in tehnologije / Materials and technology 48 (2014) 5, 653–659 653
UDK 621.95:669.295 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)653(2014)
increase in the tool wear due to a higher temperature.
Pujana et al.12 applied ultrasonic vibrations on workpiece
samples while drilling the Ti-6Al-4V alloy. They investi-
gated the temperature, the chip formation and the feed
force on the drill tip by means of an infrared radiation
thermometer. They observed that a higher force was
required for cutting because of a high-temperature occur-
rence when the vibration amplitude was increased.
Shyha et al.13 drilled titanium/carbon-fiber-reinforced
plastic (CFRP)/aluminum stacks at different cutting
parameters with a coated (a CVD diamond and a hard
metal) tungsten-carbide drill. Due to out of roundness,
they measured the hole size, cylindricity, burr height,
hole-edge quality, average surface roughness (Ra), micro-
hardness (of the metallic elements) and swarf
morphology. The burr height was observed to be larger at
the hole exit (Al 7050) compared to hole entry (Ti 6–4),
while the delamination was significantly reduced by wen
machining the CFRP in the stack configuration as
opposed to the stand-alone configuration. However, the
chip formation occurring while drilling Al 7050 was
obtained using a similar cutting fluid and low cutting
parameters of the titanium alloy. Isbilir and Ghasse-
mieh14 investigated the drilling of the Ti-6Al-4V alloy
using the 3D Lagrange finite-element software. By
analysing the materials, they reported that the thrust
force, stresses and burr height of the drilled materials
increased when the feed rate was increased. Guu et al.15
analyzed the stresses in the micro-drilling of Ti and its
alloy using finite-element methods. Kývak and Þeker16
investigated the effect of the cutting parameters on the
cutting forces in drilling the Ti-6Al-4V alloy with coated
and uncoated drills under dry and wet cutting conditions.
At the end of the experimental study, they found that the
coating materials provided about a 17 % decrease in the
cutting forces. It was seen that the feed rate had a bigger
effect than the cutting speed on the change in the cutting
forces.
The main objective of metal cutting is to provide the
surface quality and the burr height along with the geome-
tric and dimensional completeness of the workpiece to
be produced economically and within the desired limits.
Nowadays, drilling the Ti-6Al-4V alloys that are used
widely is very difficult because of their excellent pro-
perties such as high resistance to corrosion, high strength
and high-temperature resistance. Therefore, in this study,
the effects of the cutting parameters, together with diffe-
rent point angles of HSS drills, on the surface roughness,
hole-diameter deviation, burr height, cutting temperature
and structure of the chip formation were investigated
during drilling a Ti-6Al-4V alloy at various combina-
tions of the feed rates and cutting speeds.
2 MATERIALS AND METHODS
2.1 Experimental study
For the experimental study, the Ti-6Al-4V alloy was
provided by Sincemat Co. Ltd. The nominal chemical
composition of the Ti-6Al-4V alloy with the size of 100
mm × 150 mm × 10 mm is given in Table 1.
In the experiments, a HUMMER VMC-1000 CNC
vertical machining centre with the maximum speed of
8000 r/min and the spindle power of 15 kW was used.
The experiments were performed using 10 mm
diameter HSS drills with the helix angle of 35°, with the
point angles of 90°, 118°, 130° and 140°. For the experi-
ments where no cutting fluids were used, the cutting
parameters are given in Table 2.
2.2 Measurement of the surface roughness
A measurement of the surface roughness is quite
important to see the effects of the cutting parameters on
the material of a machined surface. With this regard, the
surface-roughness measurements of the drilled alloy
were performed for different cutting parameters using
Taylor-Hobson’s Surtonic 3 and a surface-roughness
measuring device in accordance with ISO standards. The
measurement sampling length was chosen as 5.6 mm.
The measurement process was carried out parallel to the
axis hole. Four surface roughness values (Ra) for the
machined surfaces were obtained and then averaged.
2.3 Determining the burr height and the hole-diameter
deviation
To determine the burr height and the hole-diameter
deviation, a three-dimensional coordinate measuring
device of the SIDIO XR brand, with a 1.4 megapixel
IDIO Neo sensor, 340 mm scan area and Manfrotto
Studio Tripod was used. Before scanning with this
device, a non-destructive testing spray (BT 70) was
sprayed on the surface of the Ti-6Al-4V alloy, where the
drilling process would be applied, in order to make a
better measurement and determine the reference points
that were affixed to certain portions of the alloy to be
drilled. For the determination of the burr height and the
hole-diameter deviation, the Ti-6Al-4V alloy was
scanned with a 3-D optical scanning system. For the
determination of the effects of different cutting para-
meters and point angles of the drills on the burr height,
the Ti-6Al-4V alloy scanned with the three-dimensional
coordinate measuring device was transferred to the
PolyWorks program. This program automatically gives
minimum, mid and maximum values of the burr height at
a hole exit. The burr height was determined by taking the
arithmetic means of these values.
Y. H. ÇELIK: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE HOLE QUALITY ...
654 Materiali in tehnologije / Materials and technology 48 (2014) 5, 653–659
Table 1: Chemical composition of the Ti-6Al-4V alloy (w/%)
Tabela 1: Kemijska sestava zlitine Ti-6Al-4V (w/%)
Ti Al V Fe C N H O
89.52 6.1 4.1 0.056 0.019 0.025 0.08 0.1
Table 2: Cutting parameters and their values
Tabela 2: Parametri rezanja in njihove vrednosti
Parameters Values
Cutting speed (m/min) 12.5, 18.75 and 25
Feed rate (mm/r) 0.05, 0.10 and 0.15
3 RESULTS AND DISCUSSION
3.1 Surface roughness
The surface roughness of the holes occurring as a
result of drilling the Ti-6Al-4V alloy is given in Figures
1 and 2. As seen in these figures, the surface roughness
increased when the feed rate was increased and it
decreased when the cutting speed was increased. In
addition, it was observed that the surface roughness
increased with an increase in the point angle. While the
lowest surface roughness was obtained with the 90°
point-angle drills, 0.05 mm/r feed rate and 25 m/min
cutting speed, the largest surface roughness was obtained
with the 140° point-angle drills, 0.15 mm/r feed rate and
12.5 m/min cutting speed. Similar results were found by
Kim et al.17 and Sharif and Rahim.18
During the machining at a high cutting speed, the
cutting temperature increases due a small contact length
between the tool and the workpiece. This could be due to
a decrease in the value of the coefficient of friction,
resulting in a low friction at the tool-workpiece interface.
These factors could contribute to an improvement in the
surface-roughness values.19 In addition, as the cutting
speed increases, more heat is generated, thus softening
the workpiece material, which, in turn, improves the sur-
face roughness. However, a low cutting speed may lead
to the formation of a built-up edge and, hence, deterio-
rates the machined surface. The investigation revealed
that at a high feed rate the surface roughness is poor,
probably due to distinct feed marks produced at the high
feed rate.10
3.2 Burr height
In Figure 3, the burr heights obtained with the Poly-
Work program for the alloy scanned with the three-
dimensional coordinate measuring device are shown. In
addition, the burr heights obtained depending on feed
rates, cutting speeds and point angles are given in Fig-
ures 4 and 5.
The burr height occurring at the hole exit at the low
feed rate was found to be larger than the burr height
occurring at the hole exit at the high feed rate. However,
when the cutting speed was increased, the burr height at
the hole exit also increased. For this reason, the biggest
burr height was obtained at the high cutting speed and
low feed rate. The burr that occurred in the drilling
process carried out with large point-angle drills was
found to be smaller. Similar results were found by Kim
et al.20 and Dornfeld et al.21
Y. H. ÇELIK: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE HOLE QUALITY ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 653–659 655
Figure 4: Effect of the feed rate on the burr height
Slika 4: Vpliv hitrosti podajanja na vi{ino zarobka
Figure 2: Effect of the cutting speed on the surface roughness
Slika 2: Vpliv hitrosti rezanja na hrapavost povr{ine
Figure 1: Effect of the feed rate on the surface roughness
Slika 1: Vpliv hitrosti podajanja na hrapavost povr{ine
Figure 3: PolyWork burr measurement
Slika 3: Merjenje zarobka s PolyWork-om
The cutting tool, processed materials, cutting para-
meters and the cooling fluid affect the burr formation.22,23
Especially the drilling of metal materials causes the
formation of burrs, resulting from the plastic deforma-
tion of workpiece materials both at the entrance and the
exit of a hole.7 ISO 13715 standards define the burr di-
mensions that deviate from the ideal geometry. However,
the measurement of chip characteristics is rather difficult
and takes time.24 Therefore, the easiest way of characte-
rising chips is to measure their height and thickness.25
The burr height and thickness are very important for
deburring an unwanted formation off the workpiece.
These unwanted burrs are harder than the workpiece.
Therefore, deburring is quite costly forcing us to use
extra tools. This causes the costs to rise by up to 30 %,
especially for aircraft engines. For automobiles, on the
other hand, the cost rise varies between 15 % and 20 %.26
3.3 Hole-diameter deviation
A hole-diameter deviation is usually a result of the
effects such as deflection, vibration, wear and lack of
lubrication. On the other hand, a deviation from the
circularity represents the fluctuations on the surface, and
it is defined as the difference between the largest and the
smallest radius measured from the center point. How-
ever, there are various ways of determining the center of
a hole. In this study, the hole diameter and the circularity
were determined with the coordinates taken at different
points. These measurements used in the evaluation of the
hole diameter and circularity were transferred directly to
the PolyWorks program (Figure 6). Thus, the hole-dia-
Y. H. ÇELIK: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE HOLE QUALITY ...
656 Materiali in tehnologije / Materials and technology 48 (2014) 5, 653–659
Figure 6: PolyWork hole-diameter deviation
Slika 6: Merjenje odmika premera izvrtine s PolyWork-om
Figure 5: Effect of the cutting speed on the burr height
Slika 5: Vpliv hitrosti rezanja na vi{ino zarobka
Figure 9: Effect of the drill-point angle on the hole diameter
Slika 9: Vpliv vr{nega kota svedra na premer izvrtine
Figure 8: Effect of the cutting speed on the hole diameter
Slika 8: Vpliv hitrosti rezanja na premer izvrtine
Figure 7: Effect of the feed rate on the hole diameter
Slika 7: Vpliv hitrosti podajanja na premer izvrtine
meter deviation was detected at the entrance, the middle
and the exit of a hole.
The evaluation of the hole diameter was performed
by comparing the values obtained with the cutting tool
depending on different cutting parameters. The diameters
affected by the feed rate, the cutting speed and the point
angle are given in Figures 7 to 9. With an increase in the
feed rate and drill-point angle, the diameter also in-
creased significantly at the entrance, middle and exit of a
hole. On the other hand, the diameter deviation de-
creased with an increase in the cutting speed. In addition,
it was found that the deviation at the entrance of a hole
was smaller than that at the exit.
3.4 Cutting temperature
During the drilling process, 90 % of work is conver-
ted to heat as a result of the plastic deformation.27
Therefore, a very high temperature occurs in the drilling
zone. This temperature affects a specific region of the
chip, tools and workpiece.28 In connection with the ther-
mal properties of the workpiece and cutting, either the
cutting tool or the workpiece is affected by this. Since Ti
and its alloys, which are about 1/6 of steels,29 have low
thermal properties, a great deal of heat, as much as 80 %,
is absorbed by the tool.30 50 % to 60 % of the heat gene-
rated during the drilling of steel is absorbed by the tool.31
When machining Ti and its alloys, the tool reaching a
high temperature wears quickly because of the high cutt-
ing temperature and a strong adhesion between the tool
and the workpiece; and the high stresses developed at the
cutting edge of the tool may cause a plastic deformation
and accelerate the tool wear.32
The images of the Ti-6Al-4V alloy taken by a ther-
mal camera at various cutting parameters are given in
Figure 10. As seen in this figure, a low cutting speed
leads to a low temperature because of the interaction bet-
ween the material and the tool.
3.5 Chip types
A chip formation of titanium alloys during the drill-
ing process is quite difficult when compared with the
other metals. There are three types of chip. These are the
continuous chip, the continuous chip with a build-up
edge, and the discontinuous chip.33 When the feed rate
was low, the Ti chip was long and continuous, and when
the feed rate was increased, it became shorter and
stiffer.17 However, as the cutting speed was raised, the
chip-serration frequency was enhanced.34
A few chip samples were taken from each drilling
process to detect the effects of the cutting parameters and
drill-point angle on them. From these chips, it was
observed that the chips were formed under difficult
conditions and that high cutting force and temperature
took place; depending on these factors, the shapes of the
Y. H. ÇELIK: INVESTIGATING THE EFFECTS OF CUTTING PARAMETERS ON THE HOLE QUALITY ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 653–659 657
Figure 11: Some chip shapes: a) continuous chip, b) continuous chip with a build-up edge, c) discontinuous chip
Slika 11: Nekaj oblik izvrtkov: a) kontinuirni izvrtek, b) kontinuirni izvrtek s poudarjenim robom, c) razlomljeni izvrtek
Figure 10: Temperature measurements with thermal cameras: a) low
cutting speed, b) high cutting speed
Slika 10: Merjenje temperature s toplotno kamero: a) majhna hitrost
rezanja, b) velika hitrost rezanja
chips were irregular. At a low cutting speed, feed rate
and drill-point angle, the chip was observed to be ductile
and continuous (Figure 11a). The chip was hardened
and it became fragile with an increase in the cutting
speed, feed rate and drill-point angle (Figure 11b). How-
ever, with an increase in the cutting speed, ridges
appeared on the chip surface (Figure 11c).
The effect of the temperature on the chip hardening is
very important due to the friction. Due to the tempera-
ture effect, the chip had a dark colour.
Acknowledgement
The author would like to thank Ýhsan Pilatin, an
experienced lecturer at Batman University for his con-
tributions and corrections to the paper.
4 CONCLUSION
In this study, the parameters such as surface
roughness, burr height, hole-diameter deviation, cutting
temperature and structure of a chip formation were
investigated during the drilling of the Ti-6Al-4V alloy
under different feed rates, cutting speeds and drill-point
angles using HSS drills with 90°, 118°, 130 ° and 140°
point angles and the following conclusions were reached:
• When the feed rate and drill-point angle were
increased, the surface roughness increased as well;
when the cutting speed increased, the surface
roughness decreased.
• When the feed rate and drill-point angle were
increased, the burr height decreased, but an increase
in the cutting speed increased the burr height.
• Generally, when the feed rate and drill-point angle
were increased, the hole diameter increased; when
the cutting speed was increased, the diameter was
close to the nominal value.
• For all the drill-point angles, the cutting speed and
feed rate combinations, the diameters obtained were
larger than the nominal values.
• For the low cutting speed and feed rate, the chip form
was found to be more regular and ductile.
• With an increase in the cutting speed and feed rate,
the chip formation displayed a shift from a ductile
structure towards a more rigid and fragile form.
• An increase in the cutting speed caused the material
to overheat due to the friction and the thermal
properties of Ti-6Al-4V.
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Materiali in tehnologije / Materials and technology 48 (2014) 5, 653–659 659

S. PRABAGARAN et al.: INFLUENCE OF GRAPHITE ON THE HARDNESS AND WEAR BEHAVIOR ...
INFLUENCE OF GRAPHITE ON THE HARDNESS AND
WEAR BEHAVIOR OF AA6061–B4C COMPOSITE
VPLIV GRAFITA NA TRDOTO IN VEDENJE KOMPOZITA
AA6061–B4C PRI OBRABI
Subramaniam Prabagaran1, Govindarajulu Chandramohan2, Palanisamy Shanmughasundaram3
1Karpagam University, Department of Mechanical Engineering, Coimbatore-641021, India
2P.S.G. Institute of Technology and Applied Research, Coimbatore-641042, India
3Karpagam University, Department of Automobile Engineering, Coimbatore-641032, India
prabagaran.s@karpagam.com
Prejem rokopisa – received: 2013-09-21; sprejem za objavo – accepted for publication: 2013-12-10
Dry-sliding-wear behavior of AA6061, AA6061-B4C composite and AA6061-B4C-Gr hybrid composite was investigated by
employing a pin-on-disc wear-test rig. Hardness tests were also carried out. Graphite was used as a solid lubricant since it is a
soft, slippery and greyish-black substance. Because of the cleavage (crystal) loose interlamellar coupling, graphite has good
lubricating properties. A comparative analysis was made on the hardness and wear behavior of AA6061, AA6061-B4C
composite and AA6061-B4C-Gr hybrid composite. Boron carbide is known as a robust material having a high hardness.
Worn-out surfaces of the wear specimens after the wear tests were examined with a scanning electron microscope to study the
morphology of the worn surfaces. Energy dispersive spectroscopy (EDS) was employed to identify the oxides formed on the
worn surfaces of the AA6061-B4C and AA6061-B4C-Gr composites after the wear test. It was found that a mutual transfer of the
material between the wearing Al-alloy and the steel counterface occurred as the load increased. Oxidative wear occurred at low
applied loads and a high velocity, whereas delamination and adhesive wear occurred at a high load and high sliding velocity.
Keywords: boron carbide, dry sliding, graphite particles, pin-on-disc, wear resistance
Preiskovano je bilo vedenje AA6061, AA6061-B4C-kompozita in hibridnega kompozita AA6061-B4C-Gr pri suhem drsenju z
uporabo naprave “pin on disc”. Kot trdno mazivo je bil uporabljen grafit, ki je mehka, spolzka in sivo-~rna snov. Zaradi cepljive
{ibke medlamelarne povezave (v kristalu) ima grafit dobre mazalne lastnosti. Izvr{ene so bile analize primerjave trdote in
vedenja pri obrabi AA6061, AA6061-B4C-kompozita in hibridnega kompozita AA6061-B4C-Gr. Borov karbid je poznan kot
zdr`ljiv material z veliko trdoto. Po preizkusu obrabe je bila z elektronskim vrsti~nim mikroskopom izvr{ena preiskava
morfologije obrabljene povr{ine vzorcev, da bi ugotovili nastale okside na obrabljeni povr{ini kompozitov AA6061-B4C in
AA6061-B4C-Gr. Ugotovljeno je bilo, da se pri nara{~anju obremenitve pojavi vzajemen prenos materiala med Al-zlitino, ki se
obrablja, in jeklom. Oksidativna obraba se je pojavila pri majhnih obremenitvah in velikih hitrostih, medtem ko se je obraba
zaradi delaminacije in adhezije pojavila pri velikih obremenitvah in veliki hitrosti drsenja.
Klju~ne besede: borov karbid, suho drsenje, grafitni delci, "pin-on-disc", odpornost proti obrabi
1 INTRODUCTION
Metal-matrix composites have received a lot of com-
mercial attention due to their enhanced mechanical
properties, wear resistance and low coefficient of thermal
expansion.1 Metal-matrix composites have many appli-
cation areas such as automotive and ballistic industries,
infrastructure, space and air vehicles, under-water vehi-
cles and deep-ocean equipment. In addition, metal-ma-
trix composites have other advantageous characteristics
such as good strength-to-weight ratio, high specific stiff-
ness, high hardness, high plastic-flow strength, good
thermal expansion, thermal stability, creep resistance,
and good oxidation and corrosion resistance.2
Boron carbide exhibits excellent physical and mecha-
nical properties. Boron carbide is a low-density ceramic
with a high hardness and Young’s modulus which make
it a valuable candidate for engineering applications.
Rama Rao and Padmanabhan3 reported that an addition
of boron carbide decreases the density of composites and
increases the hardness. Gómez et al.4 reported that the
hardness and strength of composites increased together
with the volume fraction of reinforcement, reaching its
maximum value of 10 % B4C. Wear is a removal of a
material from one or both of the two solid surfaces in a
solid-state contact. An addition of hard ceramic particles
improves the wear resistance of a matrix material. The
wear rate is associated with the sliding velocity, normal
load, particle size, hardness, particle volume fraction and
particle homogeneity.5 Jinfeng et al.6 observed that with
an addition of graphite, the friction coefficient of Al-SiC
composites decreases and the wear resistance is signifi-
cantly increased. Ames and Alpas7 reported that Al-Gr
composites had a higher wear resistance than Al-SiC
composites.
A solid lubricant provides protection from damage
during relative movement between the sliding elements
and reduces the wear. Miyazaki et al.8 analyzed the
mechanical behavior of graphite-and-boron-carbide
composites made with the hot-pressing method. They
reported that the strength of a composite increased with
an increase in the boron-carbide content. Many research
works were carried out to investigate the sliding-wear
behavior of aluminum MMCs reinforced with different
Materiali in tehnologije / Materials and technology 48 (2014) 5, 661–667 661
UDK 66.017:620.178.1:539.92 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)661(2014)
types of reinforcements such as SiC, Al2O3 and B4C. But
limited studies have been carried out on hybrid metal-
matrix composites.
Shorowordi et al.9 studied three aluminium metal-
matrix composites containing the reinforcing particles of
B4C, SiC and Al2O3 (volume fractions,  = 0–20 %) that
were manufactured with the stir-casting method followed
by hot extrusion. They reported that the B4C-reinforced
Al-composite seemed to exhibit a better interfacial bond-
ing compared to the other two composites.
Stir casting appears to be the best method for the
production of metal-matrix composites compared to the
other processing techniques because of its simplicity,
allowing an economical large-scale production. The aim
of this study is to analyze the influence of graphite parti-
cles on the hardness and wear behavior of AA6061-B4C
composites fabricated using a two-stage stir-casting
method. The wear tests were carried out by employing a
pin-on-disc wear-testing rig. Scanning Electron Micro-
scopy (SEM) was employed to study the microstructures
of the composites and the morphologies of the worn
surfaces of the composites. Energy Dispersive Spectro-
scopy (EDS) was used to characterize the Mechanically
Mixed Layer (MML) formed on a worn surface during
the sliding wear and to elucidate its influences on the
wear behavior of the composites.
2 EXPERIMENTATION
2.1 Experimental description
In this study, a two-stage stir-casting method was
employed to fabricate AA6061-B4C and AA6061-
B4C-Gr composites. Brinell hardness testing equipment
was employed to measure the hardness of the AA6061
alloy and the composites. The dry-sliding wear behavior
of AA6061 and the composites was tested with a
pin-on-disc wear-testing rig. Scanning Electron Micro-
scopy (SEM) was employed to study the microstructures
and morphologies of the worn surfaces of AA6061,
AA6061-B4C and AA6061-B4C-Gr composites.
2.2 Specimen preparation
In this study, we used AA6061 as the matrix material,
B4C particles with the average size of 20–50 μm as the
reinforcement and graphite particles of 20 μm as the
solid lubricant. The composition of AA6061 is presented
in Table 1. The stir-casting setup is shown in Figure 1.
Table 1: Composition of AA6061 (mass fractions, w/%)
Tabela 1: Sestava AA6061 (masni dele`i, w/%)
Element Si Fe Cu Mn Mg Cr Ti Ca Al
% 0.359 0.221 0.219 0.032 0.901 0.053 0.141 0.013 Bal
A two-stage stir-casting route was adopted to fabri-
cate the composites. Both B4C and graphite particulates
were preheated at 300 °C in a separate muffle furnace.
AA6061 was charged into the crucible and heated up to
650 °C in order to completely melt the aluminium and
then the stirring was done at 300 r/min. During the
stirring, degassing tablets were added to drive away the
entrapped gases from the melt. The stirring was carried
out for 3 min. During the stirring, preheated w = 10 %
B4C and 3 % graphite particulates were added. The melt
temperature was brought down to 575 °C to reach a
semi-solid state. At this stage the stirring was done for 3
min. The composite slurry was again reheated to the
temperature of 650 °C and stirred at 300 r/min for 3 min.
Finally, the composite slurry was poured into a steel die
cavity of 90 mm × 90 mm × 7 mm to solidify. The
melting was done in an electric resistance furnace (2 kW
– 1 kg capacity). The temperatures were measured with a
thermocouple with an accuracy of ± 3 K.
2.2.1 Hardness-test specimens
The poured samples of 1) the AA6061 alloy, 2) com-
posite AA6061-B4C and 3) AA6061-B4C-Gr were
machined on their four sides to obtain 50 mm squares.
All the edges and corners were made blunt/rounded to
ensure safety while handling.
2.2.2 SEM analysis (microstructure)
With electric discharge machining, three different
samples were machined to rectangular pieces of 10 mm
× 20 mm with a thickness of 6 mm. Then these samples
were mounted using a heat-conducting, quick-setting
epoxy resin. After the mounting each sample was manu-
ally polished with a 1200-grade silicon-carbide emery
sheet with the help of a diamond paste of 7 μm until
achieving a mirror-like finish. The polished surfaces
were then etched with a 2 % nitric acid and 98 % alcohol
etching solution for 10 s.
S. PRABAGARAN et al.: INFLUENCE OF GRAPHITE ON THE HARDNESS AND WEAR BEHAVIOR ...
662 Materiali in tehnologije / Materials and technology 48 (2014) 5, 661–667
Figure 1: Two-stage stir-casting setup
Slika 1: Dvostopenjska naprava za vme{avanje delcev
2.2.3 Wear-analysis specimens
Pin specimens with a width of 6 mm and a height of
30 mm were prepared with EDM (Electric Discharge
Machining). The two ends were cut with high flatness
accuracy so that the pin face was thoroughly in line with
the disc, held perpendicular to the disc surface during the
wear test. Moreover, the wear testing equipment had a
split-type holder to ensure a proper alignment throughout
the test run. Prior to the wear testing, the specimens were
polished with an abrasive paper of silicon carbide of
grade 600 followed by grade 1000, then cleaned with
ethanol and dried.
2.2.4 EDS analysis specimens
The samples prepared for the SEM analysis were also
used for the EDS analysis without any modifications.
2.3 Hardness test
The hardness test was performed on AA6061 and the
composite specimens of AA6061-B4C and AA6061-
B4C-Gr using the Brinell hardness testing equipment
with a 2.5 mm steel-ball diameter at a load of 1839.4 N.
The loading time was 30 s. Three readings were taken
for each specimen to eliminate the possibility of segre-
gation and the mean value was considered.
2.4 Microstructure analysis
A scanning electron microscope was employed to
study the distribution of boron-carbide and graphite
particles in the Al-matrix. The bonding quality between
the particulates and the matrix was also studied.
2.5 Dry-sliding wear test
The wear tests were carried out under varying loads
of (5, 10, 15 and 20) N at the sliding velocities of 1 m/s
and 3 m/s by employing a pin-on-disc wear-testing rig
(DUCOM, TR-20LE) with a data-acquisition system
shown in Figure 2. Each test was conducted for at least
30 min.
The main parts of the apparatus are a variable-speed
electric motor with a steel disc attached to it and a lever
arm to which the weight is added. The wear loss of the
sample pins was measured in terms of the height loss in
microns with the accuracy of 1.0 μm. A track diameter of
100 mm was selected for the analysis. The rotating disc
was made of the EN 31 steel having the hardness of 62
HRC. The wear tests were carried out at a room
temperature of 30 °C and a relative humidity of 60 % for
30 min, the contact pressure on the disc being 0.14
N/mm2 (at 5 N load), 0.28 N/mm2 (at 10 N load), 0.42
N/mm2 (at 15 N load) and 0.56 N/mm2 (at 20 N load).
Two samples were tested for each condition. A SEM
examination was carried out on the worn surfaces of the
specimens in order to understand the wear mechanism
under various test conditions.
2.6 Energy Dispersive Spectroscopy (EDS) analysis
An EDS analysis was done on the worn surfaces of
the AA6061 alloy and AA6061-B4C-Gr composite (Jeol)
primarily to verify the presence of oxides on the worn
surfaces. The data output plots the original spectrum
showing the number of X-rays collected at each energy
level and mapping the element distributions over the
areas of interest.
3 RESULTS AND DISCUSSION
3.1 Microstructural analysis
The microstructure of the AA6061-B4C-Gr hybrid
composite (Figure 3) was fabricated with the stir-casting
method. It can be seen that the boron-carbide and
graphite particles are distributed uniformly, bonding well
with the aluminum matrix. The interface between the
Al-matrix, boron-carbide and graphite particles is clean
allowing a strong interfacial bonding. No agglomeration
of the particles was observed in the composite.
S. PRABAGARAN et al.: INFLUENCE OF GRAPHITE ON THE HARDNESS AND WEAR BEHAVIOR ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 661–667 663
Figure 3: Microstructure of AA6061-B4C-Gr hybrid composite
Slika 3: Mikrostruktura hibridnega kompozita AA6061-B4C-Gr
Figure 2: Pin-on-disc wear-testing rig
Slika 2: Naprava "pin on disc" za preizku{anje obrabe
3.2 Hardness analysis
It can be observed from Figure 4 that the hardness of
the Al-matrix increased from 50 BHN to 83 BHN due to
the addition of w = 10 % of boron-carbide particles. This
reveals an increase in the hardness due to the addition of
boron carbide. On the other hand, the hardness of
AA6061-B4C composites decreased from 83 BHN to 66
BHN when graphite particles were incorporated. This
can be elucidated with the fact that graphite has a lower
hardness than B4C particles. The hardness of Al–10 %
B4C–3 % graphite was by about 21 % lower than the
hardness of the Al–10 % B4C composite. A similar
observation was made by Hassan et al.10
3.3 Dry-sliding wear test
Typical curves of the wear loss of matrix AA6061 at
the 3 m/s sliding velocity and of hybrid composite
AA6061-B4C-Gr at the sliding velocity of 3 m/s and the
constant load of 20 N are presented in Figures 5 and 6,
respectively. It can be observed that the wear of hybrid
composite AA6061-B4C-Gr is lower by 28 % when
compared with matrix AA6061. This is due to the
presence of the B4C and graphite particles. This can also
be attributed to a better interfacial bonding between Al
and the B4C particles. The volume of the wear debris
increases with the increasing load, resulting in a greater
loss of the material. The wear resistance of the Al-B4C
composites tends to decrease when the sliding velocity is
increased from 1 m/s to 3 m/s at the load of 20 N. In
general, a higher sliding velocity generates a higher
frictional heat which increases the wear.
It can be observed from Figures 7 and 8 that the wear
loss increased with the increasing load. The volume of
the wear debris increases with the increasing load,
resulting in a greater loss of the material. The wear
resistance of the AA6061-B4C composites tended to
decrease when the sliding velocity was increased from 1
m/s to 3 m/s at the load of 20 N. In general, a higher
S. PRABAGARAN et al.: INFLUENCE OF GRAPHITE ON THE HARDNESS AND WEAR BEHAVIOR ...
664 Materiali in tehnologije / Materials and technology 48 (2014) 5, 661–667
Figure 5: Typical curve of the wear-loss pattern of AA6061 at a load
20 N and 3 m/s sliding velocity for a duration 30 min
Slika 5: Zna~ilna krivulja obrabe AA6061 pri obremenitvi 20 N,
hitrosti drsenja 3 m/s in trajanju 30 min
Figure 6: Typical curve of the wear-loss pattern of AA6061-B4C-Gr
at 20 N and 3 m/s sliding velocity for a duration 30 min
Slika 6: Zna~ilna krivulja obrabe AA6061-B4C-Gr pri obremenitvi
20 N, hitrosti drsenja 3 m/s in trajanju 30 min
Figure 4: Hardness of AA6061, Al–10 % B4C and Al–10 % B4C–3 %
Gr (mass fractions, w/%)
Slika 4: Trdota AA6061, Al–10 % B4C in Al–10 % B4C–3 % Gr
(masni dele`i, w/%)
Figure 7: Variation of the wear as a function of normal loads and
sliding velocities for AA6061 and AA6061-B4C composite
Slika 7: Spreminjanje obrabe v odvisnosti od obremenitve in hitrosti
drsenja za kompozita AA6061 in AA6061-B4C
sliding velocity generates a higher frictional heat which
increases the wear.
Figure 8 shows that the wear loss of AA6061-B4C is
higher than that of the AA6061-B4C-Gr composite
irrespective of the load and speed. This could be due to
the soft nature of graphite particles acting as solid
lubricants with a layer lattice lamella crystal structure
consisting of hexagonal rings forming thin parallel
planes (graphenes). Graphenes are bonded to each other
with weak van der Waals forces. The layered structure
allows a sliding movement of the parallel planes, hence,
reducing the frictional forces between the pin and the
counter disc. This results in a reduction of the wear of
the AA6061-B4C-Gr composite which is 15 % lower
than that of the AA6061-B4C composite at the 3 m/s
sliding velocity and the 20 N load. A similar observation
was noticed by Shanmughasundaram et al.11 for
AA6061-Gr (w = 0–7.5 %) composites.
3.4 Morphology of worn surfaces
Figures 9 and 10 show the wear-track morphology of
AA6061 and the AA6061-B4C composite, respectively,
at the normal load of 20 N and the 3 m/s sliding velocity.
The worn surface of aluminium with large ploughing
grooves is shown. The AA6061 material is much softer
than the carbon-steel disc; the asperities on the steel
counter face pierce to a larger depth. A larger plastic
deformation is seen on the worn surface. The worn sur-
face of the AA6061-B4C composite (Figure 10) shows
that the extent of the material removal is not as large as
in the case of the AA6061 matrix (Figure 9). The wear
grooves are smaller along the sliding direction due to the
incorporation of boron carbide.
The worn surface of the AA6061-B4C-Gr composite
subjected to the 20 N load at the 1 m/s sliding velocity is
shown in Figure 11. It shows that the extent of the
material removal is not as large as in the case of the
AA6061-B4C composite. Graphite is a solid lubricant
and its particles are soft, slippery, greyish-black, having
a layer lattice lamella crystal structure. Its layered
structure allows a sliding movement of the parallel
S. PRABAGARAN et al.: INFLUENCE OF GRAPHITE ON THE HARDNESS AND WEAR BEHAVIOR ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 661–667 665
Figure 10: SEM micrograph of the worn surface of the AA6061-B4C
composite at the normal load of 20 N and 3 m/s sliding velocity
Slika 10: SEM-posnetek obrabljene povr{ine AA6061-B4C kompozita
pri obremenitvi 20 N in hitrosti drsenja 3 m/s
Figure 11: SEM micrograph of the worn surface of the AA6061-
B4C-Gr composite at the normal load of 20 N and 1 m/s sliding
velocity
Slika 11: SEM-posnetek obrabljene povr{ine kompozita AA6061-
B4C-Gr pri obremenitvi 20 N in hitrosti drsenja 1 m/s
Figure 9: SEM micrograph of the worn surface of AA6061 at the
normal load of 20 N and 3 m/s sliding velocity
Slika 9: SEM-posnetek obrabljene povr{ine AA6061 pri obremenitvi
20 N in hitrosti drsenja 3 m/sFigure 8: Variation of the wear as a function of normal loads and
sliding velocities for the AA6061-B4C and AA6061-B4C-Gr
composites
Slika 8: Spreminjanje obrabe v odvisnosti od obremenitve in hitrosti
drsenja za kompozita AA6061-B4C in AA6061-B4C-Gr
planes. Hence, it reduces the friction between the pin and
the disc, thus, reducing the wear. This observation proves
that the wear loss of the hybrid composite of AA6061-
B4C-Gr is lower compared to the AA6061-B4C com-
posite. The SEM micrograph does not indicate any
delamination. The sliding marks (Figure 12) obtained at
the 3 m/s velocity are relatively higher compared with
the low velocity (1 m/s) when viewed at the same magni-
fication. This can be noticed when the sliding velocity
increases from 1 m/s to 3 m/s. The morphologies of the
worn surfaces gradually change from fine scratches to
grooves. This shows that the transition from mild wear to
severe wear takes place with an increase in the sliding
velocity at a higher load. Generally, aluminium and
reinforcement particles react with the oxygen in the air,
forming iron oxides during the wear and a Mechanically
Mixed Layer (MML) on the worn surface. This oxide
film tends to deteriorate at the higher sliding velocity due
to a higher frictional heat and increases the wear. It can
be concluded that the SEM observation of the worn
surfaces of AA6061 and the composites validate the
results of the wear loss.
3.5 Energy Dispersive Spectroscopy (EDS) analysis
When the sliding velocity is increased to 3 m/s at the
low load (5 N), the surface of the composite pin reacts
with the oxygen and forms an oxide layer due to a higher
frictional heat. It reduces the direct metallic contact
between the sliding surfaces resulting in a lower wear
rate of the composite. As can be seen from Figure 13,
the peak of the oxide which is the main constituent of
MML is clearly detected on the worn surface of the
composite. It indicates that the oxidative wear is
predominant at the low load (5N) and high sliding
velocity (3 m/s). Moreover, a negligible amount of
oxides was observed on the worn surface of AA6061.
Figure 14 shows the EDAX spectrum of MML for
the Al–10 % B4C–3 % Gr composite when tested at 20
N, 3 m/s. It can be observed that the amount of the oxi-
des present on the worn surface tends to decrease. It can
also be noted from Figure 15 that a considerable amount
of iron is transferred from the counter steel disc to the
composite pin. However, broken and uneven oxide seg-
ments increase the wear. Hence, MML failed to sustain
S. PRABAGARAN et al.: INFLUENCE OF GRAPHITE ON THE HARDNESS AND WEAR BEHAVIOR ...
666 Materiali in tehnologije / Materials and technology 48 (2014) 5, 661–667
Figure 14: EDAX spectrum of MML for the Al–10 % B4C–3 % Gr
composite tested at 20 N, 3 m/s
Slika 14: EDAX-spekter MML za kompozit z Al–10 % B4C–3 % Gr,
preizku{en pri 20 N in 3 m/s
Figure 15: Weight percentage of Al, Fe and oxides as a function of
the load and sliding velocity on the worn surface of the Al–10 %
B4C–3 % Gr composite against the counter steel obtained with EDS
Slika 15: Masni dele`i Al, Fe in oksidov, ugotovljeni z EDS, v
odvisnosti od obremenitve in hitrosti drsenja na obrabljeni povr{ini
Al-kompozita z 10 % B4C in 3 % Gr v paru z jeklom
Figure 13: EDAX spectrum of MML for the Al–10 % B4C–3 % Gr
composite tested at 5 N, 3 m/s
Slika 13: EDAX-spekter MML za kompozit z Al–10 % B4C–3 % Gr,
preizku{en pri 5 N in 3 m/s
Figure 12: SEM micrograph of the worn surface of the AA6061-
B4C-Gr composite at the normal load of 20 N and 3 m/s sliding
velocity
Slika 12: SEM-posnetek obrabljene povr{ine kompozita AA6061-
B4C-Gr pri obremenitvi 20 N in hitrosti drsenja 3 m/s
under the high load and velocity during the sliding. A
higher sliding velocity increases the interface tempera-
ture and causes a local yielding, thereby the wear
mechanism changes into the delamination wear. This
behavior is termed as severe wear behavior, in which the
material removal occurs at a higher rate. The transition
from mild to severe wear is associated with the existence
of delamination and adhesion, which are the primary
wear mechanisms at the higher load and sliding velocity.
4 CONCLUSIONS
AA6061-B4C and AA6061-B4C-Gr composites were
successfully fabricated by employing the stir-casting
method. A SEM analysis revealed that boron-carbide and
graphite particles are distributed uniformly in the
aluminium matrix. The AA6061-B4C composite had a
higher hardness compared to AA6061. The wear resi-
stance of the AA6061-B4C-Gr hybrid composite and
Al-B4C composite increase steadily with the sliding load
and velocity. The wear resistance of the AA6061-B4C-Gr
hybrid composite is higher than that of the AA6061-B4C
composite and much higher than that of the AA6061
matrix. The formation of the oxides at the interface plays
a significant role in reducing the wear rate. The oxides
and reinforcing particles form a mechanically mixed
layer (MML) appearing on the worn surface of the
composite pin and enhancing the wear resistance. Hence,
it can be concluded that graphite particles reduce the
wear when included in an aluminium alloy or in an
AA6061-B4C composite.
Acknowledgement
I am grateful to the Faculty of Engineering, Depart-
ment of Automobile Engineering, Karpagam University,
Coimbatore, for providing the facilities for a successful
completion of this project.
5 REFERENCES
1 S. Balasivanandha Prabu, L. Karunamoorthy, S. Kathiresan, B. Mo-
han, Journal of Materials Processing Technology, 171 (2006),
268–273
2 A. A. Cerit, M. B. Karamiº, F. Nair, K. Yildizli, Tribology in
industry, 30 (2008), 3–4
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and Biomaterials Applications, 2 (2012) 3, 15–18
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of Composite Materials, 43 (2009) 9, 987–995
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Materials and Engineering, 38 (2009) 11, 1894–1898
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site Materials, 41 (2007) 4, 453–465
11 P. Shanmughasundaram, R. Subramanian, Advances in Materials
Science and Engineering, (2013), article ID 216536
S. PRABAGARAN et al.: INFLUENCE OF GRAPHITE ON THE HARDNESS AND WEAR BEHAVIOR ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 661–667 667

P. BÍLEK et al.: TRIBOLOGY OF CrAg7N COATINGS DEPOSITED ON VANADIS 6 LEDEBURITIC TOOL STEEL
TRIBOLOGY OF CrAg7N COATINGS DEPOSITED ON
VANADIS 6 LEDEBURITIC TOOL STEEL
TRIBOLOGIJA PREVLEK CrAg7N NA LEDEBURITNEM
ORODNEM JEKLU VANADIS 6
Pavel Bílek, Peter Jur~i, Mária Hudáková, ¼ubomír ^aplovi~, Michal Novák
Institute of Materials Science, Faculty of Materials Science and Technology in Trnava, Slovak University of Technology in Bratislava,
Paulínská 16, 917 24 Trnava, Slovak Republic
pavel-bilek@email.cz
Prejem rokopisa – received: 2013-09-25; sprejem za objavo – accepted for publication: 2013-11-19
Samples made from Vanadis 6 PM ledeburitic tool steel were surface machined, ground and mirror polished. Prior to the
deposition, they were heat treated to a hardness of 60 HRC. The CrAg7N coating was deposited with the magnetron-sputtering
technique, using pure-Cr and Ag targets in a composite low-pressure nitrogen/argon atmosphere and at a temperature of 500 °C
in the Hauzer Flexicoat 850 device. Tribological testing using a pin-on-disc apparatus was carried out at ambient and elevated
temperatures: (300, 400 and 500) °C, respectively. Al2O3, 100Cr6 and CuZn balls were used as the counterparts. The wear tracks
after the pin-on-disc testing were analyzed with scanning electron microscopy and a microanalysis. The experiments have
shown a strong dependence of tribological parameters on the temperature. The friction coefficient of CrN-Ag against the 100Cr6
ball at ambient temperature was μ = 0.56. Tribological sliding tests of this coating system against the alumina balls indicate a
decrease in the friction coefficient due to the increasing temperature. At ambient temperature μ = 0.68 and its minimum
occurred at the temperature of 400 °C, μ = 0.24. This is attributed to the diffusion of Ag particles to the sliding top surface at
elevated temperature. The testing against the CuZn brass ball generally gave a lower friction coefficient at ambient temperature,
μ = 0.39. In contrast, the friction coefficient slightly increased with the increasing temperature and was practically constant at
elevated temperatures, ranging between μ = 0.43–0.48.
Keywords: Cr-V ledeburitic steels Vanadis 6, PVD, chromium nitride with silver, pin-on-disc, friction coefficient
Vzorci, izdelani iz ledeburitnega orodnega jekla Vanadis 6 PM, so bili povr{insko obdelani, bru{eni in zrcalno polirani. Pred
nanosom so bili toplotno obdelani na trdoto 60 HRC. S tehniko magnetronskega napr{evanja in z uporabo tar~ iz ~istega Cr in
Ag je bila nanesena CrAg7N-nanoprevleka v sestavljeni nizkotla~ni atmosferi iz du{ika in argona pri temperaturi 500 °C v
napravi Hauzer Flexicoat 850. Tribolo{ki preizkusi s preizku{evalnikom “pin-on-disc” so bili izvr{eni pri sobni temperaturi in
pri povi{anih temperaturah (300, 400 in 500) °C. Krogle Al2O3, 100Cr6 in CuZn so bile uporabljene kot par. Sledi obrabe po
preizkusu "pin-on-disc" so bile analizirane z vrsti~no elektronsko mikroskopijo in z mikroanalizo. Eksperimenti so pokazali
mo~no odvisnost tribolo{kih parametrov od temperature. Koeficient trenja CrN-Ag proti krogli 100Cr6 pri sobni temperaturi je
bil μ = 0,56. Tribolo{ki drsni preizkusi te prevleke proti kroglam Al2O3 ka`ejo zmanj{anje koeficienta trenja z nara{~ajo~o
temperaturo. Pri sobni temperaturi je bil μ = 0,68, minimum pa se je pojavil pri temperaturi 400 °C, μ = 0,24. To se pripisuje
difuziji delcev Ag na drsno povr{ino pri povi{anih temperaturah. Preizku{anje v paru s CuZn in medeninasto kroglo je dalo na
splo{no ni`je koeficiente trenja pri sobni temperaturi, μ = 0,39. Nasprotno pa je koeficient trenja malo narasel pri povi{anju
temperature in je bil pri povi{anih temperaturah prakti~no konstanten v obmo~ju μ = 0,43–0,48.
Klju~ne besede: ledeburitno jeklo Cr-V Vanadis 6, PVD, krom nitrid s srebrom, "pin-on-disk", koeficient trenja
1 INTRODUCTION
Chromium nitrides (CrN) have been extensively in-
vestigated in the applications of protective coatings due
to their high hardness, good wear resistance as well as
excellent corrosion and high-temperature-oxidation re-
sistance.1–5 They gained great scientific interest and
industrial popularity due to these properties in copper
machining, aluminium die casting and forming, and
wood processing.6 However, in many applications, the
requirements on the coated-material surface cannot be
met by such a single coating. A further development to
adapt some of their properties to the levels required for
specific applications leads to the production of com-
posite coatings, combining different material properties
so that certain new desired properties can be created.7–9
The effect of self-lubrication has gained a great sci-
entific importance in the last few years. The main idea to
develop self-lubricating and multi-purpose coatings is
based upon the fact that commercially available lubri-
cants (sulfides, oxides, graphite) exhibit considerable
shortcomings and cannot be used effectively in tooling
applications over a sufficiently wide temperature
range.10–12 Soft noble metals, on the other hand, show a
stable chemical behavior and can exhibit self-lubricating
properties due to their low shear strength. Noble-metal
particles bring several benefits to the layer properties
compared to metal oxides or graphite. They are stable up
to relatively high temperatures, have a low hardness and
do not behave as abrasive particles. A common dis-
advantage of noble metals is their high cost, but this can
be optimized to an acceptable level. The self-lubricating
effect is based on an incorporation of a small amount of
noble metals, mostly silver, into the basic CrN film.
Silver is completely insoluble in CrN and forms nano-
particles in the basic CrN compound. Silver-containing
transition-metal-nitride films have been extensively
studied in recent years.13
Materiali in tehnologije / Materials and technology 48 (2014) 5, 669–673 669
UDK 539.92:621.793 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)669(2014)
The current paper deals with the development of
adaptive nanocomposite CrAgN coatings on the Vanadis
6 Cr-V ledeburitic tool steel. It describes and discusses
the tribological properties, such as the friction coefficient
and wear rate during the pin-on-disk testing, of the coat-
ing with the mass fraction w = 7 % content of silver.
2 EXPERIMENTAL WORK
The substrate material was PM ledeburitic steel
Vanadis 6 with nominally 2.1 % C, 1.0 % Si, 0.4 % Mn,
6.8 % Cr, 1.5 % Mo, 5.4 % V and Fe as the balance that
was soft annealed to a hardness of 21 HRC.14,15
The samples used for the investigation were plates
with the dimensions of 50 mm × 10 mm × 10 mm, heat
treated (austenitized at a temperature of 1050 °C,
quenched in a flow of nitrogen gas and double tempered
for 2 h at a temperature of 530 °C) to the final hardness
of 60 HRC and then finely ground and polished with a
diamond suspension up to a mirror finish.
The conditions for depositing CrN/Ag coatings were
reported elsewhere.16 The output power on the Cr cath-
ode was 5.8 kW and on the Ag cathode it was 0.21 kW.
Tribological properties of the coating were measured
using a CSM pin-on-disc tribometer at ambient and at el-
evated temperatures, up to 500 °C. Balls, 6 mm in
diameter, made from sintered alumina, 100Cr6 steel and
CuZn brass (55 % Cu, 45 % Zn) were used for the tests.
The testing against the 100Cr6 counterpart was carried
out only at ambient temperature due to a low thermal
stability of the 100Cr6 steel. No external lubricant was
added during the measurements. The normal loading
used for the investigation was 1 N. For each measure-
ment, the total sliding distance was 100 m. The volume
loss of the coated samples was calculated from the width
of the track using the following formula:17

 = R [r2 sin–1(d/2r) – (d/4)v(4r2 – d2)]
where R is the wear-track radius, d is the wear-track
width and r is the radius of the ball.
To relate the volume loss to normal load F and sliding
distance l wear rates W 17 were calculated.
After the testing, the wear tracks were examined with
a scanning electron microscope (SEM) JEOL
JSM-7600F and an energy-dispersive X-ray analysis
(EDX).
3 RESULTS AND DISCUSSIONS
Table 1 summarizes the results of the tribological in-
vestigations. In the case of the alumina counterpart, the
friction coefficient at ambient temperature was μ = 0.68.
Therefore, no positive effect of the Ag addition was
found at the low temperature, which is in line with the
previous investigation.18
Basnyat et al.19 and Yao et al.20 established a benefi-
cial effect of Ag on the friction coefficient at room tem-
perature. However, the loading applied in their investi-
gation was much higher, which makes the results
incomparable with our measured data.
At a higher testing temperature, the friction coeffi-
cient of the CrN coating with w = 7 % of silver became
much lower. The minimum value, μ = 0.24, was found at
the temperature of 400 °C. This phenomenon is attri-
buted to various factors. Firstly, there is an increasing
mobility of silver at elevated temperature that has been
reported.21 These atoms can diffuse to the surface at
elevated temperatures and effectively work as a solid
lubricant due to the low shear strength of silver. The
second possible contribution of the friction coefficient
decreased with the increased testing temperature can be
indentified on the basis of the assumption that an
increased temperature makes the coating softer.18
Figure 1 shows the dependence of the friction coeffi-
cient on the sliding distance. At ambient temperature, the
friction coefficient increased at the beginning of the test
and after that it slightly decreased to a stable state of μ =
0.68, typical for a steady state of sliding. At elevated
temperatures, the friction coefficient was first between μ
= 0.55–0.65 and then it immediately decreased to the
values of the steady state.
Figure 2 demonstrates the wear tracks obtained by
testing the CrAg7N coating against alumina at ambient
and elevated temperatures. At ambient temperature, the
testing gave a smooth surface without any failure of in-
tegrity. The testing at elevated temperatures led to a
creation of parallel grooves oriented along the sliding
direction on the coated material and to a much wider
wear track and a high volume loss of the coating and the
wear rate, respectively, as documented in Table 2. The
P. BÍLEK et al.: TRIBOLOGY OF CrAg7N COATINGS DEPOSITED ON VANADIS 6 LEDEBURITIC TOOL STEEL
670 Materiali in tehnologije / Materials and technology 48 (2014) 5, 669–673
Table 1: Results of tribological investigation of the CrAg7N coating
Tabela 1: Rezultati tribolo{kih preiskav nanosa CrAg7N
μ Counterpart
Temperature Al2O3 CuZn 100Cr6
20 °C 0.68 0.39 0.56
300 °C 0.40 0.43
400 °C 0.24 0.48
500 °C 0.29 0.46
Figure 1: Dependence of friction coefficient on sliding distance for
alumina counterpart
Slika 1: Odvisnost koeficienta trenja od poti drsenja za Al2O3 v paru
wider tracks can be attributed to the softening of the
coating at elevated temperatures.
At the temperature of 400 °C a partial failure of the
coating was observed. The EDX analysis showed the
presence of iron (the base element of steel) on the sur-
face of the track in some sites, while the contents of
chromium and silver were found in the other sites
(Figure 3). This result confirms a partial removal of the
coating from the substrate.
Our previous investigation of the CrN coating with w
= 3 % of silver gave similar results.22 The lowest value of
the friction coefficient was also found at the temperature
of 400 °C. However, the CrAg3N coating showed a
much higher failure, and even at the temperature of
500 °C it was completely removed from the surface of
the substrate after the test.
Figure 4 shows the surface of a track after the
pin-on-disc testing at the temperatures of 300 °C and 400
°C. In both cases, Ag particles are well visible and these
particles are responsible for a better friction and could
act as a solid lubricant. In the wear track formed during
the testing at 500 °C Ag particles were also identified,
but their population density was reduced in comparison
with the samples tested at lower temperatures.
Figure 5 shows the wear tracks formed by the sliding
of the coating against CuZn brass. Compared to the
tracks caused by the sliding of sintered alumina these
tracks are wider, but no damage of the coating is ob-
served.
On the other hand, a considerable amount of the ma-
terial transferred from the counterpart to the surface of
the coating was detected (Figure 6).
P. BÍLEK et al.: TRIBOLOGY OF CrAg7N COATINGS DEPOSITED ON VANADIS 6 LEDEBURITIC TOOL STEEL
Materiali in tehnologije / Materials and technology 48 (2014) 5, 669–673 671
Figure 4: Nanoparticles of silver on the surface of track after
pin-on-disc testing
Slika 4: Nanodelci srebra na povr{ini sledi po preizkusu "pin-on-disc"
Figure 2: Wear tracks after pin-on-disc testing against alumina at
ambient and elevated temperature
Slika 2: Sledi obrabe po preizkusu "pin-on-disc" v paru z Al2O3 pri
sobni in povi{ani temperaturi
Table 2: Width of tracks, volume loss and wear rate after pin-on-disc
testing against alumina
Tabela 2: [irina sledi, volumenska izguba in obraba pri preizkusu
"pin-on-disc" v paru z Al2O3
Heat (°C) d/μm V/m3 W/(m3/(N m))
20 85 2.72 · 10–13 2.72 · 10–15
300 228 5.17 · 10–12 5.17 · 10–14
400 250 6.85 · 10–12 6.85 · 10–14
500 276 9.24 · 10–12 9.24 · 10–14
Figure 5: Wear tracks after pin-on-disc testing against CuZn at am-
bient and elevated temperature
Slika 5: Sledi obrabe po preizkusu "pin-on-disc" v paru s CuZn pri
sobni in povi{ani temperaturi
Figure 3: Partial failure of coating CrAg7N after pin-on-disc testing
against alumina at the temperature of 400 °C: a) overview, b) EDX of
chromium, c) EDX of silver, d) EDX of iron
Slika 3: Parcialne po{kodbe nanosa CrAg7N po preizkusu "pin-on-
disc" v paru z Al2O3 pri temperaturi 400 °C: a) videz, b) EDX kroma,
c) EDX srebra, d) EDX `eleza
This is due to the very low shear strength of CuZn,
especially at a higher temperature. It should be noted that
such a material transfer was detected irrespective of the
testing temperature. These results are in good agreement
with the previous work.22
The measurement results for the friction coefficient
of the CuZn counterpart are shown in Table 1. The mini-
mum friction coefficient of μ = 0.39 was found at
ambient temperature. With the increasing temperature,
the friction coefficient becomes higher, ranging between
μ = 0.43–0.48. No effect of silver on the tribological
performance was found here – the material transfer can
be considered as a plausible explanation (Figure 6).
The friction coefficient against the 100Cr6 counter-
part was found to be μ = 0.56 (Table 1). The material of
steel transferred to the surface of the coating is docu-
mented in Figure 7. The surface of the coating stayed
smooth, without any damage.
4 CONCLUSIONS
The friction and wear characteristics of the CrAg7N
coatings prepared with the magnetron-sputter-deposition
method were examined at different temperatures and
with different counterpart materials. The results can be
summarized as follows:
• The friction coefficient rapidly decreased with the
increasing testing temperature when an alumina ball
was used as the counterpart, with its minimum at the
temperature of 400 °C. The coating showed partial
damage at elevated temperatures, but it was not
removed from the substrate.
• The friction coefficient was not positively influenced
when tested against the CuZn counterpart.
• No removal or other coating damage was established
after the sliding against the CuZn ball and 100Cr6
ball, but a considerable material transfer from the ball
to the sample was detected.
• The presence of silver in the CrN coating results in
improved tribological properties at moderate tem-
peratures. Under these conditions, the CrN coating
with w = 7 % of Ag could find its application in
industry.
Acknowledgements
This publication is the result of implementing project
CE for development and application of advanced diag-
nostic methods in processing of metallic and non-metal-
lic materials, ITMS:26220120048, supported by the Re-
search & Development Operational Programme funded
by the ERDF. The research was supported by grant pro-
ject VEGA 1/1035/12.
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P. BÍLEK et al.: TRIBOLOGY OF CrAg7N COATINGS DEPOSITED ON VANADIS 6 LEDEBURITIC TOOL STEEL
672 Materiali in tehnologije / Materials and technology 48 (2014) 5, 669–673
Figure 7: Material transferred from counterpart 100Cr6 to the surface
of coating CrAg7N after pin-on-disc, at the temperature 20 °C: a)
overview, b) EDX of chromium, c) EDX of silver, d) EDX of iron
Slika 7: Prenesen material iz para 100Cr6 na povr{ino nanosa
CrAg7N po preizkusu "pin-on-disc" pri temperaturi 20 °C: a) videz, b)
EDX kroma, c) EDX srebra, d) EDX `eleza
Figure 6: Material transferred from counterpart CuZn to the surface
of coating CrAg7N after pin-on-disc, at the temperature of 20 °C: a)
overview, b) EDX of chromium, c) EDX of silver, d) EDX of copper,
e) EDX of zinc
Slika 6: Prenesen material s para CuZn na povr{ino nanosa CrAg7N
po preizkusu "pin-on-disc" pri temperaturi 20 °C; a) videz, b) EDX
kroma, c) EDX srebra, d) EDX bakra, e) EDX cinka
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com/files/van_6_ds.pdf>
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ceedings of the 22nd Int. Conference METAL 2013, Brno, 2013
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P. BÍLEK et al.: TRIBOLOGY OF CrAg7N COATINGS DEPOSITED ON VANADIS 6 LEDEBURITIC TOOL STEEL
Materiali in tehnologije / Materials and technology 48 (2014) 5, 669–673 673

M. KAMIERCZAK et al.: MORPHOLOGY AND MAGNETIC PROPERTIES OF Fe3O4-ALGINIC ACID NANOCOMPOSITES
MORPHOLOGY AND MAGNETIC PROPERTIES OF
Fe3O4-ALGINIC ACID NANOCOMPOSITES
MORFOLOGIJA IN MAGNETNE LASTNOSTI NANOKOMPOZITOV
Fe3O4-ALGINSKA KISLINA
Ma³gorzata KaŸmierczak1,2, Katarzyna Pogorzelec-Glaser1, Andrzej Hilczer1,
Stefan Jurga2, £ukasz Majchrzycki3, Marek Nowicki3, Ryszard Czajka3,
Filip Matelski3, Rados³aw Pankiewicz4, Bogus³awa £êska4, Leszek Kêpiñski5,
Bart³omiej Andrzejewski1
1Institute of Molecular Physics, Polish Academy of Sciences, M. Smoluchowskiego 17, 60179 Poznan, Poland
2NanoBioMedical Centre, Adam Mickiewicz University, Umultowska 85, 61614 Poznan, Poland
3Poznan University of Technology, Nieszawska 13A, 60965 Poznan, Poland
4Faculty of Chemistry, Adam Mickiewicz University, Umultowska 89b, 61614 Poznan, Poland
5Institute of Low Temperature and Structure Research, Polish Academy of Sciences, Okolna 2, 50422 Wroclaw, Poland
malgorzata.kazmierczak@ifmpan.poznan.pl
Prejem rokopisa – received: 2013-09-27; sprejem za objavo – accepted for publication: 2013-11-19
The morphology, structure and magnetic properties of the nanocomposites of magnetite (Fe3O4) nanoparticles and alginic acid
(AA) are studied. Magnetite Fe3O4 nanoparticles and the nanoparticles capped with alginic acid exhibit very distinct properties.
The chemical bonding between alginic acid and the surface of magnetite nanoparticles results in the recovery of surface
magnetization. On the other hand, it also leads to the enhanced surface spin disorder and unconventional behavior of the
magnetization observed in Fe3O4-AA nanocomposites at low temperatures.
Keywords: nanocomposite, magnetite nanoparticles, alginic acid, enhanced magnetization
Preu~evali smo morfologijo, strukturo in magnetne lastnosti nanokompozitov na osnovi nanodelcev magnetita (Fe3O4) in
alginske kisline (AA). V primerjavi z magnetitnimi nanodelci izkazujejo nanokompoziti Fe3O4-alginska kislina precej druga~ne
lastnosti. Molekule alginske kisline se kemijsko ve`ejo na povr{ino magnetitnih nanodelcev in s tem povzro~ijo vrnitev
povr{inske magnetizacije. Hkrati pa se s tem v nanokompozitih Fe3O4-AA pri ni`jih temperaturah pove~a povr{inska
neurejenost spinov in nekonvencionalno vedenje magnetizacije.
Klju~ne besede: nanokompozit, nanodelci magnetita, alginska kislina, pove~ana magnetizacija
1 INTRODUCTION
The interest in the composites of polymers with
magnetic nanoparticles stems from their unique physical
properties and potential future applications for mag-
netic-data storage,1 electronic devices and sensors,2
biomedical applications in magnetic resonance imaging,3
drug delivery4 and hyperthermia agents.5 From this point
of view, one of the most preferred magnetic materials is
magnetite Fe3O4 because it is a biocompatible mineral
with a low toxicity (for example, the crystals of magne-
tite are magnetoreceptors in the brains of some ani-
mals6). It also exhibits a large magnetic moment and a
spin-polarized electric current – the features highly
desired for the applications in spintronics.
In bulk, magnetite crystallizes in the inverse spinel
AB2O4 structure with two nonequivalent Fe sites placed
in the fcc lattice of O2– ions. Tetrahedral A sites contain
Fe2+ ions, whereas octahedral B sites are occupied by
Fe2+ and Fe3+ ions. The magnetic sublattices located on A
and B sites are ferrimagnetically coupled. The mixed
valence of Fe ions and fast electron hopping between B
sites are responsible for a relatively high electric
conductivity of Fe3O4 above the Verwey transition, Tv ≈
125 K.7
Nanostructured magnetite exhibits different mag-
netic, electronic and optical properties than the bulk
material. Particularly, a significant reduction in the
magnetization at the surface of Fe3O4 nanoparticles
makes them useless for many applications. This obstacle
can be overcome by capping the magnetic nanoparticles
with polymers8 or organic acids, which allows a
restoration of the surface magnetism.9
One of the best capping material is alginic acid,
which is a cheap, common and nontoxic natural biopoly-
mer.10,11 The aim of this work is to study the effect of the
alginic-acid capping on the surface magnetization reco-
very in Fe3O4 nanoparticles.
2 EXPERIMENTAL WORK
2.1 Sample synthesis
All the chemicals used in the experiments were
purchased from SIGMA ALDRICH. To obtain the dis-
aggregated nanoparticles of magnetite Fe3O4, a portion
of 9.0 mmol of FeCl3 · 6H2O was dissolved in 200 mL of
ethylene glycol. The solution was vigorously stirred.
After 15 min 131.7 mmol of CH3COONa and
1.575 mmol of polyethylene glycol PEG 400 were added
Materiali in tehnologije / Materials and technology 48 (2014) 5, 675–678 675
UDK 620.3:66.017:621.318 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)675(2014)
and the stirring was continued until they completely
dissolved. Then, the solution was transferred into 50 mL
teflon reactors and heated using microwave radiation
(MARS 5, CEM Corporation) at 160 °C for 25 min. The
black suspension of the nanoparticles obtained as a result
of the reaction was first cooled, isolated by centrifuga-
tion and washed with absolute ethanol. The final product
was dried in a vacuum oven at 40 °C. A nanocomposite
was prepared from the aqueous dispersion of the
magnetite nanopowder and alginic acid (AA) that was
then air-dried at room temperature. The nanocomposite
had the form of flakes with flat surfaces.
2.2 Sample characterization
The crystallographic structures of the samples were
studied by means of X-ray powder diffraction (XRD)
using an ISO DEBYE FLEX 3000 instrument with a Co
lamp ( = 0.17928 nm). The morphology of Fe3O4
nanoparticles was observed using a Philips CM20
SuperTwin transmission electron microscope (TEM).
The structures of nanocomposites were studied by means
of an atomic force microscope (Dimension Icon®, Bru-
ker) using the magnetic-force-microscope (MFM) mode
and NANO-SENSORS™ PPP-MFMR probes. The mag-
netic measurements were performed using a Quantum
Design physical property measurement system (PPMS)
fitted with a vibrating-sample-magnetometer (VSM)
probe.
3 RESULTS AND DISCUSSION
Figure 1 shows the X-ray powder diffraction patterns
of the as-obtained Fe3O4 nanoparticles (panel a) and of
the Fe3O4-AA nanocomposite with the magnetite content
equal to the mass fraction w = 10 % (panel b).
The solid line corresponds to the best Rietveld profile
fit calculated by means of the FULLPROF software for
the cubic crystal structure with the Fd-3m space group
and X-ray radiation with the wavelength of 0.17928 nm,
as used in the experiment. The vertical bars correspond
to the Bragg peaks and the line below them is the
difference between the experimental data and the fit.
XRD studies verified the Fd-3m point group of the Fe3O4
nanopowder with the lattice parameters of a = 0.83641 nm
and the mean crystallite size of 20 nm determined with
the Scherrer method. For the composite, the intensity of
diffraction peaks is too low to perform an analysis, even
if the content of magnetite is high and equal to w = 10 %.
A TEM image of magnetite nanoparticles is pre-
sented in the inset to Figure 2. The magnetic nanoparti-
cles are almost monodisperse and spherical. A histogram
of the particle-size distribution of Fe3O4 nanoparticles is
presented in Figure 2. The distribution can be fitted with
a log-normal function:
f x
x
x
x
( ) exp ln= − ⎛⎝
⎜ ⎞
⎠
⎟⎡
⎣⎢
⎤
⎦⎥
1
2
1
22 2
2
π 
(1)
where x is the mean size of the nanoparticles and  is
the distribution width. The values characterizing the
distribution are: x = 20.5 nm and  = 0.11, with x
corresponding well to the XRD data.
Figure 3 shows the topography (panel a), elastic
properties (panel b) and magnetic domains (panels c and
d) of the Fe3O4-AA composite surface with the Fe3O4
content of 10 %, studied with MFM. The roughness of
the surface is below 10 nm for the scanned area of 500
nm × 500 nm. The knobs on the topography image (the
white spots) indicate the presence of small agglomerates
of Fe3O4 nanoparticles that are also seen as white areas
on the phase-contrast image (panel b). The amplitude
and phase contrast of the magnetic signal (panels c and
d) indicate the presence of magnetic domains with the
size close to 100 nm. The actual size of these domains
M. KAMIERCZAK et al.: MORPHOLOGY AND MAGNETIC PROPERTIES OF Fe3O4-ALGINIC ACID NANOCOMPOSITES
676 Materiali in tehnologije / Materials and technology 48 (2014) 5, 675–678
Figure 2: Histogram for Fe3O4 nanoparticles with a log-normal
fitting. A TEM image is shown in the inset.
Slika 2: Histogram nanodelcev Fe3O4. TEM-posnetek je prikazan v
vstavku.
Figure 1: XRD powder pattern and line profile fitting of: a) Fe3O4
nanoparticles and b) Fe3O4-AA nanocomposite
Slika 1: XRD-difraktogrami in ujemanje linij za: a) nanodelce Fe3O4
in b) nanokompozit Fe3O4-AA
can be smaller than that presented in the figures because
of the insufficient spatial resolution of the MFM method
(about 50 nm) which causes a smearing of the images.
The results of the magnetic study are presented in
Figures 4 and 5. The magnetization is normalized with
respect to the content of magnetite in the samples. For
the nanoparticles of Fe3O4, the temperature dependence
of magnetization M  T1.9 deviates from the Bloch law
M  T1.5 valid for the capped nanoparticles of magnetite
(Figure 4). The deviation from the Bloch law for the
uncapped Fe3O4 nanoparticles can be related to a
degraded magnetic ordering at the surface. Moreover, the
magnetization of the Fe3O4 nanoparticles at room
temperature is only 51 A m2/kg, i.e., much below the
saturation value for the bulk magnetite ( 90 A m2/kg),
and also lower than the magnetization of the capped
particles, equal to 60 A m2/kg. The enhancement of the
magnetization and the Bloch-like behavior of the capped
nanoparticles can be explained in terms of the recovery
of surface magnetism due to the chemical bonding
between the AA and Fe3O4 nanoparticles. This bonding
between the O atoms in the carboxylic groups and two of
the four Fe atoms in the Fe-O surface unit cell makes the
coordinations and distances close to those in the bulk.9
The remaining two Fe atoms still exhibit a reduced
magnetization because they are closer to the in-plane
oxygens, which results in partially empty dx2 – y2 orbitals.
The inhomogeneity with respect to the Fe coordination
can be responsible for the increased spin disorder or
unconventional magnetism at the Fe3O4 surface. This un-
conventional behavior is manifested as a rapid increase
in the magnetization at a low temperature observed for
Fe3O4-AA composites (Figure 4). The alternative expla-
nation of this magnetization upturn assumes a
quantization of the spin-wave spectrum due to the finite
size of the particles that occurs at low temperatures and
is responsible for the deviation from the Bloch law.12
The magnetization loops M(H) for the Fe3O4 nano-
particles and Fe3O4-AA composites are shown in Figure
5.
Both the nanoparticles and composites exhibit ferro-
magnetic (ferrimagnetic) hysteresis loops, which saturate
above about 0.3 T. The magnetization of the composites
is enhanced as compared to that of the uncapped Fe3O4
nanoparticles. At low temperatures the magnetization
loops for the composites are the superpositions of the
ferromagnetic and linear contribution from an unconven-
tional magnetism. This unconventional behavior cannot
be simply related to the paramagnetism at the degraded
M. KAMIERCZAK et al.: MORPHOLOGY AND MAGNETIC PROPERTIES OF Fe3O4-ALGINIC ACID NANOCOMPOSITES
Materiali in tehnologije / Materials and technology 48 (2014) 5, 675–678 677
Figure 5: Magnetization loops M(H) for the Fe3O4 nanoparticles and
Fe3O4-AA composites with mass fractions 5 % and 10 % of the
magnetite content
Slika 5: Histerezna zanka M(H) nanodelcev Fe3O4 in kompozitov
Fe3O4-AA z masnim dele`em magnetita 5 % in 10 %
Figure 3: a) Surface topography, b) phase contrast, c) magnetic phase
and d) magnetic amplitude images for the Fe3O4-AA composite with
the Fe3O4 mass fraction of 10 %
Slika 3: a) Povr{inska topografija, b) fazni kontrast, c) magnetna faza
in d) magnetna amplituda kompozitov Fe3O4-AA z masnim dele`em
Fe3O4 10 %
Figure 4: Magnetization M(T) of the Fe3O4 nanoparticles and
Fe3O4-AA composites containing mass fractions 5 % and 10 % of
magnetite
Slika 4: Magnetizacija M(T) nanodelcev Fe3O4 in kompozitov
Fe3O4-AA z masnim dele`em magnetita 5 % in 10 %
Fe3O4 surface because it is absent in the uncapped nano-
particles of magnetite.
4 CONCLUSIONS
The capping of Fe3O4 nanoparticles with alginic acid
leads to a partial recovery of the surface magnetization.
On the other hand, the bonding between alginic acid and
Fe3O4 nanoparticles by means of O atoms results in an
unconventional magnetism observed at low temperatures.
Acknowledgments
This project was supported by the National Science
Centre through project No. N N507 229040. M.K. was
supported through the European Union – European
Social Fund and Human Capital – National Cohesion
Strategy. We thank dr. Alja Kupec and dr. Brigita Ro`i~
for the Slovenian translation.
5 REFERENCES
1 D. Weller, A. Moser, IEEE Trans. Magn., 35 (1999), 4423
2 I. Koh, L. Josephson, Sensors, 9 (2009), 8130
3 C. W. Jung, P. Jacobs, Mag. Reson. Imaging, 13 (1995), 661
4 S. Bucak, B. Yavuztürk, A. D. Sezer, Magnetic Nanoparticles:
Synthesis, Surface Modifications and Application in Drug Delivery,
In: A. D. Sezer (ed.), Recent Advances in Novel Drug Carrier
Systems, InTech Publisher, 2012
5 P. Wunderbaldinger, L. Josephson, R. Weissleder, Bioconjugate
Chem., 13 (2002), 264
6 R. R. Baker, J. G. Mather, J. H. Kennaugh, Nature, 301 (1983), 79
7 F. Walz, J. Phys.: Condens. Matter, 14 (2002), R285
8 E. Sówka, M. Leonowicz, B. Andrzejewski, A. D. Pomogailo, G. I.
Dzhardimalieva, J. Alloys Comp., 423 (2006), 123
9 J. Salafranca, J. Gazquez, N. Pérez, A. Labarta, S. T. Pantelides, S. J.
Pennycook, X. Batlle, M. Varela, Nano Letters, 12 (2012), 2499
10 Z. Durmus, H. Sözeri, B. Unal, A. Baykal, R. Topkaya, S. Kazan, M.
S. Toprak, Polyhedron, 30 (2011), 322
11 B. Unal, M. S. Toprak, Z. Durmus, H. Sözeri, A. Baykal, J. Nano-
part. Res., 12 (2010), 3039
12 K. Mandal, S. Mitra, P. A. Kumar, Europhys. Lett., 75 (2006), 618
M. KAMIERCZAK et al.: MORPHOLOGY AND MAGNETIC PROPERTIES OF Fe3O4-ALGINIC ACID NANOCOMPOSITES
678 Materiali in tehnologije / Materials and technology 48 (2014) 5, 675–678
A. GRAJCAR, K. RADWAÑSKI: MICROSTRUCTURAL COMPARISON OF THE THERMOMECHANICALLY ...
MICROSTRUCTURAL COMPARISON OF THE
THERMOMECHANICALLY TREATED AND COLD
DEFORMED Nb-MICROALLOYED TRIP STEEL
PRIMERJAVA MIKROSTRUKTUR TERMOMEHANSKO
OBDELANEGA IN HLADNO DEFORMIRANEGA, Z
Nb-MIKROLEGIRANEGA TRIP-JEKLA
Adam Grajcar1, Krzysztof Radwañski2
1Silesian University of Technology, Institute of Engineering Materials and Biomaterials, 18a Konarskiego Street, 44-100 Gliwice, Poland
2Institute for Ferrous Metallurgy, 12-14 K. Miarki Street, 44-100 Gliwice, Poland
adam.grajcar@polsl.pl
Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2013-11-04
The work deals with a microstructural comparison of the thermomechanically processed and subsequently cold deformed
Nb-microalloyed Si-Al-type multiphase steel, showing a TRIP effect. The newly developed steel was subjected to the
thermomechanical rolling and controlled cooling under the conditions allowing us to obtain a fine-grained ferritic-bainitic
microstructure with a large fraction of the retained austenite. Subsequently, the thermomechanically rolled sheet samples were
subjected to a 10 % elongation in uniaxial tension. The comparison of the multiphase microstructures and, especially, the
identification of the strain-induced martensite were carried out using light microscopy, electron transmission microscopy and
electron scanning microscopy equipped with EBSD (electron backscatter diffraction). Morphological details influencing the
mechanical stability of the retained austenite were indicated.
Keywords: thermomechanical treatment, TRIP effect, multiphase steel, retained austenite, strain-induced martensite, EBSD
technique
Delo obravnava primerjavo mikrostruktur termomehansko obdelanega in nato hladno deformiranega, z Nb-mikrolegiranega
ve~faznega jekla Al-Si, ki izkazuje vedenje TRIP. Novo razvito jeklo je bilo termomehansko valjano in kontrolirano ohlajeno v
razmerah, ki omogo~ajo doseganje drobnozrnate feritno-bainitne mikrostrukture z velikim dele`em zaostalega avstenita.
Termomehansko izvaljani vzorci plo~evine so bili nato izpostavljeni 10-odstotnemu raztezku pri enoosni natezni obremenitvi.
Primerjava multifaznih mikrostruktur in posebno dolo~anje napetostno induciranega martenzita je bila izvr{ena s svetlobno
mikroskopijo, elektronsko presevno mikroskopijo in elektronsko vrsti~no mikroskopijo, opremljeno z EBSD (Electron
Backscatter Diffraction). Prikazane so morfolo{ke podrobnosti, ki vplivajo na mehansko stabilnost zaostalega avstenita.
Klju~ne besede: termomehanska obdelava, u~inek TRIP, ve~fazno jeklo, zaostali avstenit, napetostno inducirani martenzit,
EBSD-tehnika
1 INTRODUCTION
The mechanical properties and technological for-
mability of advanced high-strength steels (AHSS) for the
automotive industry depend on the relative proportions
and mechanical properties of individual microstructural
constituents. Ferrite forms a matrix of AHSS whereas
the strengthening phases consist of martensite and/or
bainite. A very attractive combination of high strength
and ductility can be obtained in dual-phase (DP) steels
consisting of a ferrite matrix and uniformly distributed
martensitic or martensitic-bainitic islands.1–4 A further
growth of the strength-ductility balance can be obtained
for the steels with a ferritic matrix containing bainitic-
austenitic islands, where the final mechanical properties
are formed during cold working under the conditions of
the strain-induced martensitic transformation of the
metastable retained austenite.5–8 Multiphase steel sheets
are produced with the continuous annealing of cold-
rolled sheets3,7 or they are thermomechanically hot rolled
and controlled cooled.8,9 A further increase in the
strength properties of multiphase steels requires modi-
fied chemical-composition concepts. Recently, Nb, Ti
and V microalloying has been used to enhance the
strength of multiphase steels8–10, well known for its
beneficial effect in HSLA steels.11–15
Thermomechanically processed multiphase steels are
characterized by a high-grain refinement. Therefore, the
qualitative and quantitative identifications of individual
structural constituents are especially important. A deter-
mination of the -phase volume fraction is achieved
using a computer-image analysis after the colour etching,
X-ray or neutron diffraction and magnetic methods.5,6,8,16
Recently, the EBSD technique of a scanning electron
microscope (SEM) has had an essential significance in
determining the fractions and morphological features of
various microstructural constituents.17–19 Microstructural
details of cold-rolled multiphase steels have been charac-
terized to a sufficient extent.7,17–19 However, there are
very few reports concerning morphological details of
thermomechanically rolled TRIP steels.9 Hence, the
present study addresses the microstructure evolution of a
hot-rolled Nb-microalloyed Si-Al multiphase steel.
Materiali in tehnologije / Materials and technology 48 (2014) 5, 679–683 679
UDK 620.186:621.78 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)679(2014)
2 EXPERIMENTAL PROCEDURE
The chemical composition of the newly developed
steel was designed with the focus on maximizing the
retained-austenite fraction and obtaining the carbide-free
bainite. The steel contains: 0.24 % C, 1.55 % Mn,
0.87 % Si, 0.40 % Al, 0.034 % Nb, 0.023 % Ti, 0.004 %
S and 0.010 % P. Nb and Ti were used to increase the
strength and to achieve the grain refinement during hot
rolling. The ingot was produced with vacuum induction
melting and then it was hot forged to a thickness of
22 mm. Subsequently, the flat samples were roughly
rolled to a thickness of 4.5 mm within the temperature
range between 1200 °C and 900 °C. The thermomecha-
nical rolling was conducted in 3 passes between 1100 °C
and 850 °C to the final sheet thickness of about 2 mm.
After the final deformation at 850 °C the specimens were
air cooled to 700 °C and then slowly to the temperature
of 650 °C for 50 s using furnace cooling. The next step
included immerse cooling of the sheets at a rate of about
50 °C/s, using a water-polymer medium, to the isother-
mal holding temperature (450 °C) at the bainitic trans-
formation range. The specimens were held at 450 °C for
600 s and finally cooled at a rate of about 0.5 °C/s to
room temperature. Then, standard-sized, A50 tensile-test
samples with a gauge length of 50 mm and a width of
12.5 mm were cut parallel to the rolling direction of the
sheets and deformed to 10 % of the plastic strain at a
strain rate of 5 × 10–3 s–1. Metallographic specimens were
taken at different points along the rolling direction for
both thermomechanically processed samples and those
subjected to cold deformation.
For the purpose of a detailed analysis of all the
microstructural constituents and, especially, to identify
the strain-induced martensitic transformation, light
microscopy (LM), scanning electron microscopy (SEM)
and transmission electron microscopy (TEM) were used.
Additionally, orientation imaging microscopy (OIM)
using SEM was applied. Etching in a 10 % aqueous
solution of sodium metabisulfite was used. Metallo-
graphic observations at the magnification of 1000-times
were carried out with a Leica MEF 4A light microscope.
Morphological details of microstructural constituents of
the steel were revealed with SUPRA 25 SEM using
back-scattered electrons (BSE). Observations were
performed on nital-etched samples at the accelerating
voltage of 20 kV. The EBSD technique was performed
using Inspect F SEM equipped with Shottky field
emission. After the classical grinding and polishing, the
specimens were polished with Al2O3 with a granularity
of 0.1 μm. The final stage of the sample preparation was
the ion polishing using the GATAN 682 PECS system. A
fraction of the retained austenite in both the initial state
and after cold deformation, assessed with EBSD (the
average value of five measurements) was determined at a
magnification lower than 3000-times to obtain reliable
quantitative results.
The thin-foil investigations were carried out using a
JEOL JEM 3010 at the accelerating voltage of 200 kV.
Mechanically grinded disk specimens were polished at
the voltage of 17 V and current density of 0.2 A/cm2.
The mixture of 490 mL H3PO4 + 7 mL H2SO4 + 50 g
CrO3 was used as the electrolyte.
3 RESULTS AND DISCUSSION
Applying the thermomechanical rolling and con-
trolled cooling results in a fine-grained ferritic matrix
with a volume fraction of about 60 % containing uni-
formly distributed bainitic-austenitic and austenitic
islands (Figure 1a). The amount of the retained austenite
determined by EBSD is about 13.8 %. The carbon
content (C) of the  phase determined earlier8 using the
A. GRAJCAR, K. RADWAÑSKI: MICROSTRUCTURAL COMPARISON OF THE THERMOMECHANICALLY ...
680 Materiali in tehnologije / Materials and technology 48 (2014) 5, 679–683
Figure 1: a), b) Fine-grained ferritic matrix containing bainitic-auste-
nitic and austenitic islands after the thermomechanical processing and
c) the finest regions of the retained austenite;  – ferrite, B-A – baini-
tic-austenitic islands, R – retained austenite
Slika 1: a), b) Drobnozrnata feritna osnova vsebuje po termomehanski
obdelavi bainitno-avstenitne in avstenitne oto~ke in c) najdrobnej{a
podro~ja zaostalega avstenita;  – ferit, B-A – bainitnoavstenitni
oto~ki, R – zaostali avstenit
X-ray analysis equals mass fraction 1.14 %, which
corresponds to lowering the martensite start temperature
of the  phase to about 10 °C. The mean ferrite grain size
is equal to 6 μm and the retained austenite is located
along the ferrite boundaries as blocky grains with the
size up to 5 μm. Some retained austenite forms a halo
around the -phase grains whereas another part is
located at the ferrite-bainite interfaces. A large fraction
of the retained austenite, formed as thin layers or small
blocky grains with the size of between 1 ìm and 3 μm, is
a constituent of the bainitic islands (B-A) (Figure 1b). A
utilization of TEM reveals the finest regions of the
retained austenite with the sizes from 50 nm to 200 nm
(Figure 1c).
Figure 2 is an EBSD map using colour coding for
determining individual grains. The inverse-pole figure
(Figure 2a) shows the crystal direction parallel to the
normal direction of the specimen using the colour coding
according to the unit triangle. The grains of the highest
size with a random crystallographic orientation can be
recognized as ferrite. In the grey-scale image-quality
(IQ) map (Figure 2c) they correspond to the brightest
regions of the best diffraction quality. The IQ factor
represents a quantitative description of the sharpness of
the EBSD pattern. A lattice distorted by crystalline
defects such as dislocations and subgrain boundaries has
a distorted Kikuchi pattern, leading to lower IQ values.7,8
The retained austenite, bainite and grain boundaries are
represented by different levels of dark grey because their
pattern contrast is lower than that of the ferrite. The
color-coded phase map in Figure 2b clearly shows the
distribution of the retained austenite. The BCC
constituents and the retained austenite are distinguished
on the basis of the differences in their crystal structures,
whereas for a discrimination of the bainite from the
ferrite the differences in the IQ values have to be used.
Due to a very small size of the retained austenite, its
phase map is highly fragmented. The grain area of the
retained-austenite particles covers a range of up to 14
μm2 but the majority of the grains is smaller than 6 μm2
(Figure 3). Having the knowledge of the retained-auste-
nite presence and analyzing the IQ map in Figure 2c as
well as the misorientation-angle distribution of the
retained austenite (Figure 4), it is possible to indicate the
bainite regions. Firstly, this phase always occurs in
conjunction with the retained austenite and, secondly,
having high-lattice imperfections, it corresponds to the
dark regions of the low IQ. High-angle boundaries
(> 15°) occur between the ferrite grains, BCC consti-
tuents and the retained austenite as well as between the
ferrite and the bainite (Figure 2c). A large fraction of the
retained austenite exhibits a misorientation angle close to
45° with the neighboring BCC constituents (Figure 4).
This is in line with the earlier results obtained by Zaeffe-
rer et al.19 for the 0.2C-1.4Mn-0.5Si-0.7Al steel and
Petrov et al.7, Wasilkowska et al.18 for the 0.2C-1.5Mn-
1.5Si steel, according to which bainite regions require a
Kurdjumov-Sachs (K-S) or Nishiyama-Wasserman
A. GRAJCAR, K. RADWAÑSKI: MICROSTRUCTURAL COMPARISON OF THE THERMOMECHANICALLY ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 679–683 681
Figure 3: Distribution of the retained-austenite grain area
Slika 3: Razporeditev podro~ij zrn zaostalega avstenita
Figure 2: Microstructure images of the thermomechanically pro-
cessed steel obtained with EBSD: a) inverse pole-figure map showing
the grains with different crystallographic orientations, b) marked
regions of the retained austenite, c) image-quality map with crystallo-
graphic misorientation angles
Slika 2: EBSD-posnetki mikrostrukture termomehansko predelanega
jekla: a) zemljevid inverznih polovih figur, ki prikazuje zrna z razli~no
kristalografsko orientacijo, b) ozna~ena podro~ja zaostalega avstenita,
c) zemljevid kvalitete zrn s kristalografsko razli~nimi koti
(N-W) orientation for their growth, confirming a displa-
cive growth mechanism of bainite.
The -phase content decreases to about 7.7 % after
applying a 10 % tensile strain. It corresponds to the 44 %
initial austenite volume fraction transformed into mar-
tensite due to a strain-induced transformation. Generally,
the strain-induced martensitic transformation initially
proceeds in the largest and medium-sized austenite
grains located in a ferritic matrix (Figure 5a). Martensite
usually forms in the central zones of the grains whereas
the borders remain untransformed (Figure 5b). This
confirms a higher enrichment in carbon of the regions
adjacent to the ferrite grains and a smaller enrichment of
the central austenite regions as a result of a longer
A. GRAJCAR, K. RADWAÑSKI: MICROSTRUCTURAL COMPARISON OF THE THERMOMECHANICALLY ...
682 Materiali in tehnologije / Materials and technology 48 (2014) 5, 679–683
Figure 5: a), b) Ferritic-bainitic microstructures containing the
retained austenite and strain-induced martensite of the steel strained to
the elongation of 10 %, c) plate morphology of the strain-induced
martensite;  – ferrite, B-A – bainitic-austenitic islands, R – retained
austenite, M – martensite
Slika 5: a), b) Feritno-bainitna mikrostruktura z zaostalim avstenitom
in napetostno induciranim martenzitom v jeklu z 10-odstotno natezno
deformacijo, c) plo{~ata morfologija napetostno induciranega mar-
tenzita;  – ferit, B-A – bainitno-avstenitni oto~ki, R – zaostali
avstenit, M – martenzit
Figure 6: Microstructure images of the steel strained to the elongation
of 10 %: a) inverse pole-figure map showing grains with different
crystallographic orientations, b) marked regions of the retained
austenite, c) image-quality map with the crystallographic misorienta-
tion angles corresponding to the K-S and N-W relationships;  –
ferrite, B – bainitic ferrite, B-M-A – bainitic-martensitic-austenitic
regions, SZ – retained austenite
Slika 6: Posnetki mikrostrukture jekla, natezno obremenjenega do
raztezka 10 %: a) zemljevid inverznih polovih figur, ki prikazujejo
zrna z razli~no kristalografsko orientacijo, b) ozna~ena podro~ja zao-
stalega avstenita, c) zemljevid kvalitete z razli~nimi kristalografskimi
koti, ustrezno razmerjem K-S in N-W;  – ferit, B – bainitni ferit,
B-M-A – bainitno-martenzitno-avstenitno podro~je, SZ – zaostali
avstenit
Figure 4: Distribution of the crystallographic misorientation angles of
the grains
Slika 4: Razporeditev zrn z razli~nimi kristalografskimi koti
diffusion path of carbon. The formed martensite has a
plate morphology (Figure 5c) and it contributes to a
fragmentation of the untransformed austenite fostering
its further stabilization due to a reduction in the particle
size. It is clear from Figure 6 that the strain-induced
martensitic transformation is also initiated inside the
largest bainitic-austenitic islands resulting in a further
fragmentation of -phase particles. Moreover, the
detailed analysis from Figure 6c indicates that the K-S
and N-W relationships are partially kept between the
austenite and the bainitic ferrite. However, the fraction
fulfilling the special crystallographic orientations decre-
ases (Figure 7) compared to the initial state (Figure 4).
The fragmentation of the retained austenite is revealed
through a reduction of the grain-size area below 4 μm2
(Figure 8). It should be noted that the fraction of the
grain area changes roughly in proportion to the inverse
of the particle size.
4 CONCLUSIONS
A detailed identification of the morphological
features of individual microstructural constituents of
thermomechanically processed multiphase steels is a
challenging problem due to a high dispersion of particu-
lar phases. The problem becomes even more complicated
during cold deformation, when the highly dispersed
retained austenite transforms into the strain-induced
martensite. It was shown that the investigated Si-Al
TRIP steel is characterized by a fine-grained ferritic
matrix containing bainitic-austenitic and austenitic
islands. The retained austenite occurs as small blocky
grains or thin layers forming bainitic-austenitic islands.
The strain-induced martensite initially forms in large and
medium-sized austenite grains located along the boun-
daries of the  phase. The transformation is initiated in
the central parts of the grains whereas the border regions
of the austenite remain untransformed. The essential
effect increasing the stability of the retained austenite
against the strain-induced martensite is a fragmentation
of the -phase regions. This effect is additionally
enhanced by the neighboring bainitic-ferrite laths and the
formed plate martensite creating a hydrostatic pressure
against the deformation progress.
5 REFERENCES
1 M. Pouranvari, E. Ranjbarnoodeh, Mater. Tehnol., 46 (2012) 6,
665–671
2 M. Weglowski, K. Kwiecinski, K. Krasnowski, R. Jachym, Arch.
Civ. Mech. Eng., 9 (2009), 85–97
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(2011), 857–865
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tions of Lasers, 8703 (2013), DOI: 10.1117/12.2013429
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16 C. P. Scott, J. Drillet, Scripta Mater., 56 (2007), 489–492
17 J. Wu, P. J. Wray, C. I. Garcia, M. Hua, A. J. DeArdo, ISIJ Int., 45
(2005), 254–262
18 A. Wasilkowska, R. Petrov, L. Kestens, E. A. Werner, C. Krem-
paszky, S. Traint, A. Pichler, ISIJ Int., 46 (2006), 302–309
19 S. Zaefferer, J. Ohlert, W. Bleck, Acta Mater., 52 (2004), 2765–2778
A. GRAJCAR, K. RADWAÑSKI: MICROSTRUCTURAL COMPARISON OF THE THERMOMECHANICALLY ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 679–683 683
Figure 8: Distribution of the retained-austenite grain area of the
cold-deformed steel
Slika 8: Razporeditev podro~ij zrn zaostalega avstenita hladno defor-
miranega jekla
Figure 7: Distribution of crystallographic misorientation angles of the
grains after cold deformation
Slika 7: Razporeditev zrn s kristalografsko razli~nimi koti po hladni
deformaciji

J. HAMPL et al.: CONTROL OF THE METALLURGICAL PROCESSING OF ICDP CAST IRONS
CONTROL OF THE METALLURGICAL PROCESSING OF
ICDP CAST IRONS
KONTROLA METALUR[KE OBDELAVE LITEGA @ELEZA ICDP
Jiøí Hampl1, Tomá{ Válek2, Petr Lichý1, Tomá{ Elbel1
1V[B -Technical University of Ostrava, FMMI, Department of Metallurgy and Foundry, 17. listopadu 15/2172, 708 33 Ostrava, Czech
Republic
2Vítkovické slévárny spol., s. r. o., Halasova 2904/1, Ostrava –Vítkovice, Czech Republic
jiri.hampl@vsb.cz
Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2013-11-18
The article is focused on the use of the measurement of oxygen activity for the management of cast-iron metallurgical
processing in the operating conditions of a centrifugal-roll casting foundry. The paper presents the results of the oxygen-activity
measurement recorded during the metallurgical processing of cast iron from the beginning of the melting to the inoculation of
cast iron. The measurement of oxygen activity was made with specifically developed devices and the probes with a high
sensitivity designed for measuring aO in cast iron by Heraeus Electro-Nite Celox Foundry. The oxygen activities in cast iron
correlate with the properties of cast iron.
Keywords: oxygen activity, ICDP iron, inoculation
^lanek je osredinjen na merjenje aktivnosti kisika za vodenje metalur{ke obdelave litega `eleza v livarni pri centrifugalnem
ulivanju valjev. ^lanek predstavlja rezultate meritev aktivnosti kisika, spremljane med metalur{ko obdelavo litega `eleza od
za~etka taljenja do inokulacije litega `eleza. Meritev aktivnosti kisika je bila izvr{ena s posebno razvito napravo in sondo z
veliko ob~utljivostjo za merjenje aO v litem `elezu s Celox – Foundry Heraeus ElectroNite. Aktivnosti kisika v litem `elezu so
odvisne od lastnosti litega `eleza.
Klju~ne besede: aktivnost kisika, ICDP-`elezo, inokulacija
1 INTRODUCTION
The methods for the control of the metallurgical
processing of cast iron in the molten state are based on
the analyses of solid samples, i.e., the thermal analysis,
the chill test and the spectral analysis. For a quick
interpretation of the metallurgical quality of molten cast
iron, the measurement of oxygen activity can be used
too. The measurements are made in the furnace after
melting, in the last stage before casting, after the
inoculation or modification in the ladle. It is possible to
rapidly analyse the level of metallurgical processing
(quality) of the melt on the basis of the measured oxygen
activity in real time and, if necessary, to make its adjust-
ment.
2 OXYGEN IN CAST IRON
The oxygen content in molten cast iron influences the
mechanism of solidification in the phase of eutectic
transformation and it has positive, but also negative,
effects on molten cast iron. The positive role of oxygen
is primarily to support the formation of stable oxides for
the crystallization of graphite; it also supports the hete-
rogeneous nucleation and stabilises the solidification of
cast iron. An intense formation of graphitization nuclei
occurs during the transition of a melt into the solid pha-
se, especially during the eutectic transformation. During
this solidification phase, the optimum amount of oxygen
has to be available.
A higher oxygen activity is also able to support the
formation of exogenous and endogenous gas bubbles and
pinholes in cast iron, or an increased amount of slag and
inclusions in the castings. Excessive amounts of oxides
can be unstable at elevated temperatures, and under cer-
tain conditions (the temperature, time, and viscosity of
the melt) they dissociate or coagulate, creating an in-
creased amount of slag. The oxygen activity in cast irons
is strongly dependent on the temperature. The oxygen
activity is increased by the liquidus temperature as a
consequence of a release of the crystallization heat. The
heat is secreted from the austenite in the concentration
melt of carbon. This creates good conditions for the
graphite nuclei.
A drop in the liquidus temperature starts a decrease
in the oxygen activity, taking place until the beginning of
the eutectic reaction. The decline in the oxygen activity
is simultaneously accompanied by a formation of oxides.
An increase in the oxygen activity occurs again during
the eutectic transformation, as a consequence of the cry-
stallization-heat eutectic reaction.1,2
The measurement of oxygen activity is normally used
to control the deoxidation process of steel during the
melting and casting of steel castings.
The oxygen activities in molten steel are of a higher-
order, in the range of 10 × 10–6 to 100 × 10–6, depending
on the degree of deoxidation under the temperature of
molten steel. Relatively high levels of oxygen activity in
Materiali in tehnologije / Materials and technology 48 (2014) 5, 685–688 685
UDK 621.74.042:54-145.55:669.787 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)685(2014)
steel during melting correspond to the sensitivity of the
probes used for the measurement (ppm).
Table 1: Informative chemical composition of the ICDP iron (w/%)
Tabela 1: Okvirna kemijska sestava ICDP-`eleza (w/%)
C Mn Si Pmax Smax Cr Ni Mo
3.0
3.5
0.5
1.5
0.7
1.5 0.1 0.03
1.5
2.0
3.8
4.8
0.2
1.0
The oxygen activity in cast iron is by about 3–4
orders of magnitude lower than that of steel, depending
on whether it is measured in the cast iron with lamellar
or spheroidal graphite (Figure 1). The relationship bet-
ween the oxygen activity and the shape of graphite in the
processed FeSiMg cast irons was set with the Mampay
and CELOX-Foundry equipment for measuring oxygen
activity, the Heraeus Electro-Nite company3–5.
The way of the metallurgical processing of cast iron,
i.e., the management of the melt, the holding temperature
and the time greatly influence the oxygen content, i.e.,
its activity in cast iron and the metallurgical quality of
cast iron6–8. The default level of the oxygen activity in
cast iron before an inoculation or modification conse-
quently influences its graphitization ability, the process
of crystallization, the microstructure and the resulting
quality of the castings.
3 METALLURGICAL PROCESSING OF CAST
IRON
The melting of the shell iron (ICDP – Indefinite
Child Double Pour) of the centrifugally cast rolls was
performed in 4-ton electric induction furnaces. The regu-
lation of the chemical composition of iron (alloy) was
performed in the furnaces and the subsequent inoculation
was performed in the ladles. In total, 14 melts were
analyzed. The microstructure of the ICDP iron is formed
by the ledeburite basic metal material (BMM), in which
there is extruded graphite whose surface portion is
optimized in the range of 2–5 % in the evaluated area of
the scratch pattern. The cast iron is controlled during its
melting with a spectral analysis and cooling curves
analyses7,9.
An informative chemical composition of the ICDP
iron is shown in Table 1. The samples are sampled from
the castings for metallurgical analyses, the tests of the
quantity of graphite and of the hardness.
Table 2: Timeline of the melts and the measured values of oxygen
activities, aO
Tabela 2: Potek izdelave taline in izmerjene vrednosti aktivnosti
kisika, a0
Number
of the
melt
End of
melting
(h:min)
Dwell
time of
the charge
(h:min)
aO
(10–9)
Furnace
aO
(10–9)
Ladle
aO
(10–9)
F-L
1 1:30 2:06 1159.9 699.8 460.1
2 1:40 1:40 932 721.2 210.8
3 1:40 1:30 996.1 718 278.1
4 1:40 2:00 829.3 693.5 135.8
5 1:10 3:50 877.6 766.2 111.4
6 1:40 0:45 1522.9 695.4 827.5
7 2:00 1:33 1496 708.7 787.3
8 1:20 0:55 1213.7 726.9 486.8
9 2:00 2:30 1028 700 328
10 1:20 1:20 895 518.5 376.5
11 2:25 6:20 811.4 792.63 18.77
12 1:20 0:55 1034.1 709.04 325.06
13 2:00 2:35 797.7 717.2 80.5
14 1:20 1:15 874.4 717.2 157.2
4 METHODOLOGY FOR MEASURING OXYGEN
ACTIVITY
For the measurement of the oxygen activity we used
Multi-Lab III by Heraeus Electro-Nite, which consists of
a generator, connecting cables and a vibrating lance with
a single measuring probe. The measured values are: the
temperature (°C), the Emf electromotive voltage (mV),
converted to a value of the oxygen activity by the melt
temperature and the oxygen activity, converted to a
reference temperature of 1420 °C. The reference
temperatures are shown in the measurement results. The
measured values are displayed on the display device in
real time approximately 20–30 s after the immersion of
the probe into the melt.
The aim of the measurement was to measure the
oxygen activity after every metallurgical processing of
the ICDP cast iron in real time in the interval of the
melting of the charge, while keeping the iron at the set
temperature until the inoculation in the ladle. The
timeline of the melts (h) and the measured values of the
oxygen activities (aO) are shown in Table 2.
The graph in Figure 2 presents the dependency of the
oxygen activity (aO) on the time (h) of melting in the
furnace. The highest activity was measured in the melts
with the shortest dwell time at the set temperature.
J. HAMPL et al.: CONTROL OF THE METALLURGICAL PROCESSING OF ICDP CAST IRONS
686 Materiali in tehnologije / Materials and technology 48 (2014) 5, 685–688
Figure 1: Oxygen activity in the cast iron with spheroidal, compacted
and lamellar graphite
Slika 1: Aktivnost kisika v litem `elezu s kroglastim, kompaktiranim
in lamelarnim grafitom
With an increasing dwell time in the furnace, the
oxygen activities gradually decrease from the maximum
value of 1522.9 × 10–9 at the dwell time of 45 min up to
the lowest value of 811.4 × 10–9 at the dwell time of
6.3 h.
Figure 3 shows the dependence of the oxygen acti-
vity (aO) measured in the ladle after the inoculation and
the dwell time of the melt in the furnace. For the melt
with the shortest dwell time, the oxygen activity was
measured to be aO = 1522.9 × 10–9 and after the inocu-
lation it was aO = 695.4 × 10–9. In this case, the diffe-
rence in the activities aO before and after the inoculation
is the largest (54 %).
In contrast, for the melt with the longest dwell time,
the lowest activity, aO = 811.4 × 10–9, was measured in
the furnace and after the inoculation in the ladle, it was
aO = 792.63 × 10–9. The decrease was only 18.8 × 10–9.
In this case the difference in the oxygen activity before
and after the inoculation is only 2.5 %.
The differences between the initial oxygen activity aO
in the furnace, before pouring the melt into the ladle, and
the finishing aO, after the inoculation in the ladle (Figure
4) show a similarly decreasing trend as the activity aO in
the furnace depends on the dwell time of the cast iron
(Figure 2).
5 RESULTS AND DISCUSSION
The oxygen activities were measured for the molten
ICDP iron in an electric induction furnace, showing a
large range of measured values (711.5 × 10–9). The
J. HAMPL et al.: CONTROL OF THE METALLURGICAL PROCESSING OF ICDP CAST IRONS
Materiali in tehnologije / Materials and technology 48 (2014) 5, 685–688 687
Figure 4: Difference between the initial oxygen activity aO and the
finishing aO
Slika 4: Razlika med za~etno aktivnostjo kisika a0 in kon~no
aktivnostjo kisika aO
Figure 2: Oxygen activity (aO) dependent on the dwell time (h) of the
melt in the furnace
Slika 2: Odvisnost aktivnosti kisika (aO) od ~asa (h) med dr`anjem
taline v pe~i
Figure 5: Dependence of the oxygen activity in the ladle and the
surface quantity of graphite
Slika 5: Odvisnost med aktivnostjo kisika v ponvi in koli~ino grafita
na povr{ini
Figure 3: Oxygen activity aO in the ladle after the inoculation and the
dwell time
Slika 3: Aktivnost kisika aO v ponvi po inokulaciji in ~asu zadr`anja
measurements were in the range of 1522.9 × 10–9 up to
811.4 × 10–9.
The oxygen activities were measured after the inocu-
lation in the ladle (Figure 2). A relatively narrow range
of values (99.4 × 10–9) from the lowest activity of aO =
693.5 × 10–9 to the highest activity of aO = 792.6 × 10–9
was measured.
The oxygen activities in the iron covered a relatively
wide range after the melting and the dwell time. It was
caused by different melting times. After the inoculation a
narrow range of oxygen activities was measured in the
ladle. This corresponds to the criterion set for the opti-
mum quantity of graphite (2–5 %) in the ledeburite base
of the metal mass of the ICDP iron (Figure 5). The
quantity of graphite was established on the basis of an
image analysis.
6 CONCLUSIONS
On the basis of an evaluation of these melts, it can be
concluded that the value of about aO = 700 × 10–9,
obtained after the inoculation, provides the optimum
level of metallurgical quality of this iron (type ICDP).
By measuring the oxygen activity aO (10–9 ) in real
time, metallurgical processing and quality can be eva-
luated relatively quickly so as to achieve the desired
parameters of the casting.
Acknowledgements
The paper was prepared with the financial support
from the Ministry of Industry and Trade of the Czech
Republic within the TIP project Nr TIP: FR-TI2/188
"Research and development of working-layer materials
of spun-cast multi-layered rolls focused to modern trends
of rolling mills".
7 REFERENCES
1 Z. Bù`ek, Základní termodynamické výpo~ty, Hutnické aktuality,
Informetal, 1988, VÚH@
2 I. C. Kulikov, Raskislenije metalov, Metalurgia, Moskva 1975
3 F. Mampaey, D. Habets, J. Plessers, F. Seutens, On-line oxygen
activity measurements to determine optimal graphite form during
compacted graphite iron production, International Journal of Metal-
casting, 4 (2010) 2, 25–40
4 F. Mampaey, K. Beghym, Oxygen activity in cast iron measured in
induction furnace at variable temperature, AFS Transactions, 114
(2006), 637–656
5 HEN, Heraeus Electro-Nite, Celox-Foundry, CP 10700692, http://
heraeus-electro-nite.com/en/sensorsformoltenmetals/iron/oxygencont
rol_2/oxygencontrol_3.aspx
6 J. Hampl, T. Elbel, On modelling of the effect of oxygen on graphite
morphology and properties of modified cast irons, Archives of
foundry engineering, 10 (2010) 4, 55–60
7 T. Válek, J. Hampl, Prediction of metallurgic quality of ICDP mate-
rial before tapping, 2011 International Conference on Physics Sci-
ence and Technology, Physics Procedia, 22 (2011), 191–196
8 J. Hampl, T. Elbel, Effect of oxygen on graphite morphology and
properties of modified cast irons, Proceedings of the 10th Interna-
tional Foundry Conference, Opatija, 2010
9 T. Válek, J. Hampl, Control of microstructure of cast iron Indefinite
Chill Double Pour – ICDP, Archives of Foundry Engineering, 11
(2011) 4, 199–203
J. HAMPL et al.: CONTROL OF THE METALLURGICAL PROCESSING OF ICDP CAST IRONS
688 Materiali in tehnologije / Materials and technology 48 (2014) 5, 685–688
G. HOJNIK PODREP[EK et al.: SYNTHESIS COMPARISON AND CHARACTERIZATION ...
SYNTHESIS COMPARISON AND CHARACTERIZATION
OF CHITOSAN-COATED MAGNETIC NANOPARTICLES
PREPARED WITH DIFFERENT METHODS
PRIMERJAVA POSTOPKOV IN KARAKTERIZACIJA MAGNETNIH
NANODELCEV, PREVLE^ENIH S HITOZANOM
Gordana Hojnik Podrep{ek, @eljko Knez, Maja Leitgeb
University of Maribor, Faculty of Chemistry and Chemical Engineering, Laboratory for Separation Processes and Product Design,
Smetanova ul. 17, 2000 Maribor, Slovenia
maja.leitgeb@um.si
Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2013-11-26
In this study, magnetic maghemite nanoparticles were prepared with the coprecipitation method, due to its simplicity and
productivity. Thereafter, chitosan-coated magnetic nanoparticles were synthesized with three different methods, the
micro-emulsion process, the suspension cross-linking technique and the covalent binding. Subsequently, a comparison of the
used methods was done using various analyses such as Fourier transform infrared spectroscopy (FTIR), scanning electron
microscopy (SEM), thermogravimetry (TGA), differential scanning calorimetry (DSC), vibrating-sample magnetometry (VSM)
and dynamic light scattering (DLS). The characterization results from Fourier transform infrared (FTIR) spectroscopy and
thermogravimetric analysis (TGA) indicated a successful binding of chitosan on the magnetic nanoparticles. SEM pictures
showed that spherical structured particles with an increased particle size were obtained as the chitosan layer around the particles
was increased. Considering that the magnetic-separation technique has the advantages of rapidity, high efficiency,
cost-effectiveness and lack of negative effect on the biological activity, these carriers may be applied in enzyme immobilization.
Keywords: magnetic nanoparticles, chitosan, surface functionalization
V prispevku je opisana enostavna priprava magnetnih nanodelcev, prevle~enih s hitozanom. Postopek je potekal v dveh
stopnjah. V prvi smo sintetizirali magnetne nanodelce s koprecipitacijo `elezovih ionov. V drugi smo nanodelce prevlekli s
hitozanom, da bi prepre~ili aglomeracijo, z uporabo treh razli~nih postopkov: z metodo mikroemulzije, metodo zamre`enja in s
kovalentno vezavo. Karakterizacija tako pripravljenih nanodelcev je bila izvedena s Fourierjevo transformacijsko infrarde~o
spektroskopijo (FTIR), z vrsti~no elektronsko mikroskopijo (SEM), s termogravimetri~no analizo (TGA), z diferencialno
dinami~no kalorimetrijo (DSC), z analizo vibracijskega magnetometra (VSM) in dinami~nim sipanjem laserske svetlobe (DLS).
Rezultati analiz FTIR in TGA so potrdili vezavo hitozana na magnetne nanodelce, medtem ko je bila oblika in debelina sloja
hitozana dolo~ena s SEM-analizo. Ker ima tehnika magnetnih nanodelcev veliko prednosti pri lo~evanju, cenej{i proizvodnji in
nima negativnih u~inkov na biolo{ko aktivnost, se lahko potencialno uporablja pri encimski imobilizaciji.
Klju~ne besede: magnetni nanodelci, hitozan, povr{inska funkcionalizacija
1 INTRODUCTION
Recently, magnetic nanoparticles such as maghemite
(-Fe2O3) have attracted a great deal of attention due to
their unique, controllable sizes, shapes, other physical
properties, compositions and also for their wide applica-
tions in biomedicine, biotechnology, engineering, mate-
rial science and environmental areas1–3. They have a
magnetic response and can be manipulated with an exter-
nal magnetic-field gradient (Figure 1).
Chitosan, a deacetylated derivative of chitin, is a
polysaccharide with both a hydroxyl and an amine group
in its structure. In addition, it is non-toxic, biocom-
patible, biodegradable, and anti-bacterial4. It is insoluble
in water, but becomes soluble and positively charged in
acidic media. Nowadays, the preparation of chitosan-mo-
dified magnetic nanoparticles are of great interest5,6. In
addition, this polymer has been used successfully to
colloidally stabilize magnetic nanoparticles and has also
been used as a matrix for enzyme immobilization since it
has numerous amino groups that can interact with enzy-
me7,8. The amino groups are responsible for the distinct
characteristics attributed to this basic polymer. There-
fore, the characterization of chitosan is extremely
important with respect to the structure-property relation-
ship, defining a possible industrial application9.
Materiali in tehnologije / Materials and technology 48 (2014) 5, 689–692 689
UDK 620.3:66.017:621.318.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)689(2014)
Figure 1: Magnetic-property illustration of the maghemite nanopar-
ticles dispersed in water
Slika 1: Prikaz magnetnih lastnosti nanodelcev maghemita, dispergi-
ranega v vodi
2 EXPERIMENTS
2.1 Materials
Iron (II) chloride tetrahydrate (FeCl2 · 4H2O), iron
(III) chloride hexahydrate (FeCl3 · 6H2O) and acetic
acid were supplied from Merck, (Germany). Chitosan
(CTS, MMW, the degree of deacetylation was 75–85 %),
glutaraldehyde (GA) and Span-80 were obtained from
Sigma-Aldrich. Ammonia was purchased from Carlo
Erba and paraffin from Kreiger. All the solutions were
prepared with Milli-Q water.
2.2 Apparatus and procedures
Maghemite nanoparticles were synthesized by
coprecipitating Fe2+ and Fe3+ ions in the presence of
ammonium. A functionalization of chitosan was carried
out with three different methods: the micro-emulsion
process10, the suspension cross-linking technique11 and
the covalent binding of chitosan12. These methods differ
with respect to the chitosan concentration, the presence
and concentration of the acetic acid solution, the
glutaraldehyde concentration, the synthesis temperature,
the pH of the medium and the time of the synthesis.
The properties and structures of non-functionalized
magnetic iron-oxide nanoparticles and chitosan-func-
tionalized magnetic maghemite nanoparticles were
characterized with Fourier transform infrared spectro-
scopy (FTIR), scanning electron microscopy (SEM),
differential scanning calorimetry (DSC), dynamic light
scattering (DLS), vibrating-sample magnetometry
(VSM) and potentiometric titration.
3 RESULTS
3.1 Characterization of chitosan-coated magnetic
nanoparticles
The FTIR spectra of chitosan, chitosan-coated
maghemite and maghemite nanoparticles are shown in
Figure 2a to 2c. Undoubtedly, the FTIR spectra of
chitosan showed a broader band at 3410 cm–1, which was
attributed to the hydroxyl (OH) stretching as reported
in6. For the IR spectra of chitosan, the characteristic
absorption bands appeared at 1654 cm–1 which can be
assigned to the N-H bending vibrations, and at 1377 cm–1
assigned to the C-O stretching of the primary alcohol
group in chitosan. For the magnetic maghemite nano-
particles, the peaks at 570 cm–1 and 628 cm–1 were
related to the Fe-O group10. However, the adsorption of
chitosan on the surface of the magnetic iron-oxide nano-
particles was confirmed with the FTIR analysis.
Figure 3 shows the size distribution of the magnetic
nanoparticles determined with DLS in an aqueous
solution with the mean diameter of 22.8 nm.
The results of the TGA characterization of the
maghemite nanoparticles coated with chitosan obtained
with three different methods were used for an estimation
of the amount of the chitosan coating on the maghemite
G. HOJNIK PODREP[EK et al.: SYNTHESIS COMPARISON AND CHARACTERIZATION ...
690 Materiali in tehnologije / Materials and technology 48 (2014) 5, 689–692
Figure 4: Weight-loss curve of chitosan by sample; (MC1) maghemite
particles coated with chitosan with the micro-emulsion process,
(MC2) maghemite nanoparticles coated with chitosan with the suspen-
sion cross-linking technique and (MC3) maghemite nanoparticles
coated with chitosan with the covalent-binding method
Slika 4: Krivulje zmanj{anja mase hitozana v vzorcih; (MC1) delci
maghemita, pokriti s hitozanom s postopkom mikroemulzije, (MC2)
delci maghemita, pokriti z metodo zamre`enja in (MC3) maghemitni
nanodelci, pokriti s hitozanom z metodo kovalentne vezave
Figure 2: FTIR spectra of: a) chitosan, b) maghemite coated with
chitosan and c) maghemite
Slika 2: FTIR-spektri: a) hitozan, b) maghemit z nanosom hitozana in
c) maghemit
Figure 3: Particle-size distribution of non-functionalized magnetic
nanoparticles
Slika 3: Razporeditev velikosti nefunkcionaliziranih magnetnih nano-
delcev
nanoparticles. The actual weight loss of pure maghemite
nanoparticles was subtracted from the actual weight loss
of maghemite nanoparticles coated with chitosan to get
the mass loss of chitosan in percentages for MC1, MC2
and MC3, presented in Figure 4. The mass loss of chi-
tosan for MC3 (22.2 %) is lower than for MC2 (28.8 %)
and MC1 (60.7 %).
The presence of the amino-group amount was studied
for pure chitosan and chitosan-functionalized magnetic
nanoparticles, using a potentiometric titration. The re-
sulting charging isotherms QVt/mf versus pH are pre-
sented in Figure 5. The figure contains the titration data
for pure chitosan and the MC3 sample. The amount of
amino group in free chitosan was 4.22 mmol/g, while in
sample MC3 it was 2.48 mmol/g.
Typical SEM micrographs for maghemite nanoparti-
cles and chitosan-coated maghemite nanoparticles are
shown in Figure 6, presenting maghemite nanoparticles
(a) and maghemite nanoparticles, coated with chitosan
with the covalent-binding method (b). Figure 6 reveals
that the maghemite particles, coated with chitosan
(MC3) have spherical shapes and a size range of 50–100
nm.
Figure 7 shows the magnetization curves for the
maghemite nanoparticles, revealing superparamagnetic
properties. The magnetization of the micro- and nano-
particles, coated with chitosan with the micro-emulsion
process was found to be 3 emu/g, for the suspension
cross-linking technique it was 40 emu/g and for the
covalent-binding method it was 20 emu/g. These values
were compared to the value for the uncoated maghemite
nanoparticles. Therefore, it can be concluded that a
chitosan-surface functionalization of the maghemite
nanoparticles was achieved.
4 CONCLUSION
In this paper, chitosan-coated maghemite nanoparti-
cles were synthesized with the micro-emulsion process,
the suspension cross-linking technique and covalent
binding of chitosan. The samples exhibited clear diffe-
rences in the saturation magnetization, which can be
ascribed mainly to different chemical compositions and
magnetic moments of Fe. The chitosan-coated maghe-
mite nanoparticles appeared in granules with the average
sizes of 40–350 μm after the micro-emulsion process,
400 nm after the suspension cross-linking technique and
50–100 nm after the covalent binding of chitosan. The
size of the nanoparticles was increased with an increase
G. HOJNIK PODREP[EK et al.: SYNTHESIS COMPARISON AND CHARACTERIZATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 689–692 691
Figure 6: SEM images of: a) magnetic maghemite nanoparticles and
b) maghemite nanoparticles, coated with chitosan with the covalent-
binding method (MC3)
Slika 6: SEM-posnetka: a) magnetnih nanodelcev maghemita in b)
nanodelcev maghemita, pokritih s hitozanom s kovalentno metodo
vezanja (MC3)
Figure 7: Hysteresis loops of the maghemite and chitosan-function-
alized maghemite nanoparticles
Slika 7: Histerezna zanka maghemita in s hitozanom funkcionalizi-
ranih nanodelcev maghemita
Figure 5: Charging isotherms of pure chitosan and MC3 nanoparticles
Slika 5: Izoterme naboja ~istega hitozana in MC3-nanodelcev
in the concentration of chitosan and decreased with an
increase in the cross-linker concentration. We found that
the magnetic nanoparticles, coated with chitosan with the
covalent-binding method (MC3) are suitable for practical
applications due to their sufficiently high values of
amino groups and nanosized particles. The maghemite
nanoparticles, functionalized with chitosan, have a po-
tential to be used in assisted drug-delivery systems,
cell/enzyme immobilization, separation processes,
medical diagnosis and therapy and many other industrial
applications.
Acknowledgements
This work was entirely supported by the Slovenian
Research Agency (PhD researcher fellowship contract
No. 1938/FKKT-2010) and financially supported through
Project J2-4232.
5 REFERENCES
1 L. Zeng, R. Hu, Z. Wu, Q. He, Preparation and characterization of
amino-coated maghemite nanoparticles, Bioinformatics and Biome-
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stabilization, functionalization, characterization, and applications,
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alized Nanoscale Materials, Devices and Systems, Springer,
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perties of magnetic Fe3O4-chitosan nanoparticles, Journal of Alloys
and Compounds, 466 (2008), 451–456
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chitosan nanoparticles and its application for Cu(II) removal, Che-
mical Engineering Journal, 168 (2011), 286–292
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methods and characterization of magnetic maghemite coated with
chitosan, Journal of Nanoparticle Research, 15 (2013), 1–12
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kul, Chitosan-functionalized poly(methyl methacrylate) particles by
spinning disk processing for lipase immobilization, Carbohydrate
Polymers, 89 (2012) 3, 842–848
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Elnashar (ed.), Biotechnology of Biopolymers, InTech, 2011, 91
10 H. Y. Zhu, R. Jiang, L. Xiao, W. Li, A novel magnetically separable
-Fe2O3/crosslinked chitosan adsorbent: Preparation, characterization
and adsorption application for removal of hazardous azo dye, Journal
of Hazardous Materials, 179 (2010), 251–257
11 G. Y. Li, Y. R. Jiang, K. L. Huang, P. Ding, J. Chen, Preparation of
magnetic Fe3O4-chitosan nanoparticles, Journal of Alloys and Com-
pounds, 466 (2008), 451–456
12 D. Bhattacharya, S. K. Sahu, I. Banerjee, M. Das, D. Mishra, T. K.
Maiti, P. Pramanik, Synthesis, characterization, and in vitro biolo-
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magnetic iron oxide nanoparticles, Journal of Nanoparticle Research,
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692 Materiali in tehnologije / Materials and technology 48 (2014) 5, 689–692
M. PA[ÁK et al.: STUDY OF PHASE TRANSFORMATIONS IN Cr-V TOOL STEEL
STUDY OF PHASE TRANSFORMATIONS IN Cr-V TOOL
STEEL
[TUDIJ FAZNIH PREMEN V ORODNEM JEKLU Cr-V
Matej Pa{ák, Roman ^i~ka, Pavel Bílek, Peter Jur~i, ¼ubomír ^aplovi~
Institute of Materials Science, Faculty of Materials Science and Technology, Slovak University of Technology, Paulinska 16, 917 24 Trnava,
Slovak Republic
matej.pasak@stuba.sk
Prejem rokopisa – received: 2013-10-01; sprejem za objavo – accepted for publication: 2013-11-12
The properties of steels are very dependent on the phases present in the microstructure. The wear resistance and thermal
stability of tool steels are achieved with the presence of different types of carbides. So the chemical composition and heat
treatment play crucial roles in optimizing the properties of tool steels.
The phase transformations in Cr-V tool steel were analyzed during the heating from room temperature up to 1100 °C using DTA
and dilatometry. After soft annealing the microstructure of the investigated steel consists of a ferritic matrix and M7C3 and MC
carbides, as determined with SEM, EDX and XRD techniques. The experimental results were compared to the computational
results (Thermo-Calc).
Keywords: phase transformation, Cr-V tool steel, Thermo-Calc
Lastnosti jekla so mo~no odvisne od faz, ki so v mikrostrukturi. Odpornost proti obrabi in toplotna stabilnost orodnih jekel se
dose`eta z razli~nimi vrstami karbidov. Kemijska sestava in toplotna obdelava imata klju~no vlogo pri optimiranju lastnosti
orodnih jekel.
DTA in dilatometrija sta bili uporabljeni za preu~evanje faznih premen v orodnem jeklu Cr-V med ogrevanjem od sobne
temperature do 1100 °C. Tehnike SEM, EDX in XRD so potrdile, da po mehkem `arjenju mikrostrukturo preiskovanega jekla
sestavlja feritna osnova ter karbidi M7C3 in MC. Eksperimentalni rezultati so bili primerjani z izra~unanimi (Thermo-Calc).
Klju~ne besede: fazna premena, orodno jeklo Cr-V, Thermo-Calc
1 INTRODUCTION
At present Cr-V ledeburitic steels are often used for
the manufacturing of cutting, forming and other tools in
the industry. To meet the industrial requirements for high
stability and reliability, they have to withstand the wear
and plastic deformation. They are usually produced with
powder metallurgy. This technology enables us to obtain
a uniform carbide size and distribution and also to pro-
duce the materials with particular compositions that
cannot be prepared by conventional casting. The techno-
logy of powder metallurgy provides good results in
achieving high homogeneity. The phase composition,
microstructure and mechanical properties of ledeburitic
tool steels are determined by the matrix and the type,
quantity, size and distribution of the carbides.1 The pro-
perties of the material thus depend on the microstructure
obtained after heat treatment.
The phase transformations during heating can be
described with thermodynamic modeling using the
CALPHAD method.2–4
Thermal-analysis techniques are suitable for an
experimental determination of the phase transformations
in materials. For steels, the DTA and dilatometry tech-
niques are often used.5,6
Carbide phases have a different thermal stability and
some of them are dissolved during the austenitizing in
the solid state.7,8 Bílek et al.9 investigated the phase
composition of Cr-V tool steel after heat treatment, using
SEM+EDX, a quantitative analysis of the microstructure
and hardness measurements. They found two types of
carbides in the microstructure of Cr-V tool steel. Carbide
M7C3 was partially dissolved during the austenitizing at
1000 °C and it completely disappeared after the auste-
nitizing at 1100 °C. Carbide MC is more stable and starts
to dissolve only at a temperature higher than 1200 °C.
The aim of this paper is to enhance and explain the
results published previously by Bílek et al.9 using addi-
tional experimental techniques (XRD, DTA and dilato-
metry) and thermodynamic calculations (Thermo-Calc).
2 EXPERIMENTAL WORK
2.1 Sample material
Cr-V tool steel was used as the experimental material
with the chemical composition of w(C) = 2.1 %, w(Si) =
1 %, w(Mn) = 0.4 %, w(Cr) = 6.8 %, w(Mo) = 1.5 %,
w(V) = 5.4 %, balanced by Fe.10
The sample was initially annealed at the temperature
of 900 °C for 1 h and then slowly cooled down to room
temperature.
The X-ray diffraction (XRD) analysis was accom-
plished with a Panalytical Empyrean X-ray diffracto-
meter. A cobalt anode (U = 40 kV and I = 40 mA) with a
parallel-beam X-ray mirror was used. The sample was
measured with a PIXcel3D detector at room temperature
Materiali in tehnologije / Materials and technology 48 (2014) 5, 693–696 693
UDK 669.14:669.017.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)693(2014)
in the angular range of 45–110° with a step size of
0.0131°. The microstructure and chemical composition
of the phases were analyzed with a JEOL JSM 7600F
electron microscope equipped with a secondary and
back-scattered electron detector and a MAX 50 EDX
detector from Oxford Instruments. The differential
thermal analysis (DTA) was done using a NETZSCH
STA 409CD simultaneous thermal analyzer in an inert
gas (Ar 6.0) with a magnetic frame enabling also the
measurements of the magnetic transition in the tempe-
rature range of 100–1100 °C, at the heating rate of 10
K/min. The dilatometry measurements were performed
using a NETZSCH DIL 402C dilatometer. The inert-gas
atmosphere (Ar 6.0) in the temperature range of
100–1000 °C and different heating rates (2, 5, 8, 12, 16,
20, 25) K/min were used. The thermodynamic calcula-
tions were done using the Thermo-Calc software and
TCFE6 thermodynamic database.
2.2 Results
The XRD analysis (Figure 1) confirms that only MC
and M7C3 carbides are present as the secondary phases in
the ferrite at room temperature.
Figure 2a illustrates the microstructure of Cr-V tool
steel after annealing, with the coarse M7C3 particles
(some of these particles are in an extraordinary fine
form) and finer MC particles in the ferritic matrix.
In the EDX mapping mode the distribution of vana-
dium (Figure 2b) and chromium (Figure 2c) can be
seen.
The phase transformations in the investigated steel
during heating are shown with the DTA and thermomag-
netometry curves in Figure 3. The magnetic transition of
the ferrite from the ferromagnetic to paramagnetic state
is detected at 754.0 °C from the thermomagnetometry
curve and at 753.6 °C from the DTA curve. The next
peak on the DTA curve with the onset at 855.3 °C corres-
ponds to the austenitization process.
These results were verified with thermodynamic
calculations using the Thermo-Calc software. Figure 4
shows the temperature dependence of the volume frac-
tions of individual phases. It can be seen that the amount
of the M7C3 carbide starts to decrease already during the
ferrite-to-austenite transformation, in the temperature
range of 814–837 °C.
After the completion of the ferrite-to-austenite trans-
formation, the dissolution of M7C3 in austenite continues
until its completion at about 1200 °C.
M. PA[ÁK et al.: STUDY OF PHASE TRANSFORMATIONS IN Cr-V TOOL STEEL
694 Materiali in tehnologije / Materials and technology 48 (2014) 5, 693–696
Figure 3: DTA and TG curves of the as-prepared tool-steel powder at
the heating rate of 10 K/min
Slika 3: Krivulji DTA in TG prahu orodnega jekla pri hitrosti ogre-
vanja 10 K/min
Figure 1: Diffractogram of the annealed sample measured at room
temperature
Slika 1: Rentgenski difraktogram `arjenega vzorca, posnet pri sobni
temperaturi
Figure 2: Microstructure of the ledeburitic Cr-V tool steel in the
soft-annealing state: a) overview (SEM), b) EDX map of vanadium
from Figure 2a, c) EDX map of chromium from Figure 2a
Slika 2: Mikrostruktura ledeburitnega orodnega jekla Cr-V v mehko
`arjenem stanju: a) SEM, b) EDX-razporeditev vanadija s slike 2a,
c) EDX-razporeditev kroma s slike 2a
Figure 5 shows the dilatometry curves of Cr-V tool
steel during heating (2 K/min) and cooling (3 K/min). By
analyzing the curve of the thermal-expansion coefficient,
the characteristic temperatures of the ferrite-to-austenite
transformation were determined. The transformation
start and finish temperatures during the heating are thus
843.0 °C and 872.6 °C, respectively. The transformation
start and finish temperatures during the cooling were
determined as 760.2 °C and 732.3 °C, respectively.
The other dilatometry measurements were performed
at different heating rates and the results regarding the
transformation start and finish temperatures of the
ferrite-austenite transformation are summarized in Table
1 and Figure 6.
3 DISCUSSIONS
In this work it was confirmed that the microstructure
of Cr-V tool steel after soft annealing consists of a
ferritic matrix and M7C3 and MC carbides. M7C3 starts to
dissolve during the ferrite-to-austenite phase transforma-
tion and a complete dissolution of M7C3 in the austenite
occurs at the temperature of about 1200 °C. The MC
carbide is more stable and its amount remains unchanged
from room temperature up to 1200 °C. The temperature
range of the ferrite-to-austenite phase transformation
determined with dilatometry is 843.0–872.6 °C using a
slow heating rate (2 °C/min) and the transformation
during cooling occurs in the temperature range of
760.2–732.3 °C. The temperature range of this phase
transformation calculated with Thermo-Calc is 814–837
°C. If the heating rate in dilatometry measurements
increases (2–25 °C/min), the transformation start tempe-
rature also increases from 843.0 °C to 857.1 °C, while
the transformation finish temperature increases from
871.8 °C to 904.0 °C and the temperature range of this
transformation also increases from 28.8 °C to 46.9 °C.
The results for the thermal stability of M7C3 and MC
carbides in Cr-V steel are in agreement with and explain
M. PA[ÁK et al.: STUDY OF PHASE TRANSFORMATIONS IN Cr-V TOOL STEEL
Materiali in tehnologije / Materials and technology 48 (2014) 5, 693–696 695
Figure 6: Transformation start and finish temperatures of the
ferrite-austenite transformation in dependence on the heating rate
Slika 6: Temperature za~etka in konca pretvorbe ferit-avstenit v od-
visnosti od hitrosti ogrevanja
Table 1: Start/finish temperatures of the ferrite-to-austenite transfor-
mation at different heating rates, with the calculated temperature range
of the transformation
Tabela 1: Temperature za~etka – konca pretvorbe ferita v avstenit pri
razli~nih hitrostih ogrevanja z izra~unanimi temperaturnimi obmo~ji
pretvorbe
Heating
rate
2
°C/min
5
°C/min
8
°C/min
12
°C/min
16
°C/min
20
°C/min
25
°C/min
Ts/°C 843.0 845.5 848.2 852.3 855.5 856.3 857.1
Tf/°C 871.8 881.9 887.8 895.6 899.6 902.6 904.0
T/°C 28.8 36.4 39.6 43.3 44.1 46.3 46.9
Figure 4: Volume-phase fractions in dependence on the temperature
calculated with Thermo-Calc
Slika 4: Volumenski dele` faz v odvisnosti od temperature, izra~unane
s Thermo-Calc
Figure 5: Dilatometric curves at the heating rate of 2 °C/min and the
cooling rate of 3 °C/min
Slika 5: Dilatometrijska krivulja pri hitrosti ogrevanja 2 °C/min in pri
hitrosti ohlajanja 3 °C/min
the results published previously.10 The future work will
be focused on developing a kinetic model of the
ferrite-austenite phase transformation in this system
using the dilatometry data. The dissolution of M7C3 in
austenite will be modeled using the Dictra software to
enhance the knowledge about these kinetic processes
occurring in Cr-V tool steel during heating.
4 CONCLUSIONS
The microstructure and phase transformations in
Cr-V tool steel were analyzed using experimental and
computational techniques. The main results are sum-
marized as follows:
• the microstructure of Cr-V tool steel after soft
annealing consists of a ferritic matrix and M7C3 and
MC carbides
• M7C3 starts to dissolve during the ferrite-to-austenite
phase transformation and is completely dissolved in
the austenite at 1200 °C
• MC is more stable and its amount does not change
from room temperature up to 1200 °C
• the phase transformation of ferrite to austenite
proceeds in the temperature range of 843.0–871.8 °C,
determined at the low heating rate by dilatometry by
using the dilatometry data obtained at different heat-
ing rates, a kinetic model of the ferrite-to-austenite
phase transformation will be proposed in the near
future, and the kinetics of the dissolution of M7C3 in
austenite will be calculated using the Dictra software.
Acknowledgements
The authors wish to thank the European Regional
Development Fund (ERDF) for the financial support of
projects ITMS:26220120014 and ITMS:26220120048
"Center for development and application of advanced
diagnostic methods in processing of metallic and
non-metallic materials" funded within the Research &
Development Operational Programme.
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696 Materiali in tehnologije / Materials and technology 48 (2014) 5, 693–696
J. KLOFÁ^ et al.: MODEL ANTIMICROBIAL POLYMER SYSTEM BASED ON POLY(VINYL CHLORIDE) AND ...
MODEL ANTIMICROBIAL POLYMER SYSTEM BASED
ON POLY(VINYL CHLORIDE) AND CRYSTAL VIOLET
MODEL PROTIMIKROBNEGA POLIMERNEGA SISTEMA NA
OSNOVI POLIVINILKLORIDA IN KRISTAL VIOLETA
Jiøí Klofá~1,2, Ivo Kuøitka1,2, Pavel Ba`ant1,2, Kristýna Jedli~ková1,2,
Jakub Sedlák1,2
1Polymer Centre, Faculty of Technology, Tomas Bata University in Zlin, Nam. T. G. Masaryka 275, 762 72 Zlin, Czech Republic
2Centre of Polymer Systems, University Institute, Tomas Bata University in Zlin, Nad Ovcirnou 3685, 760 01 Zlin, Czech Republic
j_klofac@ft.utb.cz
Prejem rokopisa – received: 2013-10-03; sprejem za objavo – accepted for publication: 2013-11-11
The development of novel antimicrobial materials for biomedical application in indwelling devices such as catheters is very
important as the resistance of the pathogens responsible for nosocomial infections towards recent systems has been emerging
with an increasing rate. The work presented here is focused on the preparation and characterization of an antimicrobial
polymeric system composed of poly(vinyl chloride) in combination with crystal violet as a model compound of organic ionic
active species. The antimicrobial activity of the system against gram-negative bacteria Escherichia coli, gram-positive bacteria
Staphylococcus aureus and yeasts Candida albicans as the representative microorganisms were evaluated with a disk-diffusion
test. The release profile of the active substance was observed with UV-VIS spectrometry. The mechanical properties of the
prepared material were tested to verify that they were not altered with respect to the original medical-grade polymer matrix.
Keywords: poly(vinyl chloride), crystal violet, antibacterial, antimicrobial, release
Pomemben je razvoj novega protimikrobnega materiala za biomedicinsko uporabo notranjih pripomo~kov, kot so katetri, ker se
vedno pogosteje pojavlja odpornost patogenov, odgovornih za bolni{ni~ne oku`be. Predstavljeno delo je osredinjeno na pripravo
in karakterizacijo protimikrobnega polimernega sistema, ki ga sestavlja polivinilklorid v kombinaciji s kristal violetom kot
model za spojine organskih ionskih aktivnih vrst. S plo{~inskim difuzijskim preizkusom je bila ocenjena protimikrobna dejav-
nost sistema proti predstavnikom mikroorganizmov: gramnegativni bakteriji Escherichia coli, grampozitivni bakteriji
Staphylococcus aureus in kvasovkam Candida albicans. Profil sprostitve aktivnih snovi je bil opazovan z UV-VIS-spektro-
metrijo. Mehanske lastnosti pripravljenega materiala so bile preizku{ene, da bi potrdili, da niso druga~ne od navadnega
medicinskega polimera.
Klju~ne besede: poli(vinil) klorid, kristal violet, protibakterijski, protimikrobni, spro{~anje
1 INTRODUCTION
Polymers are known for their high versatility and
excellent physical-chemical properties and, in some
cases, they are suitable as biomaterials for the medical
sector and packaging industry. As a candidate for these
applications, the third most common polymer, poly(vinyl
chloride) (PVC), can be considered due to its high
mechanical and chemical resistance, inertness against
biological fluids and a wide range of processing
possibilities.1–3Among many products, urinal catheters,
blood bags and cardiovascular implants are typical items
used in the medical sector; nevertheless, they exhibit a
vulnerability towards surface bacterial colonization.2,4
Therefore, it is important to enhance their antimicrobial
properties by modifying the surfaces of such materials
with a plasma treatment, corona discharge and chemical
grafting.5,6 Another promising strategy of modifying
PVC is to incorporate an antimicrobial substance within
the polymer matrix.4,7,8 Such modifications have long-
term effects and are relatively easy to perform, showing
high rates of success. The solvent-cast technique allow-
ing a preparation of the films with extremely high quality
requirements and great uniformity of the thickness is
often employed due to its advantage of easy blending of
the films with the active molecular compounds soluble in
the used solvent system.9,10 The effectiveness of the anti-
microbial properties of such films is strongly dependent
on the release profile of the antimicrobial substance from
the polymer matrix.11
Organic substances migrate, over time, out of the
polymer matrix and onto the polymer surface and are
then released into the surrounding liquids. Migration
occurs as the organic molecules follow down a concen-
tration gradient and exit the plastic. The migration is
driven by the inherent compatibility differences between
the organic antimicrobials and the polymer substrates in
which they are dispersed. The losses of the organic mole-
cules into the environment are replenished by the
additives within the substrate volume. The benefit of this
mode of action is that it can have a very high activity
rate, and the migratory molecules can very quickly inte-
ract with large numbers of microbes. This does, however,
affect the lifespan of the activity, as the additives leach
out over time, emptying the polymer’s reservoir. The
concentration and choice of the respective organic
additive depend on the level of efficacy required and the
duration of the action needed.12
Materiali in tehnologije / Materials and technology 48 (2014) 5, 697–703 697
UDK 678.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)697(2014)
An antimicrobial agent release that is too fast can be
harmful under certain conditions and perceived as
undesirable. Besides inorganic fillers, ionic organic com-
pounds can provide a real option in the material design
having long-lasting mild effects due to their relatively
slow migration rates. An organic salt structure comprised
of a large bulky organic cation and a small inorganic
anion can be considered to have the migration rates in a
polyolefinic matrix slow enough in comparison with the
molecular organic antimicrobial additives. As a well-
known model representative for this class of compounds
crystal violet (CV) can be chosen. It is a triarylmethane
dye formerly used in medicine due to its antibacterial,
antifungal, anthelmintic and antiseptic properties.13–15
CV was generally considered to be safe for a long time
in the history of medicine; however, many studies have
reported that CV has potentially mutagenic and carcino-
genic effects on humans and animals.16–19 In spite of this,
CV is an excellent model compound for the release-pro-
file studies. This dye can be easily mixed with polymers;
it has a very good and broad antimicrobial activity,
giving deep violet colour to its solutions; hence, its
concentration can be easily monitored with an UV-VIS
absorption spectrometer.
This study focuses on the modification of a medical-
grade PVC with a crystal-violet (CV) addition, using the
solvent-casting technique resulting in a model organic
antimicrobial polymer system. Its composition, morpho-
logy, mechanical properties, antimicrobial tests and the
kinetics of the CV release from a polymeric matrix in
water and physiological solution used as model liquids,
were investigated in the presented work.
2 EXPERIMENTAL WORK
2.1 Materials
Medical-grade thermoplastic plasticized poly(vinyl
chloride) (PVC) compound RB3 was purchased from
Modenplast Medical (Italy). This material is in compli-
ance with the European Pharmacopeia and biocompa-
tible according to ISO 10993, USP, Class VI. Crystal
violet C25N3H30Cl (CV) and cyklohexanone C6H10O
(CYH) were purchased from PENTA (Czech Republic).
All the chemicals were of the analytical grade and used
as received without further purification. Demineralized
water was used for all of these experiments.
2.2 Sample preparation
PVC/CYH/CV films were prepared with the solvent-
casting technique. In the first step a solution of CV in
CYH was prepared: 0.2522 g of CV was added to 250
mL CYH to get a concentration of 1 g/L. 20.005 raw
PVC in the form of granules was dissolved in 300 mL of
CYH during a period 16 h at the room temperature under
continuous stirring. The amount of 250 mL of the CV
solution was added and this blend was left to mix for
another 8 h. Finally, the mixing was finished with a
sonication of the solution for 15 min in an ultrasonic
bath. The solution was then poured into glass dishes and
the solvent was allowed to evaporate at the laboratory
temperature for 10 d. The PVC/CYH control sample was
prepared with the same procedure, but without
incorporating the CV. The conditions for the film
preparation were chosen on the basis of practical
laboratory experience, with the aim to achieve the
smoothest fine films of comparable quality. The
thickness of the resultant films was about 500 μm.
2.3 Characterization
2.3.1 Infrared absorption spectroscopy
A FTIR analysis was used to compare the PVC
pellets, PVC/CYH and PVC/CYH/CV films. All the
measurements were performed with a Nicolet 6700
spectrophotometer (Nicolet, Czech Republic) with the
ATR accessory and the Ge crystal for the attenuated-
total-reflection method.
2.3.2 SEM analysis
The micrographs of the prepared materials were
taken with a Vega II LMU scanning electron microscope
(Tescan, Czech Republic). The freeze fracture surfaces
were obtained with liquid nitrogen and observed after the
coating with a thin layer of gold/palladium by an SC
7640 sputter coater (Quorum Technologies Ltd, UK).
2.3.3 Tensile tests
The effects of the CV added to the PVC matrix on the
mechanical properties were studied using a tensile test.
The specimens for the test were cut from the prepared
film samples as rectangular stripes with the width of 5
mm and the length of 36 mm. The specimens were tested
on a tensile testing machine Testometric M350-5CP
(LABOR machine, Ltd.) at 25 °C according to standard
ISO 37:2005. The speed of the moving clamp was 500
mm/min. The Young´s modulus, the stress at break and
the strain at break were determined. All the samples
were measured in 5 replicates and standard deviations
were estimated.
2.3.4 Antimicrobial tests
The antimicrobial properties of the PVC/CYH/CV
films were assessed using the agar-diffusion test. Round
specimens (8 mm in diameter) were placed on Petri
dishes with the nutrient agar inoculated with the
dispersion of microorganisms (a concentration of CFU
1.0 × 107 mL–1). The samples were tested against gram-
negative Escherichia coli (EC) 4517, gram-positive
Staphylococcus aureus (SA) 4516, and yeast Candida
albicans (CA) CCN 8215. After a incubation 72 h at
23 °C for the yeast and a incubation 24 h at 37 °C for
bacteria, the dimensions of the inhibition zones were
measured in four directions, and the average values were
used to calculate the diameter of the circle-zone inhibi-
J. KLOFÁ^ et al.: MODEL ANTIMICROBIAL POLYMER SYSTEM BASED ON POLY(VINYL CHLORIDE) AND ...
698 Materiali in tehnologije / Materials and technology 48 (2014) 5, 697–703
tion area and its standard deviation. All the tests were
done in triplicates.
2.3.5 Release of CV and plasticizers from the PVC
matrix
Round specimens with a diameter of 12.7 mm were
cut from the PVC/CYH/CV samples to be used in the
release-profile study of CV in the water and physiolo-
gical-solution environment. One specimen was always
placed in a beaker with 50 mL of elution liquid and the
beaker was shaken at 60 r/min to ensure a good homo-
genization of the liquid media. The measurement of the
CV release was performed with a UV-VIS spectropho-
tometer Cary 300 (VARIAN, USA) equipped with a
sipper (a peristaltic pump) and a flow cell (a cuvette).
The whole spectral range (200–800 nm) was monitored
and the spectra were recorded in the preselected time
intervals covering representatively the full time range of
each individual release experiment. The same procedure
was used for obtaining the reference leachate for the
specimens cut from the neat PVC sample with the
thickness of 0.5 mm obtained by hot pressing at 170 °C
for 5 min to evaluate the release of the plasticizers after
three days, which was done to investigate the influence
of either the CV addition or the preparation process on
the release of the plasticizers from the PVC matrix. The
absorbance value at the wavelength of 580 nm was
chosen for a quantitative evaluation of the observed
release profile of CV because this is the position of the
absorption maximum of CV in the elution medium. The
data were then converted to the concentration of the
released CV using calibration curves. The data were
fitted with a non-linear fitting procedure using the
Lavenberg-Marquart algorithm incorporated in the
Origin 7.0 software.
3 RESULTS AND DISCUSSION
3.1 Infrared absorption spectroscopy
The aim of the FTIR analysis was to study a possible
modification of the PVC material with the preparation
process and an addition of CV. In Figure 1, the FTIR
spectra of the neat PVC (pellets), the processed
PVC/CYH sample and the modified PVC/CYH/CV
material are plotted. The infrared absorption spectrum of
an unplasticized polyvinyl chloride contains the bands
typical for the aliphatic CH groups at their most typical
positions, except that, due to the CH2 deformation
vibration, a band is shifted by about 30 cm–1 to the lower
wavenumbers, nearly to 1430 cm–1 as typically observed
for PVC. In addition to the aliphatic CH bands, the
spectra of PVC contain contributions due to the C-Cl
vibrations that can be found as a weak band at 1425 cm–1
and as a medium-intensity band at 959 cm–1. The most
intense and significant band for the C-Cl vibration at
610 cm–1 cannot be observed due to the range of
measurement. In general, the spectrum of the neat PVC
material is greatly affected by the presence of plasti-
cizers and dominated by their absorption bands. The
manifestation of the polymer matrix is, therefore, quite
weak. The most prominent band at 1725 cm–1 can be
assigned to the carbonyl group (C=O) stretching mode
typically observed for the plasticizers. It can be expected
that the medical-grade PVC contains the additives
circumventing the crucial plasticizer migration problem
associated with the softened PVC. The position of this
peak is too low for the aliphatic low-molecular plasti-
cizers such as dioctyl-sebacate, citrate or adipate esters.
On the other hand, a common phthalic ester plasticizer
with a typical manifestation of the carbonyl group at
1720 cm–1 cannot be successfully used as a medical
material intended for modern indwelling applications. A
careful analysis of the wavenumber region between 1500
cm–1 and 1650 cm–1 revealed that there is a quadruplet of
peaks at positions (1540, 1580, 1600 and 1637) cm–1,
while the phthalic ester plasticizers only display doublets
at 1580 cm–1 and 1600 cm–1. Alkyde (based on vegetable
fatty acids) polyanhydrides were found as the highest
scoring records in the available IR spectra database20;
however, the exact identification was impossible. There
is virtually no difference between the spectra recorded
for the neat material and for the PVC/CYH sample,
proving that the RB3 composition did not change during
the solution-casting process and that no solvent residuals
were manifested.
The IR absorption spectrum of the PVC/CYH/CV
sample displays all the characteristic peaks of CV in
addition to the aforementioned spectral features of the
plasticized medical-grade PVC; namely, 1587 cm–1 due
to the C=C stretching in phenyl rings, 1365 cm–1 due to
J. KLOFÁ^ et al.: MODEL ANTIMICROBIAL POLYMER SYSTEM BASED ON POLY(VINYL CHLORIDE) AND ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 697–703 699
Figure 1: FTIR ATR spectra of neat PVC pellets (curve 3), PVC/CYH
film (curve 2) and PVC/CYH/CV film (curve 1) samples. The region
between 2800–1800 cm–1 is hidden in the graph as no absorption
peaks were manifested.
Slika 1: FTIR ATR-spektri vzorcev gladkih PVC-pelet (krivulja 3),
PVC/CYH-plast (krivulja 2) in PVC/CYH/CV-plast (krivulja 1). Pod-
ro~je 2800–1800 cm–1 je skrito v grafu, ker tam ni bilo izrazitih
absorpcijskih vrhov.
the C-H deformation vibrations in methyl groups, 1174
cm–1 due to the C-H in-plane deformation in 1,3,5
substituted aromatic ring, and 1128 cm–1 due to the C-N
stretching vibration in trisubstituted aromatic amines.
3.2 SEM analysis
SEM images were obtained for the freeze-fracture
surfaces of the films prepared with CV. The PVC/CYH/
CV film morphology before the immersion into the
liquid media is shown in Figure 2a and the morphology
of the PVC/CYH/CV film after the release-profile
measurement is shown in Figure 2b. First, there are no
observable crystals of CV in the polymer matrix, and
second, there is no observable change in the material
after the release test.
3.3 Tensile tests
The influence of the PVC modification with CV on
the mechanical properties of the material prepared by
casting from a CYH solution can be seen in Table 1
where the measured values with their standard deviations
are summarized. The mechanical properties of the PVC/
CYH film and the PVC/CYH/CV film are very similar.
The Young’s modulus of PVC/CYH and PVC/CYH/CV
is about 5 MPa and the tensile stress at break is about
11–13 MPa. The only property showing a slight diffe-
rence between the samples is the strain at break. The
PVC with CV shows a higher deformation (elongation)
ability than the pure PVC sample, which can be consi-
dered as an advantage of the material with the additive.
Moreover, the obtained result is in accordance with the
microscopic observation, both testifying a good disper-
sion (blending) of CV in the material.
Table 1: Selected mechanical properties of the PVC/CYH and
PVC/CYH/CV samples and their standard deviations
Tabela 1: Izbrane mehanske lastnosti vzorcev PVC/CYH in
PVC/CYH/CV in njihov standardni odmik
Sample Young’smodulus (MPa)
Strain at break
(%)
Tensile stress
at break (MPa)
PVC/CYH 5.0 ± 0.6 490 ± 40 11 ± 3
PVC/CYH/CV 4.8 ± 0.4 620 ± 50 13 ± 3
3.4 Antimicrobial activity
The results of the antimicrobial-activity halo-zone
test against S. aureus, E. coli and C. albicans, performed
with the agar-diffusion test method, are presented in
Table 2, while Figure 3 demonstrates the inhibition
zones around the samples of the antimicrobial material
(PVC/CYH/CV) studied with the agar-diffusion test. The
obtained values show that the pure PVC material has no
antimicrobial properties, but the PVC with CV shows an
activity against all the tested microorganisms.
Table 2: Antimicrobial activity expressed as inhibition-zone diameters
and their standard deviations for PVC/CYH and PVC/CYH/CV
Tabela 2: Protimikrobna aktivnost, izra`ena kot premer podro~ja
zaviranja in njegov standardni odmik za PVC/CYH in PVC/CYH/CV
SAMPLE SA EC CA
PVC/CYH (mm) 0 0 0
PVC/CYH/CV (mm) 14.8 ± 0.9 10.3 ± 1.0 15 ± 3
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700 Materiali in tehnologije / Materials and technology 48 (2014) 5, 697–703
Figure 4: Release profile of CV from PVC/CYH/CV in the deminera-
lised water. The experimental data points are represented by full-circle
symbols; curve Y represents equation (3) fitted into the data; curves 1,
2 and 3 represent the single-term contributions to curve Y,
respectively. The inset graph shows detailed data from the initial stage
of the experiment.
Slika 4: Profil spro{~anja CV iz PVC/CYH/CV v demineralizirani
vodi. Eksperimentalni podatki so prikazani s polnimi krogci, krivulja
Y ponazarja ena~bo (3), urejeno s podatki; krivulje, ozna~ene z 1, 2 in
3 so posamezni prispevki h krivulji Y. Vstavljeni diagram prikazuje
podrobne podatke iz za~etka preizkusa.
Figure 2: Microphotographs of: a) PVC/CYH/CV before immersion
and b) after 3 d in the liquid
Slika 2: Posnetka: a) PVC/CYH/CV pred potopitvijo v teko~ino in b)
po 3 d
Figure 3: Photographs of Petri dishes after cultivation in the agar-test
diffusion zone against: a) S. aureus, b) E. coli, c) C. albicans
Slika 3: Posnetki petrijevke po kultiviranju v difuzijski coni preizkusa
z agarjem proti: a) S. aureus, b) E. colli, c) C. albicans
3.5 Release profile of CV from PVC/CYH/CV films
The obtained release profiles are shown in Figures 4
and 5 where the dependences of the CV concentration in
the elution liquids are plotted in dependence on the
release time for demineralised water and physiological
solution, respectively. According to the literature, the
first-order kinetic model can be a suitable formal kinetic
description of the process of a water-soluble compound
release from an insoluble polymer matrix to the liquid
medium although it cannot be straightforwardly related
to the sample geometry and it is difficult to concep-
tualize this mechanism on a theoretical basis.21
The release rate of the model compound (CV in our
case) that obeys first-order kinetics can be expressed
with the following equation:
d
d
c
t
c= −
1

(1)
where c is the concentration of the model compound in
the elution media, the expression on the left side of the
equation is the release rate defined as the concentration
increase rate in the elution medium (directly obtained
from absorbance, which is the observable quantity in
this study), t is the release time, –1 is the first-order
release-rate constant. Equation (1) can be integrated into
the following form:
c C exp
t
= − −⎛⎝
⎜ ⎞
⎠
⎟⎛
⎝
⎜ ⎞
⎠
⎟
max 1 
(2)
where parameter Cmax is the integration constant repre-
senting the maximum achievable concentration of the
model compound in the elution media for the infinite
time. With respect to the second boundary condition, it
is assumed that the initial concentration of the model
compound is zero in the liquid media at the beginning of
all the experiments. The rest of the variables and con-
stants have the same meanings as in equation (1).
It can be expected that this formal description only
relates to a limited concentration range and to certain
boundary conditions. According to our observations,
several mechanisms can be active at different time scales
during the release process and, thus, it is reasonable to
extend the kinetic description by one or two more terms
for the first-order processes if they differ significantly in
their rate constants, i.e., by orders of magnitude. The
following equation represents the extension of equation
(2) for three formally independent and additive contri-
butions to the release process:
c C C exp
t
C exp
t
C exp
t
= − ⋅ −
⎛
⎝
⎜
⎞
⎠
⎟ − ⋅ −
⎛
⎝
⎜
⎞
⎠
⎟ −
− ⋅ −
max 1 2
3
  
 
⎛
⎝
⎜
⎞
⎠
⎟
(3)
where C1, C2, C3 represent the maximum contribution of
each process to the infinite time Cmax concentration. The
rest of the variables and constants have the same
meanings as in equation (2) with the indexes showing
their relations to the respective process. It is obvious
that:
C C C C1 2 3+ + = max (4)
For the two processes involved in the release, equa-
tion (3) can be simplified by omitting the third term.
Obtained equation (3) was fitted into the experimen-
tal data as can be seen in Figure 4 representing the water
and Figure 5 representing the physiological solution
where the two-term variant was used. The contributions
of each process are plotted separately with simulated
curves for a better clarity. The obtained parameters are
summarized in Table 3. This mathematical analysis was
performed with a full awareness of the fact that the terms
in equation (3) are not sequential but running parallel as
J. KLOFÁ^ et al.: MODEL ANTIMICROBIAL POLYMER SYSTEM BASED ON POLY(VINYL CHLORIDE) AND ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 697–703 701
Figure 5: Release profile of CV from PVC/CYH/CV in the physiolo-
gical solution. The experimental data points are represented by
full-circle symbols, curve Y represents simplified equation (3) without
the third term fitted into the data; curves 1 and 2 represent the single-
term contributions to curve Y, respectively. The inset graph shows
detailed data from the initial stage of the experiment.
Slika 5: Profil spro{~anja CV iz PVC/CYH/CV v fiziolo{ki raztopini.
Eksperimentalni podatki so prikazani kot simboli polnega kroga,
krivulja Y ponazarja poenostavljeno ena~bo (3) brez vklju~itve tretje-
ga izraza v podatke; krivulji z oznako 1 in 2 pomenita posamezen pri-
spevek h krivulji Y. Vstavljeni diagram prikazuje podrobne podatke iz
za~etka preizkusa.
Table 3: Fitting-equation parameters and their standard errors
describing the release profile of CV from the polymer matrix in the
PVC/CYH/CV sample
Tabela 3: Parametri urejanja ena~b in njihove standardne napake, ki
opisujejo profil spro{~anja CV iz polimerne osnove v vzorcu
PVC/CYH/CV
Environment Water Physiologicalsolution
Cmax /(μg/L) 20.19 ± 0.14 11.15 ± 0.24
C1/(μg/L) 5.20 ± 0.12 2.38 ± 0.06
T1/min 15.2 ± 0.7 66 ± 3
C2/(μg/L) 3.86 ± 0.16 8.63 ± 0.21
T2/min 256 ± 24 4340 ± 220
C3 /(μg/L) 11.09 ± 0.14 n. a.
T3/min 2488 ± 98 n. a.
they use the common time and start at t = 0. However,
the differences in the rate-constant magnitude separate
them to an acceptable level resulting in a good approxi-
mation. Each process (term) dominates its own time-
scale window and relies on its specific concentration
range as it can be seen from the graphs.
The first exponential component (the term with the
shortest time constant, 1) probably represents the release
of CV from the matrix surface, because this process is
the shortest and could only be limited by the CV solu-
bility in water that is 10 g/L as indicated by the supplier.
It is evident, that even the highest CV concentrations in
the elution medium are far from approaching this limit.
In the case of the physiological solution that shares a
chloride anion with CV the solubility must be lower due
to the solubility-product limitation; however, even here
the solubility is more than several orders in magnitude
higher than the observed concentrations. The value of the
solubility product is Ks = 6 × 10–4 mol2 dm–6 estimated
roughly from the CV solubility in water. The solubility
in the physiological solution can be derived by solving
the following equation:
(x + 0.154 mol/L)x = Ks (5)
where x is the maximum CV concentration and 0.154
mol/L is the chloride concentration in the physiological
solution. The equation gives only one positive root, x =
0.0038 mol/L, which corresponds to the CV concen-
tration of 1.55 g/L. This value is about six and a half
times lower than the limitation for the sample in
distilled water.
The second phase of the release process is slower
because the CV readily available from the matrix surface
is already depleted and the CV from the subsurface
layers of the film needs to cross an energetic barrier
before being released into the demineralised water.
The third phase is characterized by a further signi-
ficant decrease-release rate that can be ascribed to the
diminishing of the gradient between the film surface and
the solution layer in its proximity and to the depletion of
the extractable CV in the subsurface of the film. Accord-
ing to the macroscopic observation, the material changed
neither its colour nor any other property after its
immersion into the liquid for several days. No dimension
or significant mass changes were observed which con-
firms there was no swelling or matrix-component disso-
lution. Therefore, we believe that the CV located in the
deeper layers of the material is not released into the
solution in the relevant time horizon.
These three phases were observed for the sample
immersed in the water. The sample in the demineralised
water has a higher saturation value (Cmax) than the
sample in the physiological solution. In general, this
might be caused by the omnipresence of chloride anions
with a relatively high concentration diminishing all the
gradients discussed above in the case of pure water.
The first and second processes of the sample in the
physiological (saline) solution were slow and the third
process was not observed at all. In this case, the solu-
bility of CV is influenced by the presence of the chloride
anion, which is commonly shared between CV and the
physiological solution. Moreover, the CV molecule can
leave the polymer matrix as a CV+ cation and a Cl–
anion, always in a pair, i.e., in the ratio of 1 : 1 due to the
electroneutrality condition that must always be kept.
This condition is obviously satisfied in the case of water,
whereas in the case of the physiological solution this pair
would be released into a medium with a high con-
centration of chloride anions. Alternatively, a lone CV+
cation can be released into the liquid with a concurrent
counter transport of a Na+ cation to the matrix. Both
options can be considered for significantly slowing and
limiting the diffusion process, so only the first two
phases were observed within the time scale of several
days.
3.6 Plasticizer role in the release of CV from PVC/
CYH/CV films
Although the neat PVC material has been approved
for medical use and can be considered as safe from the
point of view of the release of the contained plasticizer
or plasticizers, it must be re-evaluated after being mixed
with CV as the eventual synergic effects cannot be
excluded and the release of the plasticizer could be
enhanced by adding other species to the compound. A
simplified test was performed analysing the absorption
spectra in the wavelength region where both the plasti-
cizer and CV absorb light.
The graph in Figure 6 shows the UV absorption
spectra of the leachates obtained for the PVC/CYH/CV
and neat PVC samples after a three-day elution in water.
The third curve represents the absorption spectrum of
CV in water with the same concentration as that in the
liquid media collected for PVC/CYH/CV. It can be
J. KLOFÁ^ et al.: MODEL ANTIMICROBIAL POLYMER SYSTEM BASED ON POLY(VINYL CHLORIDE) AND ...
702 Materiali in tehnologije / Materials and technology 48 (2014) 5, 697–703
Figure 6: UV absorption spectra recorded for the CV solution in
water (curve 2), leachate from the neat PVC specimen (curve 1) and
leachate obtained for the PVC/CYH/CV specimen (curve 3). For
details, see the text.
Slika 6: Posneta absorpcija UV-spektra za raztapljanje CV v vodi (kri-
vulja 2), izcedna voda iz ~istega PVC-vzorca (krivulja 1) in izcedna
voda iz PVC/CYH/CV-vzorca (krivulja 3). Za podrobnosti glej tekst.
clearly seen, that there is no enhancement of the plasti-
cizer release and that only a simple additivity of the
signals takes place as the absorption spectrum recorded
for PVC/CYH/CV is approximately the sum of the
spectra of the plasticizer released from the neat PVC
sample into the liquid medium and the CV solution. The
test for the physiological solution showed the same
result, but it is not shown here for the sake of brevity.
4 CONCLUSIONS
A model organic antimicrobial polymeric PVC/CYH/
CV system based on medical-grade poly(vinyl chloride)
and crystal violet was prepared with the solvent-casting
technique. The work was focused on investigating the
effects of the used technique for preparing and entering
substances and it was shown that the CYH solvent and
the model CV-active substance did not have any adverse
influence on the chemical structure, morphology and
mechanical properties.
The prepared solvent-cast materials can be used in
the form of a film, as a volume material or as an additive
for further compounding but, preferentially, we aim at
various coatings and thin-film applications on the surfa-
ces of medical devices or other plastic articles wherever
this technique allows hopes for a good adhesion and
compatibility with the substrate material, especially
when coated on the plastic articles made of the same neat
PVC resin.
The antimicrobial activity was investigated using the
agar-diffusion test method and the PVC/CYH/CV
material manifested a good antimicrobial activity against
gram-positive S. aureus, gram-negative E. coli and yeast
C. albicans. Although the material is an organic-doped
antimicrobial polymer system, the release profile of CV,
as the representative model compound with a large orga-
nic cation and halide anion, to the demineralised water
and physiological solution simulating body liquids is
appropriately slow allowing a long-lasting mild delivery
effect of the active species on the closest proximity of
the place of insertion or application. Next, no adverse
effect of either the CV addition to the PVC matrix or the
preparation process on the release of the plasticizers
from the PVC matrix was observed.
These results suggest that the prepared model mate-
rial has a potential in medical plastic industries and the
obtained knowledge can be generalised to a certain
degree, without losing its relevance, covering the whole
class of modelled compounds and used for a further
development of the materials or coatings for PVC
medical devices and hygienic products.
Acknowledgment
The authors wish to thank for the internal grant of
TBU in Zlin, No. IGA/FT/2013/026 funded from the
resources for specific university research.
This article was written with the support of the
Operational Program "Research and Development for
Innovations" co-funded by the European Regional
Development Fund (ERDF) and the national budget of
the Czech Republic, within the "Centre of Polymer
Systems" project (reg. number: CZ.1.05/2.1.00/03.0111).
This article was written with the support of the
Operational Program "Education for Competitiveness"
co-funded by the European Social Fund (ESF) and the
national budget of the Czech Republic, within the
"Advanced Theoretical and Experimental Studies of
Polymer Systems" project (reg. number: CZ.1.07/2.3.00/
20.0104).
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J. KLOFÁ^ et al.: MODEL ANTIMICROBIAL POLYMER SYSTEM BASED ON POLY(VINYL CHLORIDE) AND ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 697–703 703

M. K. BILICI et al.: EFFECT OF THE TOOL GEOMETRY AND WELDING PARAMETERS ...
EFFECT OF THE TOOL GEOMETRY AND WELDING
PARAMETERS ON THE MACROSTRUCTURE, FRACTURE
MODE AND WELD STRENGTH OF FRICTION-STIR
SPOT-WELDED POLYPROPYLENE SHEETS
VPLIV GEOMETRIJE ORODJA IN PARAMETROV VARJENJA NA
MAKROSTRUKTURO, VRSTO PRELOMA IN TRDNOST ZVARA
PRI TORNO-VRTILNEM TO^KASTEM VARJENJU
POLIPROPILENSKIH PLO[^
Mustafa Kemal Bilici1, Ahmet Ýrfan Yükler1, Alim Kastan2
1Department of Materials Technology, Marmara University, 34722 Istanbul, Turkey
2Department of Materials Technology, Afyon Kocatepe University, 03200 Afyon, Turkey
mkbilici@marmara.edu.tr
Prejem rokopisa – received: 2013-10-05; sprejem za objavo – accepted for publication: 2013-11-19
The effect of the tool geometry and welding parameters on the macrostructure, fracture mode and weld strength of the
friction-stir spot welds of polypropylene sheets was studied. Three fracture modes were observed: the nugget pull-out failure,
the cross-nugget failure and the mixed nugget failure under a lap-shear tensile test and nugget debonding and a pull out under a
cross-tension loading, while the lap-shear tensile load was not affected significantly by the delay time. The lap-shear tensile
load and the nugget thickness increased with the increasing tool rotation speed and dwell time. The macrostructure shows that
the welding parameters have a determinant effect on the weld strength (x: the nugget thickness, y: the thickness of the upper
sheet). Finally, when different welding parameters were used, different fracture modes of the joints were obtained in the
friction-stir spot welding of polypropylene sheets. Based on the experimental observation of the macrostructures, the effect of
the welding parameters and tool geometry on the lap-shear tensile load and the fracture mode were discussed.
Keywords: polymers (thermoplastics), polypropylene, friction-stir spot welding, polymer welding, welding parameters
Preu~evan je bil vpliv geometrije orodja in parametrov varjenja na makrostrukturo, vrsto preloma in trdnost zvara pri
torno-vrtilnem to~kastem zvaru polipropilenskih plo{~. Opa`eni so trije na~ini preloma: poru{itev z izpuljenjem jedra, poru{itev
s pretrgom jedra in me{an prelom jedra pri prekrivnem stri`nem preizkusu ter prekinitve v jedru in izpuljenje jedra pri pre~ni
natezni obremenitvi, medtem ko na prekrivno natezno obremenitev ni bilo velikega vpliva. Prekrivna stri`na napetost in
debelina jedra sta nara{~ali z nara{~anjem rotacijske hitrosti orodja in ~asa zadr`anja. Makrostruktura ka`e, da parametri
varjenja igrajo pomembno vlogo pri trdnosti zvara (x: debelina jedra, y: debelina zgornje plo{~e). Kon~no, ~e so bili uporabljeni
razli~ni parametri varjenja, so bili razli~ni tudi na~ini poru{itve pri torno-vrtilnem to~kastem varjenju polipropilenskih plo{~. Na
osnovi eksperimentalnih opa`anj makrostrukture je prikazan vpliv parametrov varjenja in geometrije orodja na prekrivno stri`no
natezno obremenitev in vrsto preloma.
Klju~ne besede: polimeri (termoplasti), polipropilen, torno-vrtilno to~kasto varjenje, varjenje polimera, parametri varjenja
1 INTRODUCTION
Recently, a new joining technique called friction-stir
spot welding (FSSW) or friction spot joining (FSJ) has
been developed1. This technique has the same advantages
as friction-stir welding (FSW) such as the solid-state
process, ease of handling, joining of dissimilar materials
and the materials that are difficult to fusion weld, a low
distortion, excellent mechanical properties and little
waste or pollution. Hence, it is expected to be used for
joining lightweight materials in order to achieve a high
performance and save the energy and costs of the
machines.
The FSSW process consists of three phases: plung-
ing, stirring and retracting2. The process starts with the
spinning of a tool at a high rotation speed. Then the tool
is forced into the workpiece until the shoulder of the tool
plunges into the upper workpiece. The plunge movement
of the tool causes the material to be expelled. When the
tool reaches the predetermined depth, the plunge motion
ends and the stirring phase starts. In this phase, the tool
rotates in the workpieces without plunging. Frictional
heat is generated in the plunging and stirring phases and,
thus, the material adjacent to the tool is heated and soft-
ened. The softened upper and lower workpiece materials
mix together in the stirring phase. The shoulder of the
tool creates a compressional stress on the softened mate-
rial. A solid-state joint is formed in the stirring phase.
When a predetermined bonding is obtained, the process
stops and the tool is retracted from the workpieces. The
resulting weld has a characteristic keyhole in the middle
of the joint created during the friction-stir spot welding.
Different welding-parameter configurations and their
consequences on the weld strength and fracture mode
can be identified. FSSW has been successfully applied to
aluminium3, magnesium4 and steel sheets5, but there are
very few publications on its applications to polymer6–9.
Materiali in tehnologije / Materials and technology 48 (2014) 5, 705–711 705
UDK 678.7:621.791 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)705(2014)
Based on the observations of the FSSW macrostruc-
tures, the weld zone of a FSSW joint is schematically
illustrated in Figure 1. From the appearance of the weld
cross-section, two particular points can be identified10.
The first point is the thickness of the weld nugget (x)
which is an indicator of the weld-bond area (Figure 1).
The weld-bond area increases with the nugget thickness.
The second point is the thickness of the upper sheet
under the shoulder indentation (y). The sizes of these
points determine the strength of a FSSW joint.
In this study, FSSW was performed to join PP sheets
in order to understand the influence of the welding para-
meters and tool geometry on welds. The modifications of
the macrostructural features (the fracture mode, the
nugget thickness and the thickness of the upper sheet)
induced by these different processing configurations are
identified and their consequences on the weld strength
and fracture mode during the lap-shear tensile tests are
discussed. The purpose of this investigation was to
improve the strength of the spot welds. Both the nugget
thickness and the upper-sheet thickness have significant
effects on the fracture modes under tensile loading.
2 EXPERIMENTS
In this investigation 4-mm-thick polypropylene sheets
were used. The polypropylene sheets were purchased
from SIMONA AG, Germany (the tensile yield stress of
34 MPa or lap-shear fracture load of 4500 N). Specimens
were made by using 60 mm by 150 mm sheets with a 60
mm by 60 mm overlap area. In order to carry out the
FSSW tests, a properly designed clamping feature was
utilized to fix the specimens to be welded on an NC
milling machine. The tool was made of the SAE 1040
steel and heat-treated to a hardness of 35 HRC. Four
different tool-pin profiles (straight cylindrical, tapered
cylindrical, threaded cylindrical and triangular) were
used to fabricate the joints (Figure 2).
Each tool had a pin length 5.5 mm and pin size 7.5
mm. The tapered pin had a 15° pin angle. For the straight
cylindrical, tapered cylindrical and threaded cylindrical
pins, the pin size was determined by measuring the
bottom diameter of the pin. For the triangular pins, the
pin size was determined by calculating the diameter of
the cross-section area formed by the turning pin.
In all the cases, the constant tool-plunge rate of 0.26
mm/s and the shoulder-plunge depth of 0.2 mm below
the upper-plate surface were applied. The tool rotation
speeds and the tool dwell times also varied, being
between 560 r/min and 1400 r/min, and between 20 s and
200 s, respectively.
The welded lap-shear specimens were tested on a
ZWICK machine at the constant cross-head speed of 5
mm/s. The load and displacement were simultaneously
recorded during the test. The fracture mode of each
specimen was then determined. The fracture-mode-
appearance observations of the joints were done with a
camera. For the weld-macrostructure studies, thin slices
(of 20 μm) were cut from the welded specimens using a
Leica R 6125 rotary microtome. These thin slices were
investigated using a video spectral comparator. The
photographs of the cross-sections were obtained.
3 EXPERIMENTAL RESULTS
During the FSSW of polyetprophlene, three types of
fracture mode were observed as shown in Figure 3: (a)
the nugget pull-out failure, P; (b) the cross-nugget fail-
ure, C; and (c) the mixed nugget failure, M. By changing
the welding parameters, these three types of fracture
mode were changed. These three types of fracture mode
are affected by the heat amount: insufficient heat, ideal
heat and excessive heat. Consequently, the orientation of
the fracture mode changes the weld strength. While the
highest lap-shear fracture load causes a nugget pull-out
M. K. BILICI et al.: EFFECT OF THE TOOL GEOMETRY AND WELDING PARAMETERS ...
706 Materiali in tehnologije / Materials and technology 48 (2014) 5, 705–711
Figure 3: Three types of fracture mode: a) nugget pull-out failure, b)
cross-nugget failure and c) mixed nugget failure
Slika 3: Tri vrste preloma: a) izvle~enje jedra, b) pre~na poru{itev
jedra in c) me{ana poru{itev jedra
Figure 2: FSSW tool profiles: a) straight cylindrical, b) tapered
cylindrical, c) threaded cylindrical, d) triangular
Slika 2: Profil FSSW-orodja: a) raven valjast, b) sto`~ast valjast, c)
valjast z navojem, d) trikoten
Figure 1: Schematic illustration of the cross-section of a friction-stir
spot weld (x: nugget thickness, y: the thickness of the upper sheet and
t: the total material thickness)
Slika 1: Shematski prikaz prereza torno-vrtilnega to~kastega zvara (x:
debelina jedra, y: debelina vrhnje plo{~e in t: skupna debelina ma-
teriala)
failure, the lowest lap-shear load causes a cross-nugget
failure.
The effect of the tool-pin geometry on the lap-shear
fracture load is shown in Figure 4. In these tests the
welding parameters were kept constant: the tool rotation
speed ranged from 360 r/min to 1400 r/min, the tool-
plunge rate was 0.26 mm/s, the dwell time was 120 s and
the tool-plunge depth was 5.7 mm. The maximum frac-
ture load (900 r/min) was obtained with the tapered
cylindrical pin (4280 N) and the threaded cylindrical pin
resulted in the lowest fracture load (3305 N). Due to the
maximum fracture load obtained with the tapered cylin-
drical pin, all the experiments were made with the
tapered cylindrical pin. The effect of the tool profile on
the welding zone is shown in Figure 5. The three photo-
graphs illustrate that the size of the keyhole formed in
the welding zone depend directly on the pin profile. The
wall slope of the keyhole changed with the pin-angle of
the tool. The nugget thickness was 8.4 mm for the
straight cylindrical pin, 9.9 mm for the 15° angled
tapered pin and 7.8 mm for the threaded cylindrical pin.
The tapered pin created a thicker nugget and a bigger
weld-bond area than the straight cylindrical pin. The
nugget thickness increased with the pin profile changes,
as shown in Figure 5. Also, the nugget thickness exhi-
M. K. BILICI et al.: EFFECT OF THE TOOL GEOMETRY AND WELDING PARAMETERS ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 705–711 707
Figure 5: Effect of tool geometry on weld-nugget formation: a)
straight cylindrical pin, b) 15° angled tapered cylindrical pin and c)
threaded cylindrical pin
Slika 5: Vpliv geometrije orodja na nastanek jedra: a) cilindri~na
konica, b) sto`~asta valjasta konica s kotom 15° in c) valjasta konica z
navoji
Figure 4: Effect of tool profile and tool rotation speed on weld
strength
Slika 4: Vpliv profila orodja in hitrosti vrtenja orodja na trdnost zvara
Table 1: Effect of tapered cylindrical pin on the fracture mode and nugget thickness
Tabela 1: Vpliv valjaste konice na na~in preloma in debelino jedra
Tool rotation speed
(r/min) Fracture mode Macrostructure
Quality of weld-metal
consolidation
Probable reason for the
formation
360 very poor
Insufficient flow of the
joining materials due to
low heat
560 poor
Although there was
insufficient heat, a weld
was formed
710 better than in theprevious case
Heat input is sufficient
for a good-quality weld
900 very good
Heat input is sufficient
for a good-quality weld
1120 good
Heat input is sufficient
for a good-quality weld
1400 worse than in theprevious case
Poor weld quality
occurred due to
excessive heat
bited a very large decrease with the threaded cylindrical
pin, as shown in Figure 5.
The best results were obtained in the experiments
involving the tapered cylindrical pin. For this reason, for
the following experiments, investigating the effect of the
welding parameters the tapered pin was used. In order to
determine the effect of the tool rotation speed on the
FSSW of polypropylene, seven different speeds were
used. For all the welds, the plunge rate was 0.26 mm/s,
the dwell time was 120 s and the tool-plunge depth was
5.7 mm. These three parameters were kept constant and
only the dwell time was allowed to vary in the welding
operations. Increasing the tool rotation speed from 360
r/min to 1400 r/min resulted in a linear progress in the
strength of the welds. The effects of these speeds are
shown in Figure 6. While the lap-shear fracture load dra-
matically increased up to the rotation speed of 900 r/min,
it decreased after 900 r/min. The maximum fracture load
obtained at 900 r/min was 4280 N. The fracture modes
obtained were the cross-nugget failure at the speeds of
360 r/min and 560 r/min, the nugget pull-out failure at
710 r/min, and the mixed nugget failure at (900, 1120
and 1400) r/min. The fracture mode and the macrostruc-
ture are shown in Table 1.
The largest nugget thickness was 9.9 mm obtained at
the tool rotation speed of 900 r/min. The lowest nugget
thickness was 4.5 mm obtained at the tool rotation speed
of 360 r/min. This nugget thickness is very important for
the lap-shear fracture load. At the tool rotation speed of
900 r/min, a thicker nugget and a larger weld-bond area
than at the other tool rotation speeds were obtained. The
nugget thickness and the fracture load increased with the
tool rotation speed as shown in Table 1 and Figure 6.
Figure 7 shows the effect of the dwell time on the
lap-shear tensile strength of FSSW joints. For all the
welds, the plunge rate was 0.26 mm/s, the tool rotation
speed was 900 r/min and the plunge depth was 5.7 mm.
These three parameters were kept constant and only the
dwell time was allowed to vary during the welding ope-
rations. The dwell-time experiments used the tapered
pin. Increasing the dwell time from 20 s to 180 s resulted
in a linear progress in the strength of the welds. The
transition time for the failure modes of the PP sheets 4
mm under the above-mentioned welding parameters was
found to be 45 s. In the period between the dwell times
of 20 s and 120 s, there was a slight increase in the weld
strength. For the dwell times of more than 120 s, the
lap-shear fracture load showed a linear decrease after
120 s. A longer tool stirring time did not affect the weld
strength. Only the fracture mode was changed during the
lap-shear tests. The effect of the dwell time on the weld
cross-sections is shown in Table 2. A thin nugget of the
weld developed over the dwell time of 20 s (Table 2).
The nugget thickness increased with the dwell time
M. K. BILICI et al.: EFFECT OF THE TOOL GEOMETRY AND WELDING PARAMETERS ...
708 Materiali in tehnologije / Materials and technology 48 (2014) 5, 705–711
Figure 6: Effect of tapered cylindrical pin on lap-shear fracture load
Slika 6: Vpliv sto`~asto valjaste konice na prekrivno stri`no obreme-
nitev
Table 2: Influence of dwell time on the fracture mode and nugget thickness
Tabela 2: Vpliv ~asa zadr`anja na vrsto preloma in debelino jedra
Dwell time (s) Fracture mode Macrostructure
Quality of weld-metal
consolidation
Probable reason for the
formation
20 very poor
Insufficient flow of the
joining materials due to
low heat
120 very good
Heat input is sufficient
for a good-quality weld
200 worse than in theprevious case
Poor weld quality due to
excessive heat
Figure 7: Effect of dwell time on lap-shear fracture load
Slika 7: Vpliv ~asa zadr`anja na prekrivno stri`no obremenitev
(Table 2). The maximum fracture load was obtained at
the dwell time of 120 s. Different fracture modes were
obtained during the dwell-time tests: the cross-nugget
failure at the dwell times of 20 s and 45 s; the pull-
nugget failure at the dwell times of (80, 100, 120 and
150) s; and the mixed nugget failure at the dwell times of
60 s and 200 s. The fracture mode and the macrostruc-
ture are shown in Table 2. The largest nugget thickness
was 9.9 mm obtained at the dwell time of 120 s. The
smallest nugget thickness was 1.1 mm obtained at the
dwell time of 20 s. The dwell time is very important for
the nugget thickness. At the dwell time 120 s, a thicker
nugget and a bigger weld-bond area were produced than
at the dwell times 20 s and 200 s, as shown in Table 2.
In order to determine the effect of the delay time on
the FSSW of polypropylene, seven different delay times
were used. Figure 8 shows the effect of the delay time
on the weld strength. In all the welds, the plunge rate
was 0.26 mm/s, the dwell time was 120 s and the plunge
depth was 5.7 mm. These three parameters were kept
constant and only the dwell time was allowed to vary
during the welding operations. The lap-shear fracture
load was on an increase up to the dwell time of 30 s.
After this dwell time, the lap-shear fracture load de-
creased very little. Although the delay time increased up
to seven fold, the weld strength changed only within the
experimental scatter limits (delay times of (0, 10, 20) s).
The highest lap-shear fracture load was obtained after
30 s. For all the delay times the fracture mode was ob-
served to be the nugget pull-out failure. The largest
nugget thickness was 9.9 mm obtained with the delay
time of 45 s. The delay time is very important for the
nugget thickness. The delay time 45 s led to a thicker
nugget and a bigger weld-bond area than the delay times
of 0 s to 30 s, as shown in Table 3.
4 DISCUSSION
The importance of the tool pin profile is shown in
Figure 4. The tapered cylindrical pin resulted in the
biggest and the threaded cylindrical pin resulted in the
lowest lap-shear fracture load. The reason for this diffe-
rence can be easily explained with the weld-nugget
thicknesses that are shown in Figure 5. The straight
cylindrical pin and tapered cylindrical pin have the same
pin size (7.5 mm), but the weld-nugget thicknesses ob-
tained with these pins are different. The nugget thickness
obtained with the straight cylindrical pin was 8.4 mm as
shown in Figure 5a. The tapered cylindrical pin pro-
vided the weld-nugget thickness of 9.9 mm (Figure 5b).
The threaded cylindrical pin provided the weld-nugget
thickness of 7.8 mm (Figure 5c). The tapered pin created
a thicker nugget and a bigger weld-bond area than the
other tools. With the straight cylindrical pin a small
amount of frictional heat was produced in the weld;
therefore, a small weld-bond area and a very low
strength were obtained. The stirring of the pin increased
with the tool rotation speed11. The frictional heat also
increased with the rotation speed. The maximum
strength was obtained with the speed 900 r/min. The
welding residual stresses of the upper sheet increased
with the tool rotation speed12. When the rotation speed
exceeded 900 r/min, the strength decreased because of
the increased residual stresses. The lap-shear fracture
load of a FSSW joint is directly proportional to the
nugget thickness and the weld-bond area10. The tapered
pin produced more frictional heat and a larger weld
thickness, as shown in Figure 5. The heat produced in
the weld area is directly proportional to the welding
parameters and the tool geometry13. Suitable welding
parameters produce more heat and a large weld area,
causing a high weld strength14. Therefore, the tapered pin
produces a higher lap-shear fracture load than the other
pins. As seen in Figure 4, the strength varies signifi-
cantly with the tool pin profile. Figure 5 shows that the
maximum nugget thickness (9.9 mm) changes with the
pin profile. In fact, a change in the pin profile results in a
more extensive stirring and a higher heat input during the
M. K. BILICI et al.: EFFECT OF THE TOOL GEOMETRY AND WELDING PARAMETERS ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 705–711 709
Figure 8: Effect of delay time on lap-shear fracture load
Slika 8: Vpliv ~asa zakasnitve na prekrivno stri`no obremenitev
Table 3: Influence of delay time on the fracture mode and nugget thickness
Tabela 3: Vpliv ~asa zakasnitve na vrsto preloma in debelino jedra
Delay time (s) Fracture mode Macrostructure
Quality of weld-metal
consolidation
Probable reason for the
formation
0 very poor
Insufficient flow of the
joining materials due to
low heat
45 very good
Heat input is sufficient
for a good-quality weld
FSSW, causing the nugget thickness to develop. The
shear fracture of the nugget takes place easily when the
nugget thickness is small, having a low tensile-shear
strength. The results presented in Figure 5 are in agree-
ment with the work up to a certain tool profile (tapered
pin). Different tool profiles resulted in three types of
fracture in the FSSW of PP. These three types of fracture
mode are affected by the heat amount: insufficient heat,
ideal heat and excessive heat. Consequently, frictional
heat occurred with different tool profiles. The fracture-
mode changes associated with the nugget thickness and
weld strength are shown in Figure 5. As a result, the
reason for this strength difference is the chain scission15.
The chain scission lowers the strength of a thermoplastic
material16. If a molten thermoplastic material is heated to
a high temperature and then a high pressure is applied to
it, a decrease in the molecular weight of the material
occurs15. The mechanical properties of thermoplastics
decrease with lowering the molecular weight17. In FSSW
the welding tool produces a compressive pressure in the
weld zone18. In the FSSW of thermoplastics, the material
in the weld area melts7.
Figure 6 shows the effect of the tool rotation speed
of the tapered pin on the weld strength. The lowest
strength was obtained at the tool rotation speed of 360
r/min. In this weld, a small amount of frictional heat was
produced; therefore, a small weld-bond area, a very low
strength and the cross-nugget failure were obtained. The
stirring of the pin increased with the tool rotation
speed13. The frictional heat increased with the rotation
speed. The maximum strength (4280 N) was obtained
with the speed 900 r/min. The welding residual stresses
of the upper sheet increased with the tool rotation
speed12. When the rotation speed exceeded 900 r/min,
the strength decreased because of the increased residual
stresses. The fracture mode changed with the increasing
tool rotation speed. The fracture mode changed due to
the frictional heat. The lowest frictional heat led to the
cross-nugget failure, while the excessive frictional heat
led to the mixed nugget failure (Table 1). A high fric-
tional heat causes a high material temperature in the
welding zone7 and a thicker nugget, as shown in Table 1.
The fracture also occurs due to the high frictional heat in
the welding zone. The tapered pin forms a thicker nugget
and a bigger weld-bond area than the other pins. In the
FSSW fracture experiments, the ideal fracture type is the
nugget pull-out failure due to the high weld strength.
Figure 7 shows the effect of the dwell time of the
tapered pin on the weld strength. The effects of the dwell
time on the weld strength and fracture mode are shown
in Table 2. Short dwell times cause a thin nugget thick-
ness and the cross-nugget failure, as shown in Table 2.
The nugget thickness, bond area and fracture mode have
a direct effect on the weld strength. Increasing the dwell
time from 20 s to 40 s resulted in a linear progress in the
strength of the welds. All these welds were fractured
under small tensile loads because of the small weld-bond
areas. As the dwell time increased, the frictional heat
increased as well. Larger weld nuggets were obtained
with longer dwell times, increasing the joint strength
(Table 2). An increase in the dwell time changed the
fracture mode. The maximum weld strength was
obtained with the nugget pull-out failure, as shown in
Figure 7 (dwell time 120 s). The shape of the weld-bond
area in FSSW is found to be of high importance. The
weld nugget represents the weld bond in FSSW. The
cross-section area of a weld nugget determines the
strength of a weld19. Very high temperatures were
recorded in the FSW of plastics8,12. High melt tempe-
ratures and high welding forces cause chain scission in
the welding zones of the plastics, decreasing the weld
strength20. The physical properties of a polymer are
strongly dependent on the size or length of the polymer
chain. For example, if a chain length is increased, the
melting and boiling temperatures increase quickly as
well. The weld strength also tends to increase with the
chain length, as does the viscosity, or the resistance to
flow, of the polymer in its melt state. The FSSW process
produces high temperature and pressure. But the
excessive heat and pressure cause the chain structure to
break. Most of the molten material is expelled, so a very
small weld stir zone is formed, resulting in a very small
fracture load. Thus, a reduction in the weld strength
occurs. In friction-stir welding it is very important to
check the excessive heat and pressure. Furthermore, the
tool geometry is very important in the production of heat
and pressure.
In this study each mechanical-test diagram shows an
extremum. The lap-shear fracture load increases with the
tool rotation speed (Figure 6), the dwell time (Figure 7)
and the delay time (Figure 8). All these diagrams indi-
cate that there is an optimum value for each welding
parameter. When a variable value exceeds the critical
value, the weld strength starts to decrease. The size of
the weld increases continuously with the welding-para-
meter variables (Figures 6, 7 and 8). For example, the
weld-nugget thickness increases with the tool rotation
speed (Figure 6). The lap-shear fracture load reaches its
highest value with the tapered pin (Figure 6). The reason
for this strength difference is the mechanical scission15.
Mechanical scission lowers the strength of a thermo-
plastic material16. If a liquid thermoplastic material is
heated to a high temperature and then a high pressure is
applied to it, a decrease in the molecular weight of the
material occurs15. The mechanical properties of thermo-
plastics decrease with lowering the molecular weight17.
In FSSW the welding tool produces a compressive pres-
sure in the weld zone18. In the FSSW of thermoplastics,
the material in the weld area melts21. High liquid tem-
peratures and high welding forces cause mechanical
scission in the welding zones of the plastics, which
lowers the weld strength20.
The frictional heat produced in the vicinity of the tool
increased with the dwell time11,22, so the temperature of
M. K. BILICI et al.: EFFECT OF THE TOOL GEOMETRY AND WELDING PARAMETERS ...
710 Materiali in tehnologije / Materials and technology 48 (2014) 5, 705–711
the material increased as well. The temperature of the
material reached the melting temperature (131 °C) in the
dwell time 45 s. The temperature rose up to 142 °C
within the dwell time 50 s and it did not change with the
extended dwell time. Similar temperatures were calcu-
lated for the friction-stir welding of PP sheets12,22. If the
tool was retracted at the end of the predetermined dwell
time, the liquid filled the space of the pin, as shown in
Table 3. This weld does not have a characteristic keyhole
in the nugget. If the pin was retracted with the delay time
45 s, the liquid in the vicinity of the pin cooled down and
solidified. Such a weld has a keyhole as shown in Table
3. During all the experiments, the nugget pull-out failure
occurred at the end of the delay time. Also, an increase
in the delay time increases the weld strength. The dwell
times from 30 s to 60 s resulted in a linear progress in
the strength of the welds. All these welds were fractured
with small tensile loads because of the small weld-bond
areas. Larger weld nuggets were obtained with longer
delay times, which increased the joint strength (Figure
8). The shape of a weld-bond area in FSSW is found to
be of high importance. The cross-section area of a weld
nugget determines the strength of a weld23. A weld with
a small bond area fractures under a low tensile force in
the zero delay time. In the FSSW of PP sheets 4 mm, the
ideal delay time was found to be 45 s.
Fracture modes were changed in accordance with the
welding parameters and tool geometry. This change in
the welds causes the melting and boiling temperatures to
increase quickly. For example, as a chain length is
increased, the melting and boiling temperatures increase
quickly. The weld strength also tends to increase with the
chain length, as does the viscosity, or resistance to flow,
of the polymer in its melt state. Due to these properties
of polypropylene, the weld strength and fracture mode
can be changed. Therefore, in the FSSW of polypro-
pylene, both the fracture mode and the weld strength
were very important.
5 CONCLUSIONS
The macrostructures, the weld strength and the frac-
ture mode of friction-stir spot welds of polypropylene
sheets were investigated.
• In the FSSW of polypropylene, three types fracture
mode were observed: the nugget pull-out failure, the
cross-nugget failure and the mixed nugget failure.
• The weld strength during FSSW was found to mainly
depend on two macrostructural features: the weld-
nugget thickness (X) and the thickness of the upper
sheet under the shoulder indentation (Y).
• With the increasing rotation speed (up to 900 r/min),
the lap-shear fracture load decreased because of the
increased amount of the heat generated by the
mechanical scission in the stir zone.
• The tool rotation speed and dwell time must be
sufficient to allow a high weld strength and the
appropriate fracture mode.
• During the FSSW of polypropylene, the pin geometry
affects the nugget formation and lap-shear fracture
load.
• The optimum tool for sheets 4 mm was found to be
the tapered cylindrical pin.
• Mechanical scission can occur during the FSSW of
polypropylene, if excessive frictional heating is
created in the weld zone, so the optimum welding
parameters should be chosen (the tool rotation speed
of 900 r/min, the dwell time of 120 s and the delay
time of 45 s).
6 REFERENCES
1 Y. Tozaki, Y. Uematsu, K. Tokaji, Fatigue Fracture Materials Struc-
ture, 30 (2007), 143–148
2 P. C. Lin, J. Pan, T. Pan, Journal Fatigue, 30 (2008), 90–105
3 M. Merzoug, M. Mazari, L. Berrahal, A. Imad, Materials Design, 31
(2010), 3023–3028
4 Y. C. Chen, K. Nakata, Materials Design, 30 (2009), 3913–3919
5 M. I. Khan, M. L. Kuntz, P. Su, A. Gerlich, T. North, Y. Zhou, Tech-
nology Welding Joining, 12 (2007), 175–182
6 M. K. Bilici, A. I. Yukler, M. Kurtulmus, Materials Design, 32
(2011), 4074–4079
7 A. Arýcý, S. Mert, Journal of Reinforce Plastics and Composite, 1
(2008), 1–4
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Letters, 64 (2010), 2098–2101
9 M. K. Bilici, A. I. Yukler, Materials Design, 33 (2012), 145–152
10 Y. Tozaki, Y. Uematsu, K. Tokaji, International Journal of Machine
Tools and Manufacture, 47 (2010), 2230–2236
11 N. Ma, A. Kunugi, T. Hirashima, K. Okubo, M. Kamioka, Welding
International, 23 (2009), 9–14
12 M. Aydin, Polymer Plastics Technology Engineering, 49 (2010),
595–601
13 M. Awang, V. H. Mucino, Z. Feng, S. A. David, SAE International, 1
(2005), 1251–1256
14 S. J. Vijay, N. Murugan, Materials Design, 31 (2010), 3585–3589
15 H. M. Costa, V. D. Ramos, M. C. G. Rocha, Polymer Testing, 24
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Engineering Science, 47 (2007), 1813–1819
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Australia Rheology Journal, 2 (2004), 57–62
18 A. Gerlich, T. H. North, M. Yamamoto, Science and Technology of
Welding and Joining, 12 (2007), 472–480
19 R. S. Mishra, Z. Y. Ma, Materials Science Engineering, 50 (2005),
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21 W. Yuan, R. S. Mishra, S. Web, Y. L. Chen, B. Carlson, D. R. Her-
ling, G. J. Grant, Journal of Materials Processing Technology, 211
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22 J. E. Mark, Polymer Data Handbook, Oxford University Press, New
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M. K. BILICI et al.: EFFECT OF THE TOOL GEOMETRY AND WELDING PARAMETERS ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 705–711 711

I. IVANI] et al.: MICROSTRUCTURAL ANALYSIS OF CuAlNiMn SHAPE-MEMORY ALLOY ...
MICROSTRUCTURAL ANALYSIS OF CuAlNiMn
SHAPE-MEMORY ALLOY BEFORE AND AFTER THE
TENSILE TESTING
ANALIZA MIKROSTRUKTURE ZLITINE CuAlNiMn Z
OBLIKOVNIM SPOMINOM PRED NATEZNIM PREIZKUSOM IN
PO NJEM
Ivana Ivani}1, Mirko Goji}1, Stjepan Ko`uh1, Borut Kosec2
1University of Zagreb, Faculty of Metallurgy, Aleja narodnih heroja 3, 44103 Sisak, Croatia
2University of Ljubljana, Faculty of Natural Sciences and Engineering, A{ker~eva cesta 12, 1000 Ljubljana, Slovenia
iivanic@simet.hr
Prejem rokopisa – received: 2013-10-07; sprejem za objavo – accepted for publication: 2013-11-18
In this paper the results of a microstructural analysis before and after fracture along with the mechanical properties and hardness
of the CuAlNiMn shape-memory alloy are presented. The melting of the alloy was carried out in a vacuum-induction furnace in
a protective atmosphere of argon. The alloy was cast into an ingot of 15 kg. After casting the alloy was forged and rolled into
rods with a diameter of approximately 10 mm. A microstructural characterization was performed with light microscopy (LM)
and scanning electron microscopy (SEM) equipped with energy-dispersive spectrometry (EDS). Martensitic microstructure was
observed in the rods after the deformation. The fractographic analysis of the samples after the tensile testing revealed some
areas with intergranular fracture. However, the greater part of the fracture surface indicated the pattern of transgranular brittle
fracture. The results of the tensile tests showed the tensile strength of 401.39 MPa and elongation of 1.64 %. The hardness of the
CuAlNiMn alloy is 290.7 HV0.5.
Keywords: shape-memory alloy, CuAlNiMn, fracture analysis, microstructure, hardness
V prispevku so predstavljeni rezultati analize mikrostrukture pred prelomom in po njem skupaj z mehanskimi lastnostmi in
trdoto zlitine CuAlNiMn z oblikovnim spominom. Taljenje zlitine je bilo izvedeno v vakuumski pe~i v za{~itni atmosferi
argona. Zlitina je bila ulita v ingot mase 15 kg. Po litju je bila zlitina kovana in zvaljana na premer pribli`no 10 mm.
Karakterizacija mikrostrukture je bila izvedena s svetlobno mikroskopijo (SM) in vrsti~no elektronsko mikroskopijo (SEM),
opremljeno z energijskim disperzijskim spektrometrom (EDS). Analizirana je bila martenzitna mikrostruktura zlitine
CuAlNiMn pred izvedenim nateznim preizkusom. Izvedena sta bila natezni preizkus in meritve trdot. Fraktografska analiza je
pokazala ve~ podro~ij z interkristalnim in pogosto transkristalnim krhkim prelomom. Rezultati nateznega preizkusa so pokazali,
da je natezna trdnost 401,39 MPa in raztezek 1,64 %. Trdota zlitine CuAlNiMn je 290,7 HV0,5.
Klju~ne besede: zlitina z oblikovnim spominom, CuAlNiMn, analiza preloma, mikrostruktura, trdota
1 INTRODUCTION
Shape-memory alloys (SMAs) based on copper such
as CuZnAl and CuAlNi are attractive for practical appli-
cations because of their special properties (shape-me-
mory effect and pseudoelasticity) which are based on the
crystallographic reversible thermoelastic martensitic
transformation. They are also suitable due to lower costs
(compared to NiTi) and the advantages with regard to
electrical and thermal conductivities.1–5
However, the polycrystalline copper-based shape-
memory alloys with coarse grains are very brittle and
they are prone to intergranural fracture because of the
high elastic anisotropy of the parent  phase, the exi-
stence of the brittle 2 (Cu9Al4) phase and the formation
of the stress-induced martensites along the grain boun-
daries upon quenching.6–10 The usual way to improve the
disadvantages mentioned above is to alloy them with the
elements that are grain refiners like Ti, B and Zr, which
create the precipitates limiting the grain size and grain
growth. Also, the production of the alloy with the
rapid-solidification technique or powder metallurgy is
very effective in obtaining a fine-grain microstructure.2,11
The mechanical properties of the polycrystalline CuAlNi
alloy can be improved effectively with grain refinement
and texture control. Both of them play important roles in
relaxing the stress concentration at grain boundaries,
which prevents intergranular fracture and improves the
plasticity of the alloy. The fatigue and memory proper-
ties of the CuAlNi alloy with fine grains are considerably
limited because the fine grains inevitably tend to grow
during hot-working or heat treatment, leading to a
degradation of mechanical properties.6,12
Manganese is added as an alloying element to
improve the ductility of the CuAlNi alloy by replacing
the partially aluminum content. It also increases the
stability domain of the  phase and allows the betatising
process to be performed at lower temperatures.8,12 Also,
manganese was provided to enhance the thermoelastic
and pseudoelastic behavior.11,13
The aim of this paper was to carry out the strength
testing, hardness measurements, microstructural
characterization and a fractographic analysis of the
CuAlNiMn alloy.
Materiali in tehnologije / Materials and technology 48 (2014) 5, 713–718 713
UDK 669.35:620.17 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)713(2014)
2 EXPERIMENTAL WORK
The CuAlNiMn shape-memory alloy was produced
by melting in a vacuum-induction furnace in a protective
atmosphere of argon and cast into a classical iron mould
with the dimensions of 100 mm × 100 mm × 200 mm.
The heating temperature was 1330 °C. After casting the
alloy was forged and rolled with step heating at 900 °C
after every reduction into the bars with a diameter of
approximately 10 mm. From the bars the samples were
prepared as standard round tensile-test probes with the
dimensions of  6 mm × 100 mm. The tensile test was
made with a Zwick/Roell Z050 universal tensile-testing
machine at room temperature. Light microscopy (LM)
and scanning electron microscopy (SEM) equipped with
energy-dispersive spectroscopy (EDS) were applied for
the microstructural characterization of the alloy. For the
microstructural analysis, the samples were grinded
(120–800 grade paper) and polished (0.3 μm Al2O3).
After polishing, the samples were etched in a solution
composed of 2.5 g FeCl3 and 48 mL methanol in 10 mL
HCl for 15 s. A fractographic analysis using a JOEL
JSM5610 scanning electron microscope was carried out
to observe the surfaces of the samples after the tensile
testing. The hardness of the alloy was carried out with
the Vickers method with the applied force of 5 N.
3 RESULTS AND DISCUSSION
The average chemical composition of the alloy
measured with EDS was Cu-8.05 % Al-3.51 % Ni-2.44 %
Mn (w/%).
3.1 Microstructural characterization before fracture
The obtained microstructures are presented on Fig-
ures 1 to 3. It can be observed that the microstructure of
the alloy after the deformation (forging and rolling) is
martensitic. Because of the plastic deformation after the
casting, it can be assumed that most of the martensite in
I. IVANI] et al.: MICROSTRUCTURAL ANALYSIS OF CuAlNiMn SHAPE-MEMORY ALLOY ...
714 Materiali in tehnologije / Materials and technology 48 (2014) 5, 713–718
Figure 3: a) SEM micrograph of the CuAlNiMn shape-memory alloy
with the positions marked for EDS analysis, b) EDS spectrum for
position 1 and c) EDS spectrum for position 3
Slika 3: a) SEM-posnetek mikrostrukture zlitine CuAlNiMn z obli-
kovnim spominom z ozna~enimi mesti za EDS-analizo, b) EDS-spek-
ter na mestu 1 in c) EDS-spekter na mestu 3
Figure 1: LM micrograph of the CuAlNiMn shape-memory alloy,
magnificaton 100-times
Slika 1: Mikrostruktura zlitine CuAlNiMn z oblikovnim spominom,
pove~ava 100-kratna
Figure 2: SEM micrograph of the CuAlNiMn shape-memory alloy
Slika 2: SEM-posnetek mikrostrukture zlitine CuAlNiMn z obli-
kovnim spominom
the structure is the stress-induced martensite. The grains
appear clearly and the martensite plates have different
orientations in individual grains. We also noticed a large
grain size which commonly appears in Cu-based alloys.12
In our previous paper14 it was observed that the
CuAlNiMn SMA in the as-cast state has a martensitic
microstructure with some areas of the 2 phase.
The CuAlNi ternary phase diagram, Figure 4, shows
the main phases that appear in the CuAlNi shape-me-
mory alloy. The  phase, which is essential for the
shape-memory effect, can exist independently and stably
above 565 °C. The eutectoid reaction (   + 2) occurs
in the alloy at 565 °C. Meanwhile, the adequate cooling
rate can suppress the eutectoid reaction, and the  phase
can be totally transformed into martensites when the
temperature decreases to Ms (the martensite start tempe-
rature).2,6
On the SEM micrographs (Figures 2 and 3), marten-
sitic microstructure is confirmed. Several inclusions can
be noticed and their chemical compositions (positions 1
and 2 on Figure 3a) are presented in Table 1. It can be
seen that the inclusions contain the highest amount of
Mn, along with Cu, Al and Ni (which are the main
constituents of the alloy), and there are also other
elements – "impurities" (Fe, Cr, P) (Figure 3b).
Table 1: Results of the EDS analysis before tensile testing for the
positions marked on Figure 3a, (w/%)
Tabela 1: Rezultati EDS-analize, izvedene pred nateznim preizkusom
na mestih ozna~enih na sliki 3a, (w/%)
Position Cu Al Ni Fe Mn Cr P
1 25.20 1.73 3.58 8.59 40.65 4.75 15.50
2 17.28 1.86 3.82 9.63 46.18 5.23 15.99
3 86.07 8.04 3.53 – 2.36 – –
4 85.95 8.06 3.48 – 2.51 – –
The usual chemical composition of the CuAlNi SMA
is Cu-(11–14) % Al-(3–4.5) % Ni (w/%). According to
the literature12 an addition of manganese replacing the
aluminum content is effective as it can improve the
ductility. The EDS analysis of positions 3 and 4 (Figure
3a) and the EDS spectrum on Figure 3c confirm such a
replacement.
U. Sari13 investigated the influence of the mass frac-
tion w = 2.5 % of manganese on the CuAlNi SMA, and
he concluded that, due to a manganese addition, the
grain size, which is over 1 mm for CuAlNi, is reduced by
75 %, to the average value of 350 μm.
3.2 Fracture analysis of the CuAlNiMn shape-memory
alloy
The results of the SEM microfractography analysis
after the tensile testing of the CuAlNiMn shape-memory
alloy are presented on Figures 5 to 8.
It can be seen that a crack often occurs at a three-fold
node of grain boundaries, Figure 5. It is known that the
brittleness of copper-based alloys arises from the high
degree of order in the parent phase with B2, DO3 and L21
structures; the brittleness was also attributed to their high
elastic anisotropy (A ≅ 13) which is a reason for the
brittle-grain-boundary cracking.15–17 The second cause is
I. IVANI] et al.: MICROSTRUCTURAL ANALYSIS OF CuAlNiMn SHAPE-MEMORY ALLOY ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 713–718 715
Figure 5: SEM microfractographs of the CuAlNiMn shape-memory
alloy after tensile testing: a) magnification 100-times and b) the
magnified section
Slika 5: SEM-posnetek preloma zlitine CuAlNiMn z oblikovnim
spominom po izvedenem nateznem preizkusu: a) pove~ava 100-kratna
in b) pove~ano obmo~je
Figure 4: Ternary phase diagram of the CuAlNi alloy; a vertical
cross-section at the mass fraction of Ni w = 3 % 9
Slika 4: Ternarni fazni diagram zlitine CuAlNi; vertikalni prerez pri
masnem dele`u Ni w = 3 % 9
the grain size of -phase alloys which is usually in order
of 1 mm.12,13 The causes mentioned above probably
constitute the essential differences between NiTi alloys
and Cu-based alloys, influencing the fracture behavior. A
large stress concentration occurs at the grain boundaries
due to a large elastic anisotropy under loading. The
result is that very brittle intergranular cracking occurs
even during elastic deformation.16
It may be assumed that the cracks nucleate at the
grain-boundary nodes where the stress concentrations
develop.15 This assumption is confirmable with the
cracks visible on Figures 5 to 7. Grain boundaries
provide the easiest crack-propagation path. The cracks
nucleate at the grain boundaries where the stress-level
concentration is high and the intergranular fracture is
obtained. It is mostly a transgranular type of fracture
with the characteristic river pattern that can be observed
(Figures 5 and 6) but sporadic intergranular fracture can
also be noticed. At a higher magnification it is visible
that the plane of fracture displays the river patterns
typical of a cleavage – like a rupture (Figures 5b and
6b).18 On Figures 7a and 7b there are parallel lines near
the grain-boundary plane that probably represent the
stress-induced martensite. There are some small and
shallow dimples on the fracture surface of the
investigated alloy, indicating that the alloy underwent a
certain plastic deformation during the fracture (Figure
7c).
The fracture surface was examined with an EDS
analysis (Figure 8), and the chemical composition of the
fracture surface is presented in Table 2. It can be noticed
that the amount of Cu was from 88.15–93.55 %, Al was
from 1.78–5.81 %, Ni was from 3.19– 3.46 % and Mn
was from 1.25–2.60 % (w/%). Lower concentrations of
alloying elements probably influence the fracture
mechanism and properties of the alloy by decreasing the
strength in the region of grain boundaries. Reduced
concentrations of the alloying elements at the grain
boundaries are probably due to slow cooling rates and
I. IVANI] et al.: MICROSTRUCTURAL ANALYSIS OF CuAlNiMn SHAPE-MEMORY ALLOY ...
716 Materiali in tehnologije / Materials and technology 48 (2014) 5, 713–718
Figure 7: SEM microfractographs of the CuAlNiMn shape-memory
alloy: a) magnification 100-times, b) the magnified section and c) the
area with transgranular brittle fracture – the magnified section
Slika 7: SEM-posnetek preloma zlitine CuAlNiMn z oblikovnim
spominom: a) pove~ava 100-kratna in b) pove~ano obmo~je ter c)
podro~je z transkristalnim krhkim prelomom – pove~ano obmo~je
Figure 6: SEM microfractographs of the CuAlNiMn shape-memory
alloy: a) magnification 100-times and b) the magnified section
Slika 6: SEM-posnetka preloma zlitine CuAlNiMn z oblikovnim
spominom: a) pove~ava 100-kratna in b) pove~ano obmo~je
low solidification velocities that are the consequences of
the alloy-casting procedure.
Table 2: Results of the EDS analysis after tensile testing for the
positions marked on Figure 7a, (w/%)
Tabela 2: Rezultati EDS-analize po izvedenem nateznem preizkusu na
mestih ozna~enih na sliki 7a, (w/%)
Position Cu Al Ni Fe Mn
1 88.96 4.89 3.46 0.22 2.47
2 93.55 1.78 3.19 0.23 1.25
3 88.15 5.81 3.44 0.00 2.60
3.3 Mechanical properties of the CuAlNiMn shape-
memory alloy
The results obtained after the tensile testing are given
in Table 3. The tensile strength/elongation curves are
presented in Figure 9. The tensile strength was 401.39
MPa, calculated as the average value of three measure-
ments. The Young’s modulus and yield strength were
67.72 GPa and 242.81 MPa, respectively. According to
the literature,19 this value of the tensile strength is
satisfactory for a Cu-based alloy. The elongation (A)
after fracture was very low (1.64 %) and without a
measurable contraction. In the literature19 the maximum
elongation after tensile testing for a continuously cast
Cu-13 % Al-4 % Ni (w/%) SMA was 1.45 % and this is
below the limit of the typical recoverable strain of 4–6 %.
U. Sari13 found that the compression strength for Cu-11.6 %
Al-3.9 % Ni-2.5 % Mn (w/%) amounts to 952 MPa and
the ductility is 15 %.
The hardness of the CuAlNiMn shape-memory alloy
was 290.7 HV0.5. As manganese favorably influences
the alloy plasticity, it is fair to assume that the hardness
of the CuAlNiMn alloy should be lower than that of the
alloy without manganese.13,14
4 CONCLUSIONS
The microstructure analysis of CuAlNiMn before the
tensile testing shows the presence of a martensitic micro-
structure. According to the plastic deformation carried
out after the casting, it is fair to assume that the marten-
site existing in the microstructure is stress induced. The
fracture surface indicates intergranular fracture and
mainly transgranular brittle fracture with the characte-
ristic river pattern. There are also some parts with
shallow dimples indicating that the alloy was plastically
deformed. The cracks nucleate at the grain boundaries
where the stress-level concentration is high. Mechanical
properties of CuAlNiMn show satisfactory results for the
I. IVANI] et al.: MICROSTRUCTURAL ANALYSIS OF CuAlNiMn SHAPE-MEMORY ALLOY ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 713–718 717
Figure 8: a) SEM microfractograph of the CuAlNiMn shape-memory
alloy with the positions marked for EDS analysis and b) EDS
spectrum of position 1
Slika 8: a) SEM-posnetek preloma zlitine CuAlNiMn z oblikovnim
spominom z ozna~enimi mesti za EDS-analizo in b) EDS-spekter na
mestu 1
Table 3: Tensile-test results for the CuAlNiMn shape-memory alloy
Tabela 3: Rezultati nateznega preizkusa zlitine CuAlNiMn z oblikovnim spominom
Mechanical properties Young’s modulus / GPa Yield strength / MPa Tensile strength / MPa Elongation / %
Measurement results
with mean value
60.76
67.72
244.04
242.81
350.78
401.39
1.41
1.6466.94 245.99 416.10 1.74
75.45 238.39 437.30 1.79
Figure 9: Tensile stress – elongation curves of the CuAlNiMn SMA
Slika 9: Krivulje natezna napetost – raztezek zlitine CuAlNiMn SMA
tensile strength (401.39 MPa) and a very low value for
the elongation (1.64 %). The hardness of the alloy is
290.7 HV0.5.
Acknowledgement
The authors want to thank professor Franc Kosel and
Brane Struna (University of Ljubljana) for the tensile
testing and technical information.
5 REFERENCES
1 M. Goji}, L. Vrsalovi}, S. Ko`uh, A. Kneissl, I. An`el, S. Gudi}, B.
Kosec, M. Kli{ki}, Journal of Alloys and Compounds, 509 (2011),
9782–9790
2 G. Lojen, I. An`el, A .C. Kneissl, A. Kri`man, E. Unterweger, B.
Kosec, M. Bizjak, Journal of Materials Processing Technology,
162–163 (2005), 220–229
3 M. ^oli}, R. Rudolf, D. Stamenkovi}, I. An`el, D. Vu~evi}, M. Jen-
ko, V. Lazi}, G. Lojen, Acta Biomaterialia, 6 (2010), 308–317
4 Y. Sutou, T. Omori, K. Yamauchi, N. Ono, R. Kainuma, K. Ishida,
Acta Materialia, 53 (2005), 4121–4133
5 Y. Sutou, R. Kainuma, K. Ishida, Materials Science and Engineering
A, 273–275 (1999), 375–379
6 Z. Wang, X. F. Liu, J. X. Xie, Progress in Natural Science, Materials
International, 21 (2011), 368–374
7 A. C. Kneissl, E. Unterweger, M. Bruncko, G. Lojen, K. Mehrabi, H.
Scherngell, Metalurgija, 14 (2008) 2, 89–100
8 C. Segui, E. Cesari, Journal de Physique IV, Colloque C2, supple-
ment au Journal de Physique III, 5 (1995), 187–191
9 K. Otsuka, C. M. Wayman, Shape memory alloys, Cambridge Uni-
versity Press, Cambridge 1998, 97–116
10 Z. G. Wei, H. Y. Peng, W. H. Zou, D. Z. Yang, Metallurgical and
Materials Transactions A, 28 (1997), 955–967
11 Z. Li, Z. Y. Pan, N. Tang, Y. B. Jiang, N. Liu, M. Fang, F. Zheng,
Materials Science and Engineering A, 417 (2006), 225–229
12 J. Van Humbeck, L. Delaney, A comparative review of the (Potential)
Shape Memory Alloys, ESOMAT 1989 – Ist European Symposium
on Martensitic Transformations in Science and Technology, Germa-
ny, 1989, 15–26
13 U. Sari, International Journal of Minerals, Metallurgy and Materials,
17 (2010) 2, 192–198
14 I. Ivani}, M. Goji}, I. An`el, S. Ko`uh, M. Rimac, O. Beganovi}, K.
Begovi}, B. Kosec, Microstructure and properties of casted CuAlNi
and CuAlNiMn shape memory alloys, Proceedings of the 13th Inter-
national Foundrymen Conference, Opatija, 2013, 153–162
15 G. N. Sure, L. C. Brown, Metallurgical Transactions, 15A (1984),
1613–1621
16 H. Sakamoto, Y. Kijima, K. Shimizu, Transactions of the Japan Insti-
tute of Metals, 23 (1982) 10, 585–594
17 Y. Sutou, T. Omori, R. Kainuma, N. Ono, K. Ishida, Metallurgical
and Materials Transactions A, 33 (2002), 2817–2824
18 L. A. Matlakhova, E. C. Pereira, A. N. Matlakhov, S. N. Monteiro, R.
Toledo, Materials Characterization, 59 (2008), 1630–1637
19 G. Lojen, M. Goji}, I. An`el, Journal of Alloys and Compounds, 580
(2013), 497–505
I. IVANI] et al.: MICROSTRUCTURAL ANALYSIS OF CuAlNiMn SHAPE-MEMORY ALLOY ...
718 Materiali in tehnologije / Materials and technology 48 (2014) 5, 713–718
B. MAŠEK et al.: EFFECT OF THE INITIAL MICROSTRUCTURE ON THE PROPERTIES ...
EFFECT OF THE INITIAL MICROSTRUCTURE ON THE
PROPERTIES OF LOW-ALLOYED STEEL UPON
MINI-THIXOFORMING
VPLIV ZA^ETNE MIKROSTRUKTURE NA LASTNOSTI
MALOLEGIRANEGA JEKLA PO PREDELAVI
MINI-THIXOFORMING
Bohuslav Ma{ek, David Ai{man, Hana Jirková
University of West Bohemia in Pilsen, Research Centre of Forming Technology – FORTECH, Univerzitní 22, 30614 Pilsen,
Czech Republic
daisman@vctt.zcu.cz
Prejem rokopisa – received: 2013-10-19; sprejem za objavo – accepted for publication: 2013-12-06
Thixoforming is an unconventional forming process based on semi-solid processing. Semi-solid processing involves heating the
feedstock to a temperature at which it becomes partly liquid and partly solid. Thanks to partial melting and rapid solidification,
unconventional microstructures can be obtained even in conventional materials. Today’s research efforts mainly involve high-al-
loyed steels with a wide freezing range at lower temperatures. By contrast, low-alloyed steels with the solidification tempera-
tures above 1400 °C are not under extensive investigation.
The main objective of the study was to find whether and how the attributes of the feedstock microstructure can be transmitted to
the final semi-solid processed microstructure of the 30MnVS6 microalloyed steel. The first stage of the experiment involved
mini-thixoforming of the as-received specimens of the steel. In order to assess the effects of the initial microstructure, speci-
mens with three types of microstructure were prepared with high-pressure torsion (HPT), a severe-plastic-deformation (SPD)
technique. Initial and final microstructures were examined using light and electron microscopes and image-analysis techniques.
The phase composition was determined with the aid of the X-ray diffraction analysis. The mechanical properties were deter-
mined with tensile and hardness testing. Information on the local properties and the properties of microstructure constituents
was obtained with microhardness measurement.
Keywords: minithixoforming, 30MnVS6, severe plastic deformation, high pressure torsion
Thixoforming je neobi~ajen postopek preoblikovanja, ki temelji na preoblikovanju v testastem stanju. Preoblikovanje v
testastem stanju vklju~uje ogrevanje preoblikovanca do temperature, pri kateri postane material delno v staljenem in delno v
strjenem stanju. Po zaslugi delnega taljenja in hitrega strjevanja se lahko dose`e neobi~ajne mikrostrukture celo pri navadnih
materialih. Sedanje raziskave se izvajajo ve~inoma na mo~no legiranih jeklih s {irokim intervalom strjevanja pri ni`jih
temperaturah. Nasprotno se malolegirana jekla s temperature strjevanja nad 1400 °C ne preiskuje intenzivno.
Glavni cilj te {tudije je ugotoviti, ali in kako atribute mikrostrukture surovca prenesti po preoblikovanju v testastem stanju pri
mikrolegiranem jeklu 30MnVS6. Prva stopnja eksperimentov je vklju~evala minithixoforming preoblikovanje dobavljenih
vzorcev jekla. Za oceno vpliva za~etne mikrostrukture so bili pripravljeni vzorci s tremi vrstami mikrostrukture s torzijo pri
visokem tlaku (HPT), to je s tehniko velike plasti~ne deformacije (SPD). Za~etne in kon~ne mikrostrukture so bile preiskane s
svetlobno mikroskopijo, z elektronsko mikroskopijo in s tehnikami analize slik.
Sestava faz je bila dolo~ena z rentgensko difrakcijsko analizo. Mehanske lastnosti so bile dolo~ene z nateznim preizkusom in z
merjenjem trdote. Informacija o lokalnih lastnostih in lastnostih posameznih mikrostrukturnih faz je bila dobljena z meritvami
mikrotrdote.
Klju~ne besede: minithixoforming, 30MnVS6, plasti~na deformacija, torzija pri visokem tlaku
1 INTRODUCTION
Mini-thixoforming was developed for processing
small volumes of a material using highly dynamic heat-
ing and cooling processes.1,2 Thanks to these characte-
ristics, the processing induces transformations that are
not commonly encountered in the conventional proces-
ses. As a consequence, various types of microstructure
are obtained which contain metastable constituents, such
as austenite in ferritic-pearlitic steel.3 Owing to the
highly dynamic nature of the phenomena involved, the
initial condition of a material cannot be neglected
because the material’s history is often reflected in its
final microstructure. This fact frequently fails to be
acknowledged and has not been explored adequately.
The objective of the present experiment was to compare
the character of the resulting microstructure with the
initial condition. The method used to modify the
microstructures was a severe-plastic-deformation
technique known as high-pressure torsion (HPT)4 and it
involved various amounts of the applied strain.
2 EXPERIMENTAL PROGRAMME
The experimental programme was conducted with the
30MnVS6 microalloyed steel (Table 1). This steel,
alloyed predominantly with manganese and silicon, is
typically used for making forged parts. The as-received
material was in the form of rolled and peeled normalized
bars. The as-received microstructure consisted of ferrite
and pearlite (Figure 1). The ultimate and yield strengths
in the as-received condition were 770 MPa and 570 MPa,
respectively, and the hardness was 237 HV10.
Materiali in tehnologije / Materials and technology 48 (2014) 5, 719–723 719
UDK 669.018.258:539.374 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)719(2014)
With regard to the low content of alloying elements, a
tentative simulation using the JMatPro program was used
to ascertain that the freezing range of the material was
relatively narrow and lying at high temperatures.5 The
temperature interval was 1425–1490 °C. The tempera-
ture to be used for achieving the semi-solid state region
was set at 1455 °C.
2.1 Refinement of the initial microstructure using
HPT
The structure of the material was refined using the
HPT method, one of the severe-plastic-deformation tech-
niques. Its principle is a torsional deformation between
stationary and rotating dies that compress the material
(Figure 2).6 The magnitude of the strain introduced
depends on the number of revolutions. The feedstock
was a disc with a 35 mm diameter. The forming process
reduced its height to 7.5 mm and expanded its diameter
to 38 mm (Figure 3). The forming was performed at
ambient temperature. The material did not undergo any
recrystallization. The effect of the initial microstructure
was explored using three specimens upon various
numbers of revolutions: 3.5, 4.25 and 7.
2.2 Microstructure after HPT
As the amounts of the strain in the centre and by the
edge of the formed disc differ, the metallographic section
was prepared on the disc face. The observations with
both the light and electron microscopes confirmed that
the resulting microstructure was heavily distorted, in
contrast to the as-received material. The most severe dis-
tortion was found in the specimen upon the highest num-
ber of revolutions, which was seven (Figure 4a). Its
microstructure still contained the ferrite and pearlite
mixture but both phases were heavily deformed (Figure
4b). Its hardness, up to 502 HV10, was higher than that
of the as-received material. The surprising finding was
that the pearlite lamellae underwent a heavy deformation
without any fracturing. Another interesting feature was
B. MAŠEK et al.: EFFECT OF THE INITIAL MICROSTRUCTURE ON THE PROPERTIES ...
720 Materiali in tehnologije / Materials and technology 48 (2014) 5, 719–723
Figure 4: a) Microstructure of 30MnVS6 steel upon HPT processing,
7 revolutions,; ) detail of distorted pearlite lamellae
Slika 4: a) Mikrostruktura jekla 30MnVS6 po HPT-predelavi, 7 vrtlja-
jev, b) detajl izkrivljene lamele perlita
Figure 2: Schematic of the HPT process6
Slika 2: Shematski prikaz HPT-postopka6
Figure 1: As-received microstructure of 30MnVS6 steel
Slika 1: Mikrostruktura jekla 30MnVS6 v dobavljenem stanju
Figure 3: Final products of the HPT process
Slika 3: Kon~ni proizvod HPT-postopka
Table 1: Chemical composition of 30MnVS6 steel in mass fractions (w/%)
Tabela 1: Kemijska sestava jekla 30MnVS6 v masnih dele`ih (w/%)
C Mn Si P S Cu Cr Ni Al N Mo V Ti
0.31 1.50 0.62 0.018 0.024 0.03 0.20 0.02 0.02 0.0122 0.007 0.098 0.0261
that upon three or even four revolutions, the microstruc-
tures in the centre of a disc and by its edge showed little
difference. This was evidenced by similar hardness levels
for all the areas of the deformed disc. Upon 3.5 revolu-
tions, the hardness was 421 HV10. Upon 4.25 revolu-
tions, it rose to 465 HV10.
2.3 Preparation of the workpieces for semi-solid pro-
cessing
The non-uniform strain intensity was reflected in the
not quite uniform microstructures of the discs upon HPT.
The difference was most notable between the outermost
edge and the centre of a disc. For this reason, the edge
and the centre of a disc were not used for further experi-
ments. Cylinders of a 15 mm height were cut from the
discs. This dimension is consistent with the length of the
active part of the workpiece for mini-thixoforming. The
remaining conical parts of the workpiece were made for
low-carbon steel with a high melting temperature. The
entire workpiece for semi-solid processing was thus as-
sembled from three parts and electron-welded in a vac-
uum (Figure 5). This low-energy welding method left
the special microstructure in the active part of the
workpiece unaffected. The workpieces from the as-re-
ceived material were made in one piece. Their length
was 48 mm and their diameter was 6 mm (Figure 6).
The purpose of their conical ends is to carry the electric
current for heating and to centre the workpieces inside
the die.
3 SEMI-SOLID PROCESSING
All the workpieces were formed in a titanium die
with a groove-shaped cavity with a 20 mm length and a
5 mm × 1.9 mm cross-section. This die was specially
B. MAŠEK et al.: EFFECT OF THE INITIAL MICROSTRUCTURE ON THE PROPERTIES ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 719–723 721
Figure 7: Schematic illustration of the whole process with the microstructure development after individual steps
Slika 7: Shematski prikaz celotnega postopka z razvojem mikrostrukture po vsakem koraku
Figure 6: Final workpiece welded together with an electron beam
Slika 6: Kon~na oblika obdelovanca, zvarjenega z elektronskim
curkom
Figure 5: Schematic of the workpiece manufacturing sequence
Slika 5: Shematski prikaz korakov priprave obdelovanca
developed for the mini-thixoforming process.7 With
respect to the high heating temperature, an R-type
thermocouple (rhodium-platinum) was used. The heating
temperature was found using stepwise optimization.
Despite the temperature being high, it was maintained
accurately, as was the heating rate (Figure 7).
The initial experiments were conducted with the
as-received specimens. Only after all the mini-thixo-
forming parameters were determined, the HPT-processed
specimens were employed. First, a suitable heating tem-
perature was sought. A total of three temperatures were
tried: (1450, 1455 and 1460) °C. The heating time was
61 s. The deformation applied within 0.3 s was followed
by rapid solidification enhanced by the contact with the
metal cavity surface. The initial cooling rate reached
300 °C/s.
4 DISCUSSION OF RESULTS
At the first stage, the semi-solid processing parame-
ters were optimized for the 30MnVS6 steel employed
outside its typical domain. Mini-thixoforming of the
as-received non-refined material produced a martensitic
microstructure with some amount of bainite. This is un-
usual, considering the typical thixo-formed micro-
structure, which consists of polyhedral austenite grains
embedded in a ledeburite network. In the specimen
heated to 1460 °C, the etchant used for outlining prior
austenite grains (the picric acid) revealed that the major-
ity of the product and the interior of the workpiece had a
dendritic morphology. In response to this finding, the
heating temperature was reduced. However, with the
temperature reduced by 10 °C, the cavity was not filled
completely because the liquid fraction was inadequate.
The most suitable heating temperature proved to be
1455 °C. Once the processing parameters were opti-
mized, the product filled the die cavity completely. The
microstructure in the centre of the product contained
mostly martensite and some bainite (Figures 8a and 8b).
Its hardness was approximately 767 HV10. The etchant
for the prior austenite grain boundaries revealed that the
proportion of the dendrites in the product decreased.
Besides the dendrites, there were polygonal shapes of the
prior austenite grains in the microstructure. These shapes
were found predominantly in the product formed by
forcing the semi-liquid material into the groove in the
titanium die. A higher proportion of bainite was found in
the workpiece interior, whose cooling rate was lower
than that of the product.
Processing the workpieces made from the discs upon
3.5 revolutions led to a similar character of the micro-
structure. It consisted of a mixture of martensite and
bainite. The hardness in the centre of the product reached
no more than 500 HV10. The variation in the micro-
structure was only found after outlining the prior auste-
nite grain boundaries by etching, whereas the normally
etched martensitic-bainitic microstructure appeared
uniform. The morphology of the prior austenite grains in
the mini-thixoformed HPT-refined material was poly-
gonal. In contrast to the as-received material, no large
dendritic areas were found upon mini-thixoforming
(Figures 9a and 9b). Etching the material processed with
HPT with 4.25 revolutions and picric acid revealed areas
with very fine globular particles with the size of
approximately 20 μm (Figure 9c). The matrix consisted
of martensite and bainite, as in the previous cases. The
hardness in the centre of the product was 713 HV10. No
effects of a further refinement were found in the speci-
mens upon 7 revolutions.
B. MAŠEK et al.: EFFECT OF THE INITIAL MICROSTRUCTURE ON THE PROPERTIES ...
722 Materiali in tehnologije / Materials and technology 48 (2014) 5, 719–723
Figure 8: Micrographs of the as-received material upon semi-solid
processing: a) workpiece centre, b) centre of the extruded product
Slika 8: Mikrostruktura dobavljenega materiala po predelavi v testa-
stem stanju: a) sredina obdelovanca, b) sredina vzorca po ekstruziji
Figure 9: Comparison between the morphologies of prior austenite grains upon semi-solid processing of the material with two different initial
microstructures: a) micrograph of as-received material, b) micrograph of HPT-refined material, c) detail of fine globular grains
Slika 9: Primerjava videza prvotnih avstenitnih zrn po predelavi v testastem stanju materiala z dvema razli~nima za~etnima mikrostrukturama: a)
mikrostruktura dobavljenega materiala, b) mikrostruktura materiala po HPT-postopku, c) detajl drobnih globulitnih zrn
5 CONCLUSION
The effect of the initial microstructure of the micro-
alloyed 30MnVS6 steel on its final microstructure upon
semi-solid processing was explored. The material in the
two initial states was mini-thixoformed under identical
conditions. The first state was characterized by a fer-
rite-pearlite microstructure with the hardness of 237
HV10 and the other was a plastically deformed micro-
structure caused by high-pressure torsion. The hardness
level was between 421 and 502 HV10, depending on the
strain intensity.
All the products made with mini-thixoforming con-
tained martensite with a small fraction of bainite. The fi-
nal hardness was high, reaching 713 HV10 in some
cases. The refinement of the initial microstructure was
not reflected substantially in the semi-solid processed
microstructure. After the heaviest deformation, the work-
piece material had a very fine morphology of the prior
austenite grains. In order to characterize them, additional
microscopic techniques will have to be used such as
EBSD to find their size or X-ray diffraction analysis to
describe their texture.
Acknowledgement
The authors gratefully acknowledge the funding by
the German Research Foundation (Deutsche Forschungs-
gemeinschaft, DFG) and the Czech Science Foundation
(Grantové agentura ^eské republiky, GA^R) through
joint, binational projects WA 2602/2-1 and GA ^R
P107/11/J083.
6 REFERENCES
1 H. Mi{terová, A. Rone{ová, [. Jení~ek, Mini-thixoformnig of tool
steel X210Cr12, 22nd International Conference on Metallurgy and
Materials, Metal 2013, Brno, 2013
2 I. Sen, H. Jirkova, B. Masek, M. Böhme, M. F. X. Wagner, Micro-
structure and Mechanical Behavior of a Mini-Thixoformed Tool
Steel, Metallurgical and Materials Transactions A, 43 (2012) 9,
3034–3038
3 M. N. Mohammed, M. Z. Omar, J. Syarif, Z. Sajuri, M. S. Salleh, K.
S. Alhawari, Microstructural Evolution during DPRM Process of
Semisolid Ledeburitic D2 Tool Steel, The Scientific World Journal,
(2013), article ID 828926
4 T. C. Lowe, R. Z. Valiev (eds.), Investigations and Applications of
Severe Plastic Deformation, Kluwer Academic Publishers, Dordrecht
2000
5 JMatPro, The Java-based Materials Property Simulation Package,
Version 6.1, Sente Software Ltd., Surrey Technology Center, UK
6 T. Hebesberger, H. P. Stüwea, A. Vorhauer, F. Wetscher, R. Pippan,
Structure of Cu deformed by high pressure torsion, Acta Materialia,
53 (2005) 2, 393–402
7 B. Ma{ek et al., Minithixoforming of high chromium tool steel
X210Cr12 with various initial states, Annals of DAAAM for 2011 &
Proceedings of the 22nd International DAAAM Symposium, Vienna,
Austria, 2011
B. MAŠEK et al.: EFFECT OF THE INITIAL MICROSTRUCTURE ON THE PROPERTIES ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 719–723 723

A. AKKURT: EXPERIMENTAL INVESTIGATION OF THE SURFACE PROPERTIES OBTAINED BY CUTTING BRASS-353 ...
EXPERIMENTAL INVESTIGATION OF THE SURFACE
PROPERTIES OBTAINED BY CUTTING BRASS-353 (+)
WITH AN ABRASIVE WATER JET AND OTHER
CUTTING METHODS
PREISKAVA LASTNOSTI POVR[INE MEDENINE 353 (+) PO
REZANJU Z ABRAZIJSKIM VODNIM CURKOM IN DRUGIMI
METODAMI REZANJA
Adnan Akkurt
Gazi University, Faculty of Technology, Department of Industrial Design Engineering, Ankara, Turkey
aakkurt@gazi.edu.tr
Prejem rokopisa – received: 2013-10-26; sprejem za objavo – accepted for publication: 2013-11-18
In the manufacturing industry different methods are used to provide the fastest, cheapest and the most cost-effective way of
facilitating the process of cutting with the minimum surface deformation. Apart from the conventional methods, non-traditional
methods such as abrasive water jet (AWJ), laser, plasma, underwater plasma and wire erosion are used intensely for the cutting
of hard-to-cut materials and products. Research has been conducted on the AWJ method. Brass materials are widely used in
industry. In this study the results of the cutting process for brass material with AWJ were investigated. Based on the results the
ideal cutting method for the investigated material was found to be AWJ.
Keywords: cutting methods, unconventional cutting, surface properties
V industriji se uporabljajo razli~ne metode za hitro, cenej{e in stro{kovno bolj ugodne metode rezanja z minimalno deformacijo
povr{ine. Poleg navadnih metod za rezanje trdih materialov in proizvodov se uporabljajo tudi netradicionalne, kot je abrazijsko
rezanje z vodnim curkom (AWJ), laser, plazma, podvodna plazma in `i~na erozija. Izvr{ene so bile raziskave AWJ. Medenina se
pogosto uporablja v industriji. V tej {tudiji je bil preiskan postopek rezanja medenine z AWJ. Glede na dobljene rezultate je
ugotovljeno, da je za preiskovani material najbolj{a metoda abrazijsko rezanje z vodnim curkom.
Klju~ne besede: metode rezanja, neobi~ajno rezanje, lastnosti povr{ine
1 INTRODUCTION
Cutting quality can be determined by measuring the
surface roughness, dimensional tolerances, etc. In the
cutting processes for different materials, there are no sig-
nificant differences in general macro-morphological sur-
face properties. For example, the surface obtained on cut
glass is the same as on metal, ceramic and composites.
However, when examined at the micro-level, micro-qual-
ities of surfaces vary depending on the differences be-
tween the cutting mechanisms of different methods. The
properties of the surfaces obtained with an abrasive wa-
ter jet are listed below:
• The surface is not affected by thermal impacts or
heat.
• No crusting is found on brittle materials. Surface is
almost free of refractions.
• An insignificant hardness alteration may occur on the
surface.
• The width of the cut may be narrowed depending on
the diameter of the jet.
• Abrasive fragment sedimentation may occur in the
material.
• Small chamfers may occur in the holes to the surface.
The quality of the obtained surface could be im-
proved by increasing the power spent for each unit of the
cutting length. A better surface quality is obtained by in-
creasing the water pressure, decreasing the jet speed, in-
creasing the abrasive magnitude rate in the jet and select-
ing a larger nozzle. Abrasive-surface properties as well
as abrasive-particle shape and dimension are important
factors. The width of the cutting channel is controlled
with the mixture tube nozzle and the jet speed1.
Materiali in tehnologije / Materials and technology 48 (2014) 5, 725–734 725
UDK 621.96:691.73:620.179.11 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)725(2014)
Figure 1: Surfaces obtained with the jet flow and the quality zones2,3
Slika 1: Povr{ine, dobljene s curkom, in njihova kvaliteta 2,3
Characteristics of the cut surface: When examined in
order to determine the surface quality, the surfaces cut
with different methods are similar. Surface roughness is
defined with the waves on the surface and the size of the
wave is proportional to the jet diameter (Figure 1)2,3.
While the wave size depends on the jet diameter and
the penetration of the abrasive water jet, the surface
roughness is related to the micro-interaction between
each abrasive and the workpiece. The cutting quality de-
pends on the inner physical effects caused by the jet and
the external factors such as various cutting parameters,
nozzle vibration and job fragment. When a surface cut
with abrasive water is examined, three different sections
can be seen as shown in Figure 22,4.
1. In the upper corner of the cut surface there is a
small curve caused by the hitting articles departing from
the jet. This section is usually accepted as an ignorable
edge impact.
2. This is a smoother surface section located under
the first section. This section is formed by the particle
erosion caused by the abrasive particles hitting the sur-
face at a low impact angle. Experimental studies per-
formed recently have proven the fact that a 1.3 μm sur-
face-roughness quality can be obtained on this section.
3. The cutting capability is reduced as the kinetic en-
ergy of the abrasives decreases and the jet looses it regu-
larity. This is a transition section where the second cut-
ting mechanism prevails and the surface is formed by
faults due to parallel jet deviations. In this second cutting
mechanism, the impulse angle of the hitting particles
against the surface is bigger and defined as "the deforma-
tion erosion". The deformation abrasive mechanism is
realized by the particles hitting the surface at a bigger
angle. When the travelling speed of the jet is reduced,
the transition area between the second and third sections
is smaller4.
If a quality cutting process is required, the parameters
must be adjusted and the cutting process must be com-
pleted before entering the deformation abrasion section.
By adjusting the parameters, the flaking will also be
avoided. By selecting a sufficiently low lateral speed
level, a considerably smooth surface without any flaking
will be obtained on the first section. Smaller abrasive
particles and a bigger abrasive mass of the jet flow will
reduce the surface-roughness value. A particle with
bigger dimensions will consequently cause a larger cut
area and the surface will be rougher (it will have a larger
roughness value)3–5.
Increasing the abrasive mass of the jet or reducing the
jet speed will improve the quality by increasing the
number of the particles hitting against the surface being
cut. When greater cutting speeds are used in a rough
cutting operation, each of the three sections can be seen
on the surface. A deviation of the jet on the third section
and a formation of parallel lines appear as a function of
the parameters of lateral speed alterations, abrasive
feeding-flow rate, liquid pressure and nozzle geometry.
Abrasive substances form holes and pockets at the lower
parts, where they are accumulated and embedded during
the rough cutting operations. Such residual particles may
damage the nozzle during the operations. These negative
effects must be taken into consideration. When the
surface quality and energy of the particle are considered,
we find that as the cutting depth gets bigger the deviation
of the jet increases causing an increase in the energy of
the particle.6,7 Thus, a greater energy applied on the
surface show that the surface roughness and surface
waviness are more robust and there are more deviations
of the jet (Figure 3)1,4.
Comparison of the abrasive water jet with the alter-
native methods: In Figure 4, the inverse relationship
between the thickness and lateral feed rate is shown,
considering the surface quality of the cutting surface.
The AWJ method has the lowest lateral feed rate, while
the plasma method has the highest feed rate. An overall
comparison of the abrasive water jet and the alternative
cutting methods in Table 1 shows that the most efficient
cutting method is the cutting with AWJ, being inde-
pendent of the material thickness and its characteristics.
However, there are some disadvantages of this method.
The most important one is the dependency of the system
and the cutting parameters on several variables. Because
of this dependency, it is hard to provide a continuous
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726 Materiali in tehnologije / Materials and technology 48 (2014) 5, 725–734
Figure 3: Cut-face quality zones based on jet flow143
Slika 3: Podro~ja na povr{ini, rezani s curkom1,4
Figure 2: Surface sections cut by the abrasive water jet4
Slika 2: Podro~ja povr{ine reza pri abrazijskem vodnem curku4
surface quality on the cutting surface. An increasing
surface roughness is inevitable, as in the cases of laser,
plasma, underwater plasma and oxygen-flame cutting
methods4,8–12.
There are several studies that compare the AWJ
method with the other methods. The studies give differ-
ent results due to different materials used. The techni-
ques of AWJ and the other methods are compared by
Hashish2 as shown in Figures 5a and 5b. This compa-
rison is based on an evaluation of different processing
methods in terms of their power levels and typical
machining removal rates. There are various techniques
for cutting materials (Figure 6)2,9,13.
According to Hashish,4 when compared with the
traditional methods, AWJ forms a jet that is able to
perform a cutting process with a very low energy and an
intense energy distribution where most of the energy is
lost due to friction. Just as in the other unipolar, ductile
cutting operations, AWJ can be given directions perfectly
well with a low energy applied, and it can perform
cutting in all directions and can form considerably
narrow cuts. Particularly due to no thermal effects on the
cut materials, AWJ is more effective than the other
competitive methods. However, in spite of the many
advantages of the WJ and AWJ processing technologies,
there are still certain disadvantages2,14.
There are many studies comparing AWJ and the other
methods. When these are examined different results are
set forth depending on the material. Powell et al.10 per-
formed a study comparing the economical aspect of AWJ
and laser. In their analysis they discussed the technical
and commercial advantages and disadvantages of both
methods and focused on the relative productivity of both
processes. Ohlsson et al.11 studied the pressure, abrasive
flow and lateral-speed impacts on the steel cut with
AWJ, the grey-cast-iron cutting depth and surface pro-
perties. Zheng et al.14 made comparisons based on the
quality and process costs, aiming at helping the users
decide which methods would be more convenient for
various applications. They made their comparison by
using stainless steel with different thicknesses, soft steel
and aluminum12–14. In the studies by Hashish and
Ramulu,13 focusing on the mechanical properties of laser
and AWJ, they discussed the unique cutting abilities and
characteristics of both methods. The researchers drew
the attention of the users not only on the technical
performance of the methods but also on how they affect
the completed products; they evaluated the mechanical
impacts of both methods on the titanium-alloy
(Ti6Al4V) and steel (A286) materials12. As the optimum
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Materiali in tehnologije / Materials and technology 48 (2014) 5, 725–734 727
Figure 6: Comparison of the abrasive water jet with the other cutting
methods2,9,13
Slika 6: Primerjava rezanja z abrazivnim vodnim curkom z drugimi
metodami rezanja2,9,13
Figure 4: Comparison of the cutting abilities of different cutting
methods using single-orifice jet beams8
Slika 4: Primerjava zmogljivosti razli~nih metod rezanja pri uporabi
curka z eno {obo8
Figure 5: Comparison of the abrasive water jet with the other cutting
methods4,8
Slika 5: Primerjava rezanja z abrazivnim vodnim curkom z drugimi
metodami rezanja4,8
parameters have not been completely determined yet for
the vast majority of these methods, there are plenty of
other studies still being currently performed. The best
data to set forth the superiority of AWJ is probably the
figure given below. Furthermore, a graph is given indi-
cating the capability of the method with respect to
material thickness and a general comparison is given in
Tables 1 and 21,9,15–17.
Applications of various machining methods are
summarised in Tables 2 and 3. The machining characte-
ristics of different non-conventional processes can be
analyzed with respect to metal-removal rate, tolerance
maintained, surface obtained, depth of surface damage
and power required for machining. The physical para-
meters of the non-conventional machining processes
have direct impacts on the metal removal as well as on
the energy consumed for different processes. These
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Table 1: Overall comparison of abrasive water jet and the alternative cutting methods 1,9,16
Tabela 1: Primerjava abrazijskega vodnega curka z drugimi metodami rezanja1,9,16
Comparison of Disconnections by Water Jet and the Other Machining Methods
Comparison Factor AbrasiveWater Jet
Laser Cut-
ting
Plasma
Cutting
Underwa-
ter Plasma Wire EDM
Milling
Cutting Band Saw
Oxygen
Cutting
Material Thickness A C B B A B B A
Cutting Quality A A C B A B B C
Lateral Speed B A B B B B A B
Multi-Purpose Use A D B B B B B C
Heat Affected Zone (HAZ) A D D C C B B D
Sensitive Cutting A A B B A A C D
Secondary Process Requirement A B B B B B C C
Chip Formation B C C C A B D B
Production Flexibility A B C C B A C D
Overall Process Time B B D D B B A C
A: Excellent B: Good C: Acceptable D: Unacceptable
Table 2: Material applications of some machining methods1,9
Tabela 2: Uporabnost obdelovalnih metod glede na material1,9
Materials Applications
Process Aluminium Steel Super Alloys Titanium Refectories Plastics Ceramics Glass
Ultrasonic Machining C B C B A B A A
Abrasive Jet Machining B B A B A B A A
Electrochemical Machining B A A B B D D D
Chemical Machining A A B B C C C B
Electric Discharge Machining B A A A A D D D
Electron Beam Machining B B B B A B A B
Laser Beam Machining B B B B C B A B
Plasma Arc Machining A A A B C C D D
Abrasive Water Jet Machining A A A A A B A A
A: Good Application B: Fair C: Poor D: Not Applicable
Table 3: Process capabilities of non-conventional cutting processes9,12
Tabela 3: Zmogljivosti nekonvencionalnih postopkov rezanja9,12
Process Capability
Process Metal Removal Rate(mm/min)
Tolerance
(μm)
Surface
(μm) CAL
Depth of Surface
Damage (μm) Corner
Power
(W)
Ultrasonic Machining 300 75 0.2–0.5 25 0.025 2 400
Abrasive Jet Machining 0.8 50 0.5–1.25 2.5 0.100 250
Electrochemical Machining 0.15 15 0.1–2.5 50 0.025 100000
Chemical Machining 150 50 0.4–2.5 50 0.125 –
Electric Discharge Machining 800 15 0.2–1.25 125 0.025 2 700
Electron Beam Machining 16 25 0.4–2.5 250 250 150 (average),200 (peak)
Laser Beam Machining 0.1 25 0.4–1.25 125 250 2 (average)
Plasma Arc Machining 75000 125 Rough 500 – 50000
Abrasive Water Jet Machining 1.3 25 0.4–2.5 125 0.025 220
Conventional Milling of Steel 50000 50 0.4–5.0 25 0.050 3000
characteristics of different methods are given in Tables 4
and 59,17–19.
2 EXPERIMENTAL STUDIES
In this study the samples of (Figure 7) brass-353
(+) material 20 mm were cut with conventional (oxy-
gen flame, hydraulic saw and freeze) and eight uncon-
ventional methods (abrasive water jet, laser-plasma arc,
underwater plasma, wire erosion). The cutting edges
obtained with these methods were examined in terms of
their hardness and their effect on the microstructures. A
comparison was made between the initial microstruc-
tures and the microstructures of the materials after
cutting them with different methods; the effectiveness of
the methods was evaluated. Water-jet-cutting parameters
are shown in Table 6. Other cutting-process parameters
were selected according to the parameters recommended
by the lathe-manufacturing companies.
Chemical composition of the material: w/% (S 0.831,
Pb 2.21, Zn 36.37, P 0.216, Mn 0.0778, Fe 0.293, Si
0.0829, Al 0.442, Cu < 59.23, Ni 0.237)
The average hardness level was calculated by taking
the arithmetic average of the measured values at five dif-
ferent points at a given height on the surface. The value
of HV 30 was calculated with an INSTRON WOLPERT
TESTER hardness-measurement device. Additionally,
hardness was measured in intervals 1 mm from the edge
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Figure 7: Samples after cutting
Slika 7: Vzorci po rezanju
Table 4: Effects of different machining methods on equipment and tooling9
Tabela 4: Vpliv razli~nih metod obdelovanja na opremo in orodje9
Effects on Equipment and Tooling
Process Tool Wear Ratio Machining MediumContamination Safety Toxicity
Ultrasonic Machining 10 B A A
Abrasive Jet Machining – B B A
Electrochemical Machining 0 C B A
Chemical Machining 0 C B A
Electric Discharge Machining 6.6 B B B
Electron Beam Machining – B B A
Laser Beam Machining – A B A
Plasma Arc Machining – A A A
Abrasive Water Jet Machining – B B A
Tool Wear Ratio = Volume of work material removed / Volume of tool electrode removed
A: No Problem B: Normal Problem C: Critical Problem
Table 5: Economic performance of different machining methods9
Tabela 5: Ekonomi~nost posameznih metod rezanja9
Process Economy
Process Capital Invest-ment
Tooling and Fix-
tures
Power Require-
ment Efficiency
Tool Consump-
tion
Ultrasonic Machining B B B D C
Abrasive Jet Machining A B B D B
Electrochemical Machining E C C B A
Chemical Machining C B D* C A
Electric Discharge Machining C D B D D
Electron Beam Machining D B B E A
Laser Beam Machining C B A E A
Plasma Arc Machining A B A A A
Abrasive Water Jet Machining B B B C C
Conventional Milling of Steel B B B A B
A: Very Low Cost B: Low C: Medium D: High E: Very High
*Indicates cost of chemicals.
of the material towards the inner part along a linear line,
so that the hardness changes depending on the heat
distribution were observed. The microstructures of the
main material and the cut edges were viewed with a
PANASONIC WV-CP410 Model, Type N334, light
microscope, with a magnification of 280-times. Alumina
and diamond paste were used to examine the microstruc-
ture of the material in the polishing operation followed
by the etching process when dipped in the mixture of
2 mL of HNO3 and 98 mL of methane alcohol for 20 s.
Examination of different cutting methods in terms of
the structural variations created on the materials: In or-
der to perform metallographic examinations and find
structural deterioration on the cut section of the material,
a microstructure photo of the section resistant to the
cutting process was taken as shown in Figure 8. For an
accurate assessment, plenty of photos were taken from
every cutting edge, and the deformations due to the
cutting method formed on the material structure as well
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Table 6: Cutting systems and cutting parameters
Tabela 6: Sistemi rezanja in parametri rezanja
Abrasive Water Jet Cutting
Water consumption  3.5 L/min Pump piston diameter 20 mm
System temperature of water 48 °C Inlet pressure of water into thepressure booster 6 bar
Working pressure of the booster 200 bar Inlet diameter of water into thenozzle 0.25 mm
Outlet pressure of water from the pressure
booster 20 bar
Abrasive nozzle inlet diameter into
the nozzle 0.75 mm
Water flow rate 3 L/min Stand-off distance 4 mm
Outlet velocity of water from the nozzle 800 m/s Water pressure at the instance ofdischarge 400 MPa
Temperature at the instance of cutting  55 °C Jet angle at the nozzle 90°
Current consumption during work 380 V Energy consumption 58 kW h
Amount of abrasive consumed 250 g/min Material used in the nozzle orifice Sapphire
Abrasive used GMA Garnet Chemical composition Fe2O3Al2 (SIO4)3
Abrasive hardness (Mohs) 7.5–8 Abrasive particle size 300 μm
Abrasive water outlet diameter from the nozzle 0.75 mm Nozzle length 76.2 mm
Slurry content 18 % Mixing tube length 88.9 mm
Mixing tube diameter 1.27 mm Nozzle orifice life 40–50 h
Laser Beam Cutting Plasma Beam and Water Shield Plasma Cutting
Cutting rate (Lateral feed rate) 20 m/min Cutting rate (Lateral feed rate) 20 m/min
Position rate 140 m/min Plate positioning By Laser
Laser power 1550 W Current for maximum cutting 760 A
Main power supply GW 0–100 % Nozzle pressure 12 bar
Pulse type Mega pulse Operating pressure 24 bar
Pulse change frequency 2000 Hz Operating frequency 50 Hz
Pulse time NP(T) 1500 ìs Cooling capacity
16747 kJ/h
SP(t) 120 ìs
Mod type Sürekli mod
(CW)
Nominal voltage 400 V
Focus distance 7.5 mm Average sound level (A) 68 dB
Cutting gas Nitrogen Cutting gas Oxygen + Nitrogen
Cutting gas pressure 1.2 bar Maximum cutting thickness 35 mm
Cooling temperature TA = 25 °C Cutting capacity 4000 mm × 7000 mm
Oxygen Flame Cutting Wire Electrical Discharge Cutting
Cutting rate (Lateral feed
rate) 20 m/min
Processing condition C521
Current for maximum cutting 760 A Feed rate 3 m/min
Nozzle pressure 10 bar Processing conditions and parameters
Operating pressure 20 bar ON OFF IP HP MA SV V SF C
Operating frequency 50 Hz 006 15 17 2 15 0.3 0.3 005 0
Cooling capacity 16747 kJ/h Voltage 32 V
Receiver tank capacity 30 l Current 5.6 A
Nominal voltage 400 V Wire tension Level 8
Average sound level (A) 68 dB Wire feed rate 7 m/min
Cutting gas Oxygen+Propane Control system Fine APT
Parameters for each cutting methods are selected in accordance with the machine manufacturers’ recommendations.
as their changes were evaluated at the end of examining
these photos.
Stripe (hydraulic) saw cutting: A considerably rough
surface profile was obtained on the cut section. This cut
profile was similar to the gear shape and it had a rather
rough surface. Again, in the distance of approximately
25 μm a hard rough surface was observed due to the ef-
fect of cool deformation. No remarkable morphological
change is seen, but there is a structural change caused by
the deformation (Figure 9a).
Cutting with a milling cutter: A very flat profile was
gained on the cut section. In the distance of nearly 10
μm, there is a remarkable, strong surface affected by the
impact of cool deformation. Also, no strong structural
change due to heat is noticed, but there is a change due
to deformation (Figure 9b).
Cutting with underwater plasma: In the distances of
approximately 75 μm the grain size of  and  phases
got smaller and thinner but on the remaining section the
grain size remained same. Also, structural deformations
exist on the cut section due to excessive warming and
fast cooling in the water (Figure 9c).
Laser cutting: On the cut section a rough cut profile
is visible and due to a high temperature and cooling in
air, the geometries of  and  phases turned into acicular
forms. At the same time, the structure of the cut section
was entirely deformed and  grains became different
from their original forms. On the cut section and around
it, a new rigid and fragile form emerged (Figure 9d).
Cutting with plasma: Owing to the heat effect, in the
distances of approximately 75 μm, the grain size of 
and  phases got smaller and over the main metal section
the size of these grains got infinitesimally small as well.
Moreover, due to excessive heat and fast cooling, the
grains forming the structure got thinner. This trend con-
tinues towards the inner sections. A new hard and fragile
structure was formed (Figure 9e).
Cutting with abrasive water jet: A very flat cut sur-
face was obtained and in the distance of approximately
10 ìm, a layer affected by cool deformation was ob-
served. Apart from that, no structural alteration was ob-
served on the cut section (Figure 9f).
Cutting with wire erosion: In the distance of approxi-
mately 20 μm the particle size of  and  phases got
smaller and over the remaining part the particle size re-
mained the same. Also, on the cut section, the particles
forming the structure got thinner, more rigid and fragile
due to excessive heat and rapid cooling (Figure 9g).
Cutting with oxygen flame: Over the cut section, the
structure was entirely deformed and there was a new
form, different from the original one. Due to an exces-
sive heat input and rapid cooling in air, the geometries of
 and  phases changed and  particles, apparently
acicular, were also formed around the cutting section
(Figure 9h).
In Figure 10, the surface-roughness values obtained
by cutting thick brass-353 20 mm with different methods
are compared. If this graph is carefully analyzed, it is
clear that the roughest surface is obtained by cutting the
material with the oxygen-flame method and the smooth-
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Materiali in tehnologije / Materials and technology 48 (2014) 5, 725–734 731
Figure 8: Microstructure of brass-353 ()
Slika 8: Mikrostruktura medenine 353 ()
Figure 9: a) Stripe-saw cutting, b) milling cutting, c) underwater
plasma cutting, d) laser cutting, e) plasma cutting, f) abrasive water jet
cutting, g) wire-erosion cutting, h) oxygen-flame cutting
Slika 9: a) Rezanje s tra~no `ago, b) rezanje z rezkanjem, c) podvod-
no rezanje s plazmo, d) rezanje z laserjem, e) rezanje s plazmo, f)
abrazijsko rezanje z vodnim curkom, g) rezanje z `i~no erozijo, h) pla-
mensko rezanje s kisikom
est surface is obtained by cutting it with the wire-erosion
method.
The obtained outcomes of the study were evaluated
using the unprocessed surface-microstructure photo-
graphs of the material shown in Figure 11 and the sur-
face microstructures of different methods shown in
Figure 12. With the conventional cutting methods (in
this study they include the milling cutter and the band
saw) nearly the entire energy used for the machining was
liberated as heat and a very small percentage of the
energy turned into lost energy in the form of an elastic
loss14,15,19.
If the heat liberated in this way is not controlled, it
will lead to a change in the metallurgical properties of
the material. When the temperature is higher than the
recrystallization heat of the material, it will lead to sig-
nificant changes in the metallurgical properties of the
material. The cooling conditions applied during
machining will also affect the metallurgical forms of the
material. Transformation of the energy into heat and the
cooling conditions can be interpreted as the underlying
reasons of the main metallurgical and mechanical chan-
ges such as the hardness of the material. The fundamen-
tal principle of the oxygen-flame cutting operation relies
on rising the temperature of the material to the melting
point. Rising the heat to the melting temperature and the
successive cooling conditions will lead to significant
changes in the mechanical and metallurgical formation
of the material. This study also gave the expected result,
according to which both metallurgical and hardness
properties revealed the most significant changes to the
material cut with this method.
The causes for the metallurgical changes and hard-
ness variations in the materials are based on the frame-
works of the methods applied. The laser, plasma and
wire-erosion methods are based on the principle of
cutting the material at the melting heat level. Different
energy inputs and cooling conditions are the main causes
for different metallurgical and hardness formations.
Among the traditional methods, the hardness values
obtained with the underwater-plasma (focusing) and
wire-erosion methods were a little better than those
obtained with the laser and plasma methods, because
they were implemented in a preserving liquid and, thus,
the temperature level was controlled. If a comparison is
to be made between the executed cutting methods in
terms of metallurgical properties and hardness factors on
the basis of the original material structure and the hard-
ness alteration, the best outcome is obtained for the AWJ
cutting method. The hardness values for the surfaces cut
with AWJ are fairly close to the original hardness ratios
(for all the materials). This can be explained in terms of
abrasion mechanisms. When cutting with AWJ, the heat
variation remains very low (around t = 75 °C)1,4,7. This
shows that no section (HAZ) is affected by the heat
factor when using the AWJ cutting method. Taking this
feature into account, it is clear that the AWJ cutting
method is outstanding, not causing any form of metal-
lurgical and mechanical alteration of the original mate-
rial.
For the bras-353 materials used in this study, the
hardness differences caused by different methods on the
cut surfaces are shown in Figure 11 and the impact rates
of these effects are shown in Table 7. Following the
AWJ cutting method, the second lowest change in the
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732 Materiali in tehnologije / Materials and technology 48 (2014) 5, 725–734
Figure 11: Comparison of the hardness values for the brass-353 (+)
samples cut with different methods in comparison with the original
hardness of the material core
Slika 11: Primerjava trdot medenine 353 (+), odrezane z razli~nimi
metodami, v primerjavi s trdoto jedra materiala
Table 7: Hardness variations for brass-353 (+) cut with different
methods
Tabela 7: Spreminjanje trdote medenine 353 (+)
Cutting Method
Brass-353
Hardness
(HV30)
Change
(%)
Base material 115.17 -
Cutting by Abrasive Water Jet 116.50 1.15
Cutting by Milling Cutter 118.17 2.60
Stripe (Hydraulic) Saw Disconnection 118.00 2.46
Cutting by Oxygen Flame 128.50 11.57
Cutting by Laser 122.67 6.51
Cutting by Plasma 125.50 8.97
Cutting by Underwater Plasma 119.33 3.61
Wire EDM Cutting 118.50 2.89
Figure 10: Comparison of the roughness values of cut faces obtained
by cutting brass-353 with different methods
Slika 10: Primerjava hrapavosti povr{ine reza pri rezanju medenine
353 z razli~nimi metodami
hardness is observed for the conventional methods such
as the milling cutter and the band saw. This finding may
be attributed to the fact that for the classical methods the
cutting parameters are selected so as to avoid excessive
recrystallization heat levels.
The depth of the section exposed to the heat also
changes depending on the properties of the cutting
method. Due to the changes in the metallurgical struc-
tures caused by the method, the measurement of the
hardness, in the distance 1 mm starting from the cut
surface towards the inner part, provides the information
on the width of the section affected by the heat factor.
The results of these measurements for brass-353 are
shown in Figure 12. The most outstanding result ob-
served from the graphs is the fact that there is a linear
slope for the AWJ cutting method and, thus, no section
on the brass material is affected by heat. The AWJ
cutting method appears to be a process causing almost
no change in the material hardness and metallurgical
properties. On the other hand, the oxygen-flame cutting
causes the highest level of change to the metallurgical
and hardness properties. With this method, the hardness
varies significantly from the surface to the core, and the
whole material is affected by the heat factor. With the
laser and plasma-cutting methods known as the biggest
rivals to the AWJ cutting method, the hardness changes
from the surface to the core, indicating that a large
percentage of the surface of the material is affected by
the heat factor. With respect to the metallurgical
properties of the material, these methods cannot compete
with AWJ.
When all the methods are taken into consideration,
the hardness of brass changes constantly. This tendency,
which is higher up to some point in steel materials, is re-
duced after a certain point18,19. This circumstance may be
explained as a dependency on the heat conductivity of
the material. For brass-353, the heat conductivity is
higher than that of steel and, thus, the section affected by
the heat factor is larger.
3 CONCLUSION
When the effects of different cutting methods on the
metallurgical properties of the surface are taken into con-
sideration, the AWJ cutting prevails outstandingly over
the other cutting methods.
While different cooling and heat impacts caused by
different cutting methods have important effects on the
metallurgical properties of the material, in the AWJ cut-
ting method, no section is affected by the heat as the
temperature on the surface (HAZ) is not very high and
there is no destruction of the original properties of the
material. This finding shows that the mechanical proper-
ties of the material will remain unchanged as well.
Depending on the changes in the microstructure
properties of the material, the section affected by the
heat factor and the width of this section are subjected to
structural change because of the high heat and cooling of
some methods. Depending on the features of the cutting
methods, some methods cause a rough particle formation
and others cause a thin particle formation, due to instant
cooling. Again, due to the effects of the methods, gas
holes in the structure and microcracks are likely to
emerge. In the AWJ cutting method, a high heat and in-
stant cooling are the fundamental reasons for the
microstructures not being destroyed.
When evaluating the eight different methods
examined in this study on the basis of the changes in the
microstructure properties of the section affected by the
heat factor, it is clear that the least effective method is
the oxygen-flame cutting and the most effective one is
the AWJ cutting. Among the applied methods, the oxy-
gen-flame cutting is viewed as the poorest method
because of the variation in the hardness of the material it
creates.
Depending on the effects of different methods on the
metallurgical forms of the material, the mechanical prop-
erties of the material also change. In the experimental
studies, the hardness values of the material, after using
different methods, are different from the original values.
This finding proves that the other cutting methods
change the mechanical properties of the materials.
All the cutting methods tested in this study change
the hardness of the material. This variation is changeable
depending on the heat, temperature and cooling condi-
tions occurring during the cutting operation.
When comparing different cutting methods with re-
spect to the metallurgical properties and the hardness of
the material, the best method is the AWJ cutting. This
finding proves that during the AWJ cutting no section is
affected by the heat factor (HAZ).
When the hardness changes caused by the heat factor
are examined from the surface to the centre of the mate-
rial cut with different methods, the AWJ cutting stands
out as the most effective cutting method because with
AWJ no section is affected by the heat factor and the cut-
A. AKKURT: EXPERIMENTAL INVESTIGATION OF THE SURFACE PROPERTIES OBTAINED BY CUTTING BRASS-353 ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 725–734 733
Figure 12: Hardness variations for brass-353(+) from the cutting
edge to the center due to various cutting methods
Slika 12: Spreminjanje trdote medenine 353 (+) od roba rezanja
proti sredini pri razli~nih metodah rezanja
ting operation does not cause any metallurgical and me-
chanical changes to the material.
In the laser and plasma methods, considered as the
most important alternatives to the AWJ cutting, the
changes in the hardness from the surface to the center of
the material show that, with these methods, the section
affected by the heat is much larger than in the case of
AWJ.
When compared with the other methods, AWJ is an
effective and contemporary alternative cutting method in
terms of the surface properties of the materials pro-
cessed.
4 REFERENCES
1 A. Akkurt, Surface Properties of Various Materials in Abrasive Water
Jet, Comparative Examination of Hardness and Micro Structure
Changes in Different Methods, Doctoral Thesis, G. U. Science
Department Institution, Ankara, 2002 (In Turkish)
2 M. Hashish, Application of Abrasive-Waterjets to Metal Cutting,
Proceedings of The Nontraditional Machining Conference, ASM,
Cincinnati, OH, 1985, 1–11
3 N. S. Guo, H. Louis, G. Meier, A Surface structure and kerf geo-
metry in abrasive water jet cutting: formation and optimization, Proc.
7th American Water Jet Conf., 1 (1993), 1–25
4 M. Hashish, On the Modeling of Surface Waviness Produced by
Abrasive-Waterjets, Proceedings of the 11th International Con-
ference on Jet Cutting Technology, St. Andrews, Scotland, 1992,
17–34
5 N. S. Guo, H. Louis, G. Meier, Surface structure and kerf geometry
in abrasive water jet cutting: formation and optimization, Pro-
ceedings of the 7th American Waterjet Conf., Seattle WA, USA,
1993, 1–26
6 A.W. Momber, R. Kovacevic, Principles of AWJ Machining, Sprin-
ger-Verlag, London Limited 1988, 284–324
7 A. W. Momber, H. Kwak, R. Kovacevic, An alternative method for
the evalution of the abrasive water jet cutting process in gray cast
iron, J. of Mater. Process. Technol., (1997), 65–72
8 http://www.aliko.fi/english/watercut.htm
9 M. K. Singh, Unconventional Manufacturing Process, http://www.
newagepublishers.com/samplechapter/001566.pdf
10 J. Powell, L. Ohlsson, E. M. Olofsson, An Economic Comparison of
Laser and Abrasive Water Jet Cutting, Lulea University of Technol-
ogy, Division of Materials Processing, Sweden, 1995
11 L. Ohlsson, J. Poweell, A. Ivarson, C. Magnusson, Comparison bet-
ween abrasive water jet cutting and laser cutting, J. Laser Appl., 3
(1991) 3, 46–50
12 S. A. Brown, J. E. Lemons (Eds.), Medical Applications of Titanium
and its alloys, ASTM STP 1272, ASTM, Coshohochen, PA 1996
13 M. Ramulu, M. Hashish (Eds.), Machining Characteristics of
Advanced Engineering Materials, MD-Vol.16, ASME Publ, New
York NY 1989
14 H. Y. Zheng, Z. Z. Han, Z. D. Chen, W. L. Chen, S. Yeo, Quality and
Cost Comparisons between Laser and Waterjet Cutting, Journal of
Materials Processing Technology, 62 (1996) 4, 294–298
15 Y. W. Seo, M. Ramulu, D. Kim, Machinability of Titanium Alloy
(Ti-6Al-4V) by Abrasive Waterjets, Proc. Instn Mech. Engrs, Part B:
Journal of Engineering Manufacture, 217 (2003), 1709–1721
16 A. Akkurt, M. K. Külekçi, U. ªeker, F. Ercan, Effect of feed rate on
surface roughness in abrasive waterjet cutting applications, Journal
of Materials Processing Technology, 147 (2004), 389–396
17 J. A. McGeough, Advanced Methods of Machining, Chapman and
Hall Ltd, London UK 1988
18 A. Akkurt, Surface properties of the cut face obtained by different
cutting methods from AISI 304 stainless steel materials, Indian
Journal of Engineering and Materials Sciences, 16 (2009), 373–384
19 D. Arola, M. L. McCain, M. Ramulu, S. Kunaporn, Waterjet and
abrasive waterjet surface treatment of titanium: a comparison of
surface texture and residual stress, Wear, 249 (2001) 10–11, 943–950
A. AKKURT: EXPERIMENTAL INVESTIGATION OF THE SURFACE PROPERTIES OBTAINED BY CUTTING BRASS-353 ...
734 Materiali in tehnologije / Materials and technology 48 (2014) 5, 725–734
U. ARTI^EK et al.: INFLUENCE OF THE WORKING TECHNOLOGY ON THE DEVELOPMENT ...
INFLUENCE OF THE WORKING TECHNOLOGY ON
THE DEVELOPMENT OF ALLOYS H13-w(Cu) 87.5 %
VPLIV TEHNOLOGIJE IZDELAVE NA RAZVOJ ZLITINE
H13-w(Cu) 87,5 %
Uro{ Arti~ek1, Marko Bojinovi~2, Mihael Brun~ko3, Ivan An`el3
1EMO – Orodjarna, d. o. o., Be`igrajska cesta 10, 3000 Celje, Slovenia
2TIC-LENS, d. o. o., Be`igrajska cesta 10, 3000 Celje, Slovenia
3Faculty of Mechanical Engineering, University of Maribor, Smetanova 17, 2000 Maribor, Slovenia
uros.articek@gmail.com
Prejem rokopisa – received: 2013-11-04; sprejem za objavo – accepted for publication: 2013-11-22
Most dies in the casting industry for injection moulding are machined from the premium-grade H13 tool steel. They provide
excellent performance in terms of mechanical properties and service life; however, these dies are characterised by a relatively
low thermal conductivity. The tool-and-die industry is interested in depositing a material of a high thermal conductivity onto
steel in order to improve the thermal management and productivity. We have explored the possibility of using copper with a new
technology. In this study, the microstructure evolution and mechanical properties are discussed using the Laser Engineered Net
ShapingTM (LENSTM) technology. For a better understanding of the solidification, the microstructure of a LENS sample was
compared with the microstructure of a reference alloy produced with the ingot-casting technology having the same chemical
composition of H13-w(Cu) 87.5 %. We carried out light microscopy, scanning electron microscopy, an EDS microchemical
analysis, the tensile test and microhardness testing. The results show a successful fabrication of LENS samples; their
microstructure is more homogeneous compared to the castings; they show better mechanical properties and represent a good
potential for further development and use.
Keywords: LENS, casting, microstructure evolution, mechanical properties
V industriji tla~nega litja se za izdelavo matric orodij pogosto uporablja visokokakovostno orodno jeklo H13, ki ima odli~ne
mehanske lastnosti in dolgo trajnostno dobo, vendar je zanj zna~ilna relativno nizka toplotna prevodnost. Orodjarska industrija
`eli dodati jeklu material z visoko toplotno prevodnostjo za dosego bolj{e porazdelitve toplote in ve~je produktivnosti. Z novo
tehnologijo, imenovano Laser Engineered Net ShapingTM (LENSTM), smo raziskali mo`nost uporabe bakra, spremljali razvoj
mikrostrukture ter ugotovili mehanske lastnosti. Za bolj{e razumevanje strjevanja smo primerjali mikrostrukturo vzorca LENS z
mikrostrukturo referen~ne zlitine z enako kemijsko sestavo H13-w(Cu) 87,5 %, izdelano s tehnologijo litja. Karakterizacija
zlitin je potekala s svetlobno mikroskopijo, vrsti~no elektronsko mikroskopijo, mikrokemi~no EDS-analizo, z nateznim
preizkusom in merjenjem mikrotrdote. Rezultati ka`ejo uspe{no izdelavo vzorcev LENS, katerih mikrostruktura je bolj
homogena in ima bolj{e mehanske lastnosti v primerjavi z odlitki. Tehnologija LENS je dober potencial za nadaljnji razvoj in
uporabo v praksi.
Klju~ne besede: LENS, litje, razvoj mikrostrukture, mehanske lastnosti
1 INTRODUCTION
In the injection-moulding industry, new materials and
technologies are required for mould dies in order to
optimize the production and keep the costs as low as
possible. Despite their excellent mechanical properties,
tool steels that are nowadays used as the materials for
moulds limit the productivity due to their low thermal
conductivity. To solve this problem, designers have been
focusing on how to design the tool geometry and con-
struction to achieve higher cooling rates. Complex cool-
ing channels are being designed to enable the cooling
liquid to extract the heat from a mould. Ejector pins,
slides and air-stream gates are used to eject a part from a
mould cavity so the space left in it is small. An alter-
native way to solve this problem is the use of copper-
beryllium inserts1, which have a higher thermal conducti-
vity. However, they can leave marks on the part and are
not environmentally friendly.
Therefore, new technologies and materials, like a
Cu-deposition on tool steel, or functionally graded mate-
rials (FGM) with a combination of high strength, wear
resistance and thermal conductivity are explored as
potential candidates for a more efficient injection-mould-
ing tool. The use of a direct-metal-deposition (DMD)
fabrication process makes it possible to create an opti-
mum configuration of fin trees and cooling channels as
well as to use FGM without being concerned with the
limitations of the traditional manufacturing method2,3.
The laser-engineered-net-shapingTM (LENS) process is
seen as one of the most promising DMD technologies for
the production of such materials4–6. LENS is a relatively
new technology capable of rapidly producing complex,
fully dense parts directly from a computer-aided design
(CAD). A high-power Nd-YAG laser is used to heat and
melt a metal powder, thereby producing a melt pool on a
substrate attached to an X-Y table. The metal powder
from coaxial powder-feed nozzles is injected into the
melt pool as the table is moved along pre-designed 2-D
tool paths generated using the sliced CAD models. The
additions of multiple layers produce a 3-D net or a near-
net shape. The fabrication takes place under a controlled,
Materiali in tehnologije / Materials and technology 48 (2014) 5, 735–742 735
UDK 621.74.043:669.018.258 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)735(2014)
inert atmosphere of argon. Some of the important pro-
cess parameters are laser power, powder flow rate, layer
thickness, hatch width, deposition speed and oxygen
level.
The iron-copper (Fe-Cu) alloying system7 is one of
the most suitable systems to produce an efficient FGM
mould material. The thermal conductivity of copper is
approximately 13 times higher than that of the H13 tool
steel at the operating temperatures between 220–600 °C.
Unfortunately, a large solidification temperature range
and a high amount of the Cu-rich terminal liquid over a
wide range of Cu concentration promote solidification
cracking in Fe-Cu alloys8,9. Additionally, a Fe-Cu phase
diagram contains two peritectics and a nearly flat
liquidus; they exhibit a high tendency for non-equili-
brium solidification that probably has a significant influ-
ence on the susceptibility to cracking7,8. At the same
time, LENS is considered to be a technology that in-
volves a high velocity of solidification with the possibi-
lities of specific reactions, new phases and a metastable
microstructure formation. While the properties of mate-
rials mainly depend on the microstructures, it is import-
ant to know the microstructure development during the
LENS process as well as the final microstructure of the
layers, depending on the thermal influence of the addi-
tional layers.
Our main objective is to explore the microstructure
development with the LENS technology. As it is evident
from the previous research8, some chemical composi-
tions are more susceptible to the formation of cracks
when using DMD in a Fe-Cu system. Based on the
current data, it is not yet possible to conclude how the
LENS technology will influence the formation of cracks.
In the first part of our research, the composition of
H13-w(Cu) 87.5 % was selected. Despite the fact that
this composition belongs to the crack-free composition
range, the question of how the LENS technology influ-
ences the microstructure evolution is still open. Namely,
the thermal impact of the solidifying layers on the
microstructure of the previously solidified layers and the
irregularities of the previous layers is not yet known. In
our research, the resulting microstructures and mecha-
nical properties of the samples produced in this way
were compared with the conventional-casting samples
with the same composition, which enabled us to evaluate
the effect and the rationality of the technology.
2 EXPERIMENTAL WORK
The conventional solidification of the alloying system
consisting of the H13 tool steel and the mass fraction of
Cu w = 87.5 % was studied using the mould casting. For
the preparation of the alloy, oxygen-free high-conduc-
tivity copper (w (Cu) = 99.99 %) and the H13 tool steel
with the chemical composition presented in Table 1
were used. The material system was vacuum-induction
heated at 10–2 mbar. Before melting, the chamber was
backfilled with the high-purity Ar gas up to 1150 mbar.
The melt was homogenized at 1450 °C and then cast at
1400 °C into, cylindrically shaped, grey-cast-iron
moulds 50 mm. The inner walls of the moulds were
protected with a thin layer of ZrO2. Before casting, the
moulds were preheated to 400 °C to decrease the cooling
rate and lower the melt undercooling prior to the primary
solid nucleation.
For the layered manufacturing experiments a LENS
850-R machine made by Optomec Inc. with a high-
power Nd:YAG laser with a capacity of 1000 W was
used. The machine consists of a dual-powder feeder
system that allows a simultaneous delivery of two
different material mixtures. Several cylindrically shaped
samples (D = 10 mm, L = 100 mm) were successfully
produced using the following parameters: the laser
power of 530 W, the traverse speed of 5.3 mm/s, the
layer thickness of 0.35 mm, the hatch spacing of 0.46
mm, the hatch angle of 60° and the powder-flow rate of
2.75 g/min. In our experiments, we used the powders
produced using gas atomisation with the particle sizes
ranging from 45 μm to 160 μm (Figure 1). The greater
U. ARTI^EK et al.: INFLUENCE OF THE WORKING TECHNOLOGY ON THE DEVELOPMENT ...
736 Materiali in tehnologije / Materials and technology 48 (2014) 5, 735–742
Figure 1: SEM images of the powders: a) Cu, b) tool steel H13
Slika 1: SEM-posnetka prahov: a) Cu, b) orodno jeklo H13
Table 1: Chemical composition of tool steel H13 in mass fractions, w/%
Tabela 1: Kemijska sestava orodnega jekla H13 v masnih dele`ih, w/%
Element C Si Mn Cr Ni Mo V Fe
Composition (w/%) 0.32 - 0.45 0.8 - 1.2 0.2 - 0.5 4.75 - 5.5 0.3 max 1.1 - 1.75 0.8 - 1.2 bal.
parts of the powders of both materials were spherical,
providing the solution to the porosity problem10,11. The
powders were delivered by the argon carrier gas (2 L/min)
to the focus of the laser beam. The oxygen level during
all the experiments was maintained below 10 · 10–6.
The influence of the microstructure on the mecha-
nical properties was evaluated using uniaxial tensile
testing and microhardness measurements. The static
tensile tests were performed on a Zwick/Roell ZO 10
tensile-testing machine with a load cell capacity of 10
kN at a constant position and a controlled speed of the
crosshead of v = 1.5 mm/min at ambient temperature.
The shapes and dimensions of the tensile-test samples
complied with the SIST EN 10002-1 standard. The
mechanical properties of several testing samples cut out
from the cylinder-shaped casting produced with the
conventional casting technology were compared with the
tensile bars obtained with the LENS technology,
machined in the longitudinal orientation so that the axis
of the tensile bars was parallel to the cladding direction.
The hardness measurements were carried out according
to the 6507-1:1998 standard by means of the Vickers test
on a Zwick 3212 microhardness-measurement device.
A microstructural characterisation of the conven-
tionally cast and LENS materials was carried out with
light microscopy, LM (Nikon Epiphot 300) and scanning
electron microscopy, SEM (FEG Sirion 400 NC) as well
as the energy dispersive X-ray analysis, EDX (INCA
6650). The samples for light and electron microscopies
were grinded, polished and etched according to the
standard metallographic procedures. The etching in a
solution consisting of 5 g FeCl3, 10 mL HCl and 100 mL
ethanol for about 20 s was used to reveal the
microstructure.
3 RESULTS AND DISCUSSION
The microstructure evolution during solidification
depends on the alloy characteristics, its chemical compo-
sition and it is primarily a function of the solidification
condition. For a better understanding of the solidification
process under the LENS conditions, the microstructure
obtained with conventional solidification in the mould
casting of an alloying system with the same composition
was studied first.
3.1 Microstructure of conventional-casting samples
The typical microstructure of the conventionally cast
material composed of the H13 tool steel and w(Cu) =
87.5 % is shown in Figure 2. The microstructure con-
sists of the Fe-rich dendritic-like primary phase (P1) and
Cu-rich matrix (M1). However, the size, morphology and
distribution of the primary-phase particles are very
different throughout the volume of the casting, indicating
a high inhomogeneity of the microstructure.
A detailed microstructural analysis of the samples at
a higher magnification revealed very fine dendrites (P2)
in the Cu-rich matrix and the Cu-rich zone (M2) around
the Fe-rich particles (Figure 3). An elemental EDX
analysis performed within the FE SEM indicates that the
Fe-rich primary phase and the surrounding zone contain
the alloying elements of the H13 tool steel, while the
Cu-rich matrix and fine dendrites consist only of copper
and iron (Table 2). Because of the inaccuracy of the
EDX method used for a quantitative analysis of light
U. ARTI^EK et al.: INFLUENCE OF THE WORKING TECHNOLOGY ON THE DEVELOPMENT ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 735–742 737
Figure 2: Microstructure of the as-cast sample (SEM back-scattered
electron image – BEI)
Slika 2: Mikrostruktura litega vzorca (SEM-posnetek povratno sipa-
nih elektronov – BEI)
Figure 3: SEM BEI of the: a) as-cast sample, b) showing the Cu-rich
zone around the Fe-rich particle and c) fine dendrites in the Cu-rich
matrix
Slika 3: SEM BEI: a) litega vzorca, b) podro~je, bogato z bakrom,
okrog delca, bogatega z `elezom, in c) fini dendriti v matrici, bogati z
bakrom (c)
Table 2: Chemical compositions of different microstructural regions
in mass fractions, w/%
Tabela 2: Kemijska sestava razli~nih podro~ij mikrostrukture v
masnih dele`ih, w/%
Microstructural
region
Chemical composition (w/%)
Fe Cu Cr Mo Si V Mn
P1 78.9 14.5 3.8 1.2 0.6 0.9 0.1
M1 3.6 96.4
M2 4.5 95.0 0.3 0.1 0.1
elements, the detected carbon concentration was not
taken into consideration in these results. Also, in the
microchemical analysis of fine dendrites, the size of the
interaction volume was too large for a correct quantita-
tive evaluation of the elements. Therefore, a line analysis
is presented in Figure 4 to show the detected chemical
elements in the region of fine dendrites.
It is well known that the alloys from the ternary
Cu-Fe-Cr system12,13 indicate a high tendency for non-
equilibrium solidification with metastable transforma-
tions. The phase diagrams of the constituent binary
Cu-Fe and Cu-Cr systems have flat parts of the liquidus
lines, and a metastable liquid-phase separation has been
established for these systems13,14. The Cu-Fe system
displays a large solidification-temperature range and a
peritectic reaction at both ends of the phase diagram.
Under the near-equilibrium condition, the solidification
of the hypo-peritectic composition with w(Cu) = 87.5 %
starts with the -Fe dendrite nucleation. A further cool-
ing leads to the growth of the primary phase, and the
solidification ends with a peritectic reaction, where the
rest of the remaining liquid reacts with an equivalent part
of the primary solid, i. e., -Fe + L  	-Cu. On the other
hand, if the melt is undercooled below the metastable
miscibility gap, the metastable liquid-phase separation
takes place, i. e., L  L1 (Fe-rich) + L2 (Cu-rich). In this
case, the L1 phase solidifies as the leading phase under
the non-equilibrium conditions and the solidification of
the L2 phase proceeds under the near-equilibrium con-
dition.
According to the Cu-Fe binary-phase diagram with
metastable miscibility lines, the minimum undercooling
which is required for the liquid-phase separation depends
on the chemical composition15. For the alloy with w(Cu)
= 87.5 %, the estimated critical value of T based on the
metastable miscibility line in the phase diagram is about
70 K. An addition of Cr increases the critical tempera-
ture of the miscibility gap of the Cu-Fe binary system15,16
and decreases the necessary undercooling for the meta-
stable liquid separation.
The results of our microstructural analysis indicate
that the obtained undercooling in the mould casting of
the material composed of the H13 tool steel and w(Cu) =
87.5 % exceeded the critical value for the melt separation
into two liquids: Fe-rich and Cu-rich. After the separa-
tion, both liquid phases were undercooled. In accordance
with different copper concentrations, the Fe-rich liquid
was more undercooled, having a larger driving force for
nucleation. Consequently, the solidification started with
the Fe-rich primary-phase nucleation in the Fe-rich
liquid. The lower undercooling as well as the recale-
scence event during the Fe-rich liquid solidification
enabled the Cu-rich liquid to solidify at a much smaller
deviation from equilibrium. The composition and frac-
tion of each liquid changed as the sample was being
continuously cooled and probably followed the meta-
stable miscibility-gap phase boundary as the time was
allowed for the transfer of the atoms between the two
liquids. The solidification of both liquids was terminated
by a peritectic reaction, which resulted in the formation
of a Cu-rich zone around the Fe-rich primary phase.
3.2 Microstructure of LENS samples
The LENS process can be analysed as a sequence of
discrete events, given that it is a layer-by-layer process.
Each layer of the melted powders composed of
H13-w(Cu) 87.5 % was highly supercooled below the
liquidus temperature; the liquid entered an immiscibility
gap and separated into two liquids17. The phase sepa-
ration generally appeared as dispersed Fe-rich liquid
spheres (L1) in the Cu-rich liquid matrix (L2). Figure 5
shows a typical microstructure of a LENS sample in the
longitudinal direction. The microstructure is not fully
homogeneous, but still much more homogeneous than in
the cast samples. It consists of dark and bright areas
U. ARTI^EK et al.: INFLUENCE OF THE WORKING TECHNOLOGY ON THE DEVELOPMENT ...
738 Materiali in tehnologije / Materials and technology 48 (2014) 5, 735–742
Figure 5: Typical microstructure of a LENS sample layer deposition
in the longitudinal direction
Slika 5: Tipi~na mikrostruktura vzdol`nega prereza vzorca plastne
gradnje procesa LENS
Figure 4: Line microanalysis across a fine dendrite in the Cu-rich
matrix
Slika 4: Linijska mikroanaliza preko finega dendrita v matrici, bogati
z bakrom
depending on the distribution and size of the Fe-rich
phase, which are a consequence of the heat-affected zone
(HAZ) and re-melting zone (RMZ). The dark (wide) area
consists of fine copper grains, dispersed fine spherical
particles of the Fe-rich phase and some individual coarse
Fe-rich particles. In the bright (tight) area – the interlayer
zone – the copper grain sizes are bigger, and the area
includes a considerably lower amount of Fe-rich
spherical particles that are also of a smaller size. There
are no individual coarse Fe-rich particles.
If we compare the average size of the dendritic-like
primary phase in the as-cast microstructure, being in the
range of 100 μm, with small Fe-rich spherical-type
particles of the LENS samples, we can see that they are
much smaller (2–8 μm) and also more uniformly distri-
buted. The size of these Fe-rich particles depends on the
dark/bright area of the microstructure. Some individual
coarse Fe-rich particles (A) can also be observed, but
they occur less often, normally in the dark area of the
microstructure, and belong to the size range of 100 μm
(Figure 5).
In the microstructure, a uniformly distributed micro-
spherical-type gas porosity is present. The gas dissolved
or entrapped in the melt did not have sufficient time to
escape to the top of the melt pool due to a rapid solidifi-
cation. However, the concentration of porosity is much
smaller than it can often be because of the spherical
morphology of atomized powders10,11. An light micro-
graph (Figure 5) clearly shows the presence of porosity,
indicated by small dark spots, having a diameter smaller
than 50 μm.
For a better understanding of the microstructure
evolution, we clad one single layer on the substrate. In
this case, the primary microstructure was preserved
because there was no HAZ or RMZ caused by additional
layers. The temperature distribution was different. The
absence of HAZ and RMZ, and significantly higher cool-
ing rates because of the cold-plate substrate prevented
the Fe-rich liquid spheres (L1) from coarsening. There
were no individual coarse Fe-rich particles. The results
show that the resultant microstructure with a very fine
grain size of approximately 1 μm is homogeneous, and
the minority Fe-rich phase spheres are homogenously
dispersed in the Cu-rich matrix so that there are no dark
or bright areas (Figure 6).
A further deposition of layers led to the changes in
the primary single-layer microstructure. As the previous
layer was re-melted, some of the Fe-rich particles lifted
up into the new layer due to the lower density of iron
compared to copper. RMZ provides a directional solidifi-
cation and a grain growth of Cu-crystals in the bright
area. Ahead of the solid-liquid interface, liquid L1
becomes highly undercooled and begins to solidify
forming a solid and dark area. The latent heat is released
to the previous layer slowing down the growth of Cu in
RMZ (Figure 7a). The figure shows the distribution and
size of the Fe-rich particles in the Cu-rich matrix in the
dark and bright areas of an individual clad layer, in
which the bright area is in the upper part and the dark
area in the lower part of Figure 7b.
The theoretical layer thickness is 355 μm. For the
cyclical layer deposition technique, an interlayer zone –
the bright area between two deposition layers – is
characteristic, being a consequence of re-melting. The
bright and dark areas depend on the local heat input and
U. ARTI^EK et al.: INFLUENCE OF THE WORKING TECHNOLOGY ON THE DEVELOPMENT ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 735–742 739
Figure 7: Border area of an individual layer recorded with: a) LM,
and b) SEM micrograph of the transitional zone between a dark and a
bright area
Slika 7: Mejno podro~je posamezne plasti, posneto z: a) LM in b)
SEM-posnetek prehoda iz temnega v svetlo podro~je
Figure 6: Single-layer weld build-up shows a very fine-grained micro-
structure: a) LM micrograph, b) SEM BEI micrograph
Slika 6: Ena sama navarjena plast izkazuje zelo finozrnato mikro-
strukturo: a) LM-posnetek, b) SEM BEI-posnetek
the mechanism of the turbulence in the melt due to
cladding the next layer (Figure 5).
After the liquid-phase separation, Fe-rich droplets
grow and coagulate in order to reduce the interface area
with the Cu-rich phase. Since the system still remained a
complete liquid during the period between the separation
and solidification, the liquid Fe-rich spheres can grow
and move relative to the Cu-rich matrix and to each
other.
Several mechanisms have been proposed to explain
the evolution of size distribution of the dispersion phase,
including the wetting behaviour, Ostwald ripening and
coarsening of droplets due to collisions, whereby the size
of the dispersed L1 spheres increases with the increasing
undercooling. Generally, as the cooling rates become
higher, the solidification time becomes shorter and the
microstructure becomes finer18. The area with a higher
amount of dispersed Fe-rich particles exhibits very fine
grains of the Cu-rich matrix (Figure 7).
High undercooling favours a long interval between
the separation and nucleation temperatures. As shown in
the upper part of Figure 8, there was enough time for the
spheres to grow through coagulation or coalescence.
With the coalescence, some droplets collided with each
other so that they may have mutually lost surface energy
by joining to form a larger single one. Coagulation, on
the other hand (a lower temperature, a higher viscosity)
is a process when particles come together irreversibly,
i.e., they get stuck together and cannot be separated.
Coagulation and coalescence constitute a process, in
which fine, dispersed, primary spherical Fe-rich particles
(2–8 μm) aggregate together to form individual coarse
Fe-rich particles that belong to the size range of 100 μm.
Furthermore, low undercooling provides a short
coarsening time, in which a free movement and coagu-
lation cannot occur. Therefore, the coarsening of the L1
spheres in this undercooling range should be attributed to
Ostwald ripening19. This is a mechanism, allowing the
droplets to coalesce due to the solute diffusion between
them. The solute solubility depends on the curvature of a
droplet; the smaller the curvature, the larger is the radius
and the lower is the solubility. This diffusion-dependent
coarsening mechanism plays a dominant role when the
dispersion phase has a small diameter. With an increase
in the droplet radius and a decrease in the temperature,
the solid-liquid interface tension increases, whereas the
solubility decreases, resulting in the weakening of
Ostwald ripening. This mechanism can be observed in
the lower part of Figure 8, in the bright area.
It should also be noted that the secondary phase
separation was observed in the highly undercooled areas
of the H13-w(Cu) 87.5 % alloy, i.e., inside the L1 phase,
and some Cu-rich spheres occurred due to the secondary
phase separation (Figure 9), which is a monotonic
increasing function of undercooling. Multi-phase separa-
tion has been rarely observed in stable metallic
immiscibles. Since liquid metals exhibit a low viscosity
and high diffusion coefficient, less time is needed to
adjust the composition of the primary phases20. However,
as undercooling increases, the viscosity rises and diffu-
sion coefficient declines. In this case, a complete diffu-
sion is absent and a multi-phase separation in the liquid
becomes more likely to occur.
3.3 Mechanical properties
The solidification parameters of the alloys directly
affect the microstructures of the alloy systems, and also
significantly influence their mechanical behaviours. The
LENS samples have higher tensile-strength values and
the cast samples are more ductile (Figure 10). This is
due to their microstructures, which are finer and more
homogenous in the LENS samples that solidified at
much higher rates. The uniformly distributed particles of
U. ARTI^EK et al.: INFLUENCE OF THE WORKING TECHNOLOGY ON THE DEVELOPMENT ...
740 Materiali in tehnologije / Materials and technology 48 (2014) 5, 735–742
Figure 8: Coagulation and coalescence mechanisms in liquid-liquid
mixtures and the Ostwald ripening mechanism in solid-liquid mixtures
Slika 8: Mehanizem koagulacije in koalescence v zmesi teko~e-teko~e
in mehanizem Ostwaldovega zorenja v zmesi trdo-teko~e
Figure 9: SEM BEI micrograph of a secondary phase separation
inside an Fe-rich particle
Slika 9: SEM BEI-posnetek prikazuje sekundarno lo~itev faze v
delcu, bogatem z `elezom
the Fe-rich phase, which are of a smaller size and in a
larger quantity, represent the regions that require an
increased amount of energy for the dislocations to pass
through.
The elongation (	) of cast samples is greater than the
elongation of LENS samples. This is due to different
solidification rates of the alloys. A LENS microstructure
is the result of faster cooling and solidification – a
metastable solidification where the Cu-rich matrix with a
higher strength is formed (containing more alloying
elements and precipitates due to which 	 is smaller). In
the case of cast samples where the solidification is
slower, the matrix – practically pure copper – has a lower
strength, resulting in a larger 	.
There is a trend of the LENS samples to have a
slightly higher value of elastic modulus E than the cast
samples because they are more metastable due to a
greater number of the alloying elements in the Cu-rich
matrix, affecting the bond strength.
These characteristics of the microstructure are
reflected on the results of the microhardness measure-
ments. A very fine-grained microstructure, a smaller
size, a larger quantity of the uniformly distributed
Fe-rich spherical phase and the amounts of the alloying
elements in the Cu-rich matrix result in significantly
higher average microhardness values of the LENS
samples. Table 3 shows the average measures of the
microhardness values of the cast samples in the areas of
dendrites, near dendrites and without dendrites and the
microhardness of the LENS samples in Fe-rich phases,
Cu-rich phases and the re-melted HAZ.
4 CONCLUSIONS
In this research, two different technologies were
compared, i.e., the LENS technology and the conven-
tional mould casting to produce an alloy with the same
chemical composition, H13- w(Cu) 87.5 %. The obtained
results can be summarized as follows:
• During the solidification of the conventional castings,
the melt is undercooled and separated into two pha-
ses. The Fe-rich phase solidifies as the leading phase
under the non-equilibrium condition and the solidi-
fication of the Cu-rich phase proceeds under the
near-equilibrium condition according to the phase
diagram.
• Large local variations in the temperature gradient of
the castings cause a very inhomogeneous microstruc-
ture.
• The microstructure of the LENS samples is much
more homogeneous than that of the cast samples.
During the LENS process, where a much higher
undercooling takes place, the undesired dendritic
morphology of primary Fe-rich crystals becomes
spherical.
• There is a significant change in the microstructure of
a LENS sample between the first layer and the
following layers that are re-melted and heat affected,
resulting in dark and bright areas of the microstruc-
ture. The largest differences are due to HAZ, where a
bright area occurs, containing a very small amount of
Fe-rich spherical particles. The bright area represents
the weaker part of the material, which should be
minimized as much as possible by optimizing the
energy intake. However, the microstructure is suffi-
ciently homogeneous to allow homogeneous mecha-
nical properties of the bulk samples.
• The porosity found in the LENS samples was con-
sidered low and uniformly distributed due to the
spherical morphology of the atomized powders.
• The tensile data show that the as-deposited yield
strength of the LENS fabricated materials is substan-
tially higher than that of the cast materials; better
properties are obtained as a result of a rapid solidi-
fication and grain refinement. The tensile data and
microstructure characterisation indicate a good
metallurgical bonding of individual layers of LENS
deposits.
• It has been shown that the H13-w(Cu) 87.5 % alloy
fabricated with the LENS technology can perform
well in real-life applications. A good thermal conduc-
tivity of copper and a high wear resistance of steel
can be achieved without any cracks as well as better
U. ARTI^EK et al.: INFLUENCE OF THE WORKING TECHNOLOGY ON THE DEVELOPMENT ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 735–742 741
Figure 10: Stress-strain curves of the cast and LENS alloys at room
temperature
Slika 10: Krivulja napetost – raztezek lite in LENS-zlitine pri sobni
temperaturi
Table 3: Results of the average microhardness-measurement values
according to Vickers
Tabela 3: Rezultati povpre~nih vrednosti merjenja mikrotrdote po
Vickersu
CASTING
of dendrites near thedendrites
without
dendrites
194 HV 0.01 67 HV 0.01 49 HV 0.01
LENS
Fe Cu Cu (HAZ)
734 HV 0.01 102 HV 0.01 81 HV 0.01
mechanical properties compared to conventional
mould casting.
Acknowledgments
The research is partially funded by the European
Social Fund. Invitations to tenders for the selection of
the operations are carried out under the Operational
Programme for Human Resources Development for
2007–2013, 1. development priority: Promoting entre-
preneurship and adaptability, the priority guidelines, 1.1:
Experts and researchers for enterprises to remain com-
petitive.
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742 Materiali in tehnologije / Materials and technology 48 (2014) 5, 735–742
B. [U[TAR[I^ et al.: MICROSTRUCTURE CHARACTERISTICS OF THE Al-w(Cu) 4.5 % MODEL ALLOY
MICROSTRUCTURE CHARACTERISTICS OF THE
Al-w(Cu) 4.5 % MODEL ALLOY
MIKROSTRUKTURNE ZNA^ILNOSTI MODELNE ZLITINE
Al-w(Cu) 4,5 %
Borivoj [u{tar{i~1, Monika Jenko1, Bo`idar [arler1,2
1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2Laboratory for Multiphase Processes, University of Nova Gorica, Vipavska 13, 5000 Nova Gorica, Slovenia
borivoj.sustarsic@imt.si
Prejem rokopisa – received: 2014-04-14; sprejem za objavo – accepted for publication: 2014-05-12
Samples of the model alloy Al-w(Cu) 4.5 % with a controlled microstructure obtained at different cooling rates were
synthesized. We investigated the microstructure and the microchemistry using light and scanning electron microscopy in order
to determine the segregation on the macro and micro scales, depending on the cooling rate of the synthesized model alloy.
Theoretical thermodynamic and kinetic analyses of the model alloy were also performed.
Keywords: model alloy Al-w(Cu) 4.5 %, thermodynamics, kinetics, microstructure, microchemistry, segregation
Pripravili smo vzorce modelne zlitine Al-w(Cu) 4,5 % s kontrolirano mikrostrukturo, dose`eno pri razli~nih hitrostih ohlajanja.
[tudirali smo mikrostrukturo in mikrokemijo s svetlobnim (LM) in vrsti~nim elektronskim mikroskopom (FE SEM) ter
uporabili analizni metodi EDS in AES za dolo~itev obsega izcejanja na makro- in mikronivoju v odvisnosti od hitrosti ohlajanja.
Naredili smo tudi teoreti~ne termodinamske in kineti~ne preiskave modelne zlitine.
Klju~ne besede: zlitina Al-w(Cu) 4,5 %, termodinamika, kinetika, mikrostruktura, mikrokemija
1 INTRODUCTION
The selected alloy is of the eutectic type (Figure 1)
and is positioned mainly in the region of the homoge-
neous solid solution of -Al and the secondary -phase
(Al2Cu). This type of alloy can be heat treated by the
so-called precipitation hardening (combination of ho-
mogenisation annealing, fast cooling and natural/artifi-
cial ageing) due to the improvement in the mechanical
properties. CALPHAD-based (Calculation of Phase
Diagrams)1,2 software based on theoretical thermodyna-
mic and kinetics also enable a determination of the
equilibrium and metastable phases in more complex
systems. The theoretical binary phase diagram (Figure
2) calculated by ThermoCalc1 is in a good agreement
with the experimental diagram,3,4 shown in Figure 1. A
theoretical calculation (at standard pressure 1 bar) for the
pure binary alloy predicts the existence of only three
phases in the whole (20 °C to 700 °C) temperature
region. Between 20 °C and 521 °C two phases coexist:
crystals of the -Al solid solution and the eutectic
-phase (Al2Cu). Between 521 °C and 564 °C only the
-Al phase is present (complete solid solubility of Cu in
Al) and then between 564 °C and 648 °C exists the
two-phase -Al + L region. Above 648 °C only the
liquid L is still present. A theoretical calculation shows
that the Cu content in the Al2Cu phase increases with
temperature from mass fractions (w) 46 % to 47.7 %.
The solubility of Cu in the -Al solid solution is minimal
at 20 °C (only approx. 5.3 × 10–4) and then increases up
to approximately w = 5 % at 560 °C. Theory also
predicts the existence of the ’ (theta prime) phase up to
339 °C, containing approximately 54 % Cu and 46 % Al
in metastable equilibrium.
The selected alloy is a simple binary alloy (nominal
w(Cu) = 4.5 %); however, bulk and microchemical anal-
yses by SEM/EDS have shown that the synthesised alloy
also contains some impurities and trace elements be-
cause of the use of impure raw materials (technical pu-
rity wp > 99.7 %). Therefore, in the alloy, besides Cu
some Si is also present (approx. w(Si) = 0.14 %), Fe
(approx. w(Fe) = 0.09 %) and traces of Ni and Mg. The
Materiali in tehnologije / Materials and technology 48 (2014) 5, 743–752 743
UDK 669.715'3:620.187:536.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)743(2014)
Figure 1: Experimentally determined equilibrium binary phase
diagram for Al-Cu with a designated position of the model alloy3,4
Slika 1: Eksperimentalno dolo~en ravnote`ni binarni fazni diagram
Al-Cu z ozna~enim mestom, kjer se nahaja modelna zlitina3,4
Al, Mg and Si have a large affinity for oxygen (Gf =
–277.8 kJ/mol at the melting point of Al). Therefore,
some surface oxidation and a thin Al2O3 film is formed
during the alloy synthesis if a very pure protective atmo-
sphere (Ar or N2) is not used. Some complex inclusions
containing Al2O3, MgO and SiO2 can also be formed.
For this type of alloy there is a typical dendrite
morphology of solidification during the cooling and
casting into a sand, metal or graphite model at normal
cooling rates (0.1 K/s to 100 K/s). The measure for
cooling rate and segregation is the so-called Secondary
Dendrite Arm Spacing (SDAS). It generally follows
exponent law,  = k · –n; i.e., the larger is the cooling
rate , the finer are the dendrites and the smaller is the
space between the secondary dendrite arms , as well as
the alloy being more micro homogeneous over its
volume. Some theoretical models exists5,6 for a
determination of , and recently some researchers have
also tried to predict  with ANNs (Artificial Neuron
Networks)7. For Al alloys different values for the
alloy-dependent parameter k (approx. 110) and exponent
n (approx. 0.30) are reported. Generally, the parameter k
depends on the chemical composition; i.e., decreasing
with an increasing concentration of alloying elements.
Figure 3 shows an example of the secondary dendrite
arm spacing  vs. cooling rate  diagram for the selected
values of k  109 and n  0.33. One can clearly see that
for the selected parameters  decreases from approxi-
mately 500 μm at cooling rate  = 0.01 °C/s, over 110
μm at  = 1 °C/s, and finally at  = 100 °C/s it is only
approximately 24 μm.
The physical and mechanical properties of metal-
based alloys are dependent on the chemical composition
(alloying and trace elements) and the microstructure
controlled by the solidification rate connected with the
type and geometry of the model and product, respecti-
vely. A higher initial solidification rate (cooling rate in
the mushy zone) produces a finer dendrite morphology
of solidification, a smaller segregation of alloying ele-
ments, as well as better are mechanical properties
(higher yield/tensile strength and hardness). The refine-
ment of the microstructure also improves the ductility
and toughness for a given strength level. Recently, new
software6 has been developed for the relatively accurate
prediction of the mechanical properties of Al alloys over
a wide range of chemical compositions and solidification
conditions.
Figure 4 shows the theoretical equilibrium thermo-
dynamic phase stability of the model alloy with actual
chemical composition determined at the Institute of
Metals and Technology (IMT), Ljubljana Slovenia
(Table 1). It is clear that in this case seven phases are
stable in the temperature region between 20 °C and 700
°C. However, also in this system the main phases remain
-Al (fcc-A1 solid solution), Al2Cu and liquid L. The
theoretical calculation predicts the existence of the
pro-eutectic -phase (Al2Cu) between 20 °C and 514 °C.
The solid solution -Al is stable up to 646.4 °C. In this
case the autonomous one-phase region of the -Al solid
B. [U[TAR[I^ et al.: MICROSTRUCTURE CHARACTERISTICS OF THE Al-w(Cu) 4.5 % MODEL ALLOY
744 Materiali in tehnologije / Materials and technology 48 (2014) 5, 743–752
Figure 3: Prediction of the secondary dendrite arms space vs. cooling
rate for the Al-based alloy and selected values of the alloy-dependent
parameters
Slika 3: Napoved razdalje med sekundarnimi dendritnimi vejami v
odvisnosti od hitrosti ohlajanja za zlitino na osnovi Al in izbrane
zlitinske parametre
Figure 2: Theoretical equilibrium binary phase diagram Al-Cu, calculated by ThermoCalc1 (TCBin database)
Slika 2: Teoreti~ni ravnote`ni binarni fazni diagram Al-Cu, izdelan s ThermoCalc-om1 (podatkovna baza TCBin)
solution does not exist, because the Fe-based inter-
metallic phase Al7Cu2Fe is stable up to 579 °C. The
optimal temperature interval for homogenization anneal-
ing is between 547 °C and 559 °C. But in this tempera-
ture region there is also a small content of Al7Cu2Fe.
Only liquid is present above 646.4 °C. Silicon has a low
solubility in the present alloy. Therefore, it appears in the
temperature region between 20 °C and 223 °C as -phase
(AlFeSi) and between 223 °C and 340 °C as elementary
crystals of Si, respectively. As already mentioned, in the
temperature region between 223 °C and 579 °C, the
intermetallic phase Al7Cu2Fe is also stable, because of
the presence of iron in the model alloy. Besides this, the
presence of traces of Ni can stabilize the Al7Cu4Ni phase
in the temperature region between 20 °C and 548 °C.
Some traces of Mg are also detected in the model alloy.
In this case we can also expect the presence of the
Al5Cu2Mg8Si6 intermetallic phase. However, it should be
noted that the content of Ni and Mg in the model alloy is
very low and these phases could be neglected.
2 EXPERIMENTAL
The model alloy Al-w(Cu) 4.5 %, was prepared in the
frame of the project entitled ½Advanced modelling and
simulation of liquid-solid state processes½8. In the frame
of this project the development of the microstructure dur-
ing solidification in the mushy zone are studied on the
macro and micro levels for the selected Al- and Fe-based
complex alloys. The microscopic model is to be solved
by the cellular automata concept.
2.1 Synthesis of the model alloy
The model alloy was synthesised in a graphite pot of
a laboratory melting furnace 10 kW under an Ar protec-
tive atmosphere. The furnace is one of the basic compo-
nents of the Melt-Spinner M-10 Marco Materials Inc.
(Figure 5a) which serves primarily for the preparation of
rapidly solidified ribbons with the casting of the melt on
a rotating Cu-wheel.
As the basic raw materials for the preparation of the
model alloy, commercially available materials of techni-
cal purity (Al 99.7 % of manufacturer Impol Slovenska
Bistrica and commercial Cu 99.9 %) were used. Four
batches of 200 g were prepared. Each weight was melted
and then cooled under natural cooling conditions down
to the room temperature. The chamber of the melt-spin-
ner was evacuated with a rotary vacuum pump (absolute
pressure approx. 10 Pa) and then filled with Ar (over
pressure 60 kPa abs.) before melting. Individual weights
were heated up in the inductive melting furnace to 800
°C in 20 min, homogenized for 10 min and then cooled
down. The temperature was followed with a DataLogger
and measured with a Pt-PtRh10 thermocouple, located in
the middle of the graphite melting pot. It was protected
against the melt by a ceramic (alumina) protective tube.
The final bulk chemical composition of the prepared
alloys was checked using a fast portable XRF (X-Ray
Fluorescence) analyser XL3Thermo Fischer Scientific
Niton and a more accurate classical ICP OES (Ion
Coupled Plasma – Optical Emission Spectroscopy)
Agilent 720 instrument with a lower limit of detection (w
< 0.001 % of individual element). Table 1 shows the
results of both chemical analyses. Some impurities are
detected because of the use of technically pure raw
B. [U[TAR[I^ et al.: MICROSTRUCTURE CHARACTERISTICS OF THE Al-w(Cu) 4.5 % MODEL ALLOY
Materiali in tehnologije / Materials and technology 48 (2014) 5, 743–752 745
Figure 5: a) Chamber of melt-spinner with inductive melting furnace
and b) laboratory tube furnace for heat treatment of the model alloy
Slika 5: a) Laboratorijska naprava Melt-Spinner z induktivno talilno
pe~ico in b) laboratorijska cevna pe~ za toplotno obdelavo modelne
zlitine
Figure 4: Theoretical equilibrium thermodynamic stability of the
phases in the alloying system with the real chemical composition of
the model alloy, calculated by ThermoCalc1.
Slika 4: Teoreti~na ravnote`na termodinamska stabilnost faz v
sistemu z dejansko kemijsko sestavo modelne zlitine, izra~unana s
ThermoCalc-om1
materials. Besides Al and Cu, some Fe and Si, as well as
traces of Ni, Zn and Mg, were detected.
Table 1: Average bulk chemical composition of prepared batches of
model alloy in mass fractions, w/%
Tabela 1: Povpre~na kemijska sestava izdelanih {ar` modelnih zlitin v
masnih dele`ih, w/%
Chemical
composition Cu Si Zn Fe Mg Ni Al
XRF 4.30 0.14 0.010 0.12 – – balance
ICP OES 4.45 0.14 0.002 0.09 < 10–3 0.006 balance
Five cylinders of diameter 45 mm and height 40 mm
weighing approximately 200 g were prepared in this way
(Figure 6) for further experiments and microstructure
investigations. The cylinders were cleaned of surface
oxidation by smooth drilling. Further heat treatments of
the cylinder were performed due to the formation of an
appropriate microstructure at different cooling rates. It is
very difficult to obtain the required cooling rate because
the real cooling conditions vary over the cross-sections
(volume) of relatively large samples, and they are not
constant in different temperature regions. The laboratory
equipment also does not enable the use of a larger
number of thermocouples at different locations of the
samples. It can be considered that below 200 °C signi-
ficant microstructure changes do not happen in a relati-
vely short period of time. Therefore, we can assess the
average cooling rate above this temperature. However, in
the literature6 one can find different approaches to
cooling-rate definition (initial, average, for a given
temperature interval etc.) depending on how it can influ-
ence the microstructure formation. During the planning
of the present project we did not have in mind that for
the formation of a selected Al-4.5 Cu binary alloy that
the initial cooling rate and cooling rate in the mushy
zone are important. Therefore, it was planned that three
different characteristic average cooling rates will be
obtained; i.e., very slow (< 0.1 °C/min or 0.0017 °C/s),
natural cooling rate (30 °C/min to 40 °C/min or 0.5 °C/s
to 0.7 °C/s) and very fast quenching (300–400 °C/min or
5.0 °C/s to 6.7 °C/s). Actually, the following experimen-
tal cooling rates are obtained:
Sample 0 to 4 – reference materials – only synthe-
sised alloy (melted in graphite pot of inductive furnace
and cooled down). Figure 7 shows the temperature pro-
file of the alloy synthesis, i.e., heating and cooling in a
graphite melting pot. During cooling there is a clearly
visible solidification interval between approximately 645
°C and 550 °C. This is in relatively good agreement with
the theoretical prediction, which predicts that the solidi-
fication starts at 648 °C and ends at 564 °C (Figure 4).
The obtained cooling rate between 645 °C and 100 °C is
approximately 10 °C/min and approximately 0.17 °C/s,
respectively. For the formation of the dendrite morphol-
ogy of solidification the cooling rate in the mushy zone
is important. This was assessed at approximately 30
°C/min and 0.5 °C/s, respectively. In this case, one can
estimate from Figure 3 the SDAS on 130 μm.
Samples 0 and 1 were retained in the original state
for metallographic investigations. However, the samples
2, 3 and 4 were additionally heat treated in order to ob-
tain the planned cooling rates. The samples were heated
up to 610 °C into the mushy zone (semi-solid state) and
then cooled down.
With sample 2 we tried to simulate natural cooling
conditions in the tube furnace (Figure 5b). A half part of
the cylinder was heated up to 610 °C for 10 min and then
a ceramic tube was pulled from the heating chamber of
the furnace. The obtained cooling rate in the temperature
interval between 610 °C and 200 °C was approximately
12 °C/min (0.2 °C/s) and below 200 °C it was approxi-
mately 0.7 °C/min (0.012 °C/s) (Figure 8). The maxi-
mum cooling rate in the mushy zone was estimated to be
approximately 0.4 °C/s (SDAS  138 μm). But it has to
be noted that the prepared alloy was not completely
melted. As one can see the average cooling rate of the
material is rather lower than planned 30–40 °C/min
(0.5–0.7 °C/s) and similar to that obtained during the
alloy synthesis.
Sample 3 was heated to 610 °C for 10 min and then
fast cooled, i.e., quenched in water. The estimated aver-
B. [U[TAR[I^ et al.: MICROSTRUCTURE CHARACTERISTICS OF THE Al-w(Cu) 4.5 % MODEL ALLOY
746 Materiali in tehnologije / Materials and technology 48 (2014) 5, 743–752
Figure 7: Temperature profile obtained during the synthesis of the
model alloy Al-w(Cu) 4.5 %
Slika 7: Temperaturni profil sinteze vzorca modelne zlitine Al-w(Cu)
4,5 %
Figure 6: a) Schematic presentation of the cast cylinder of the model
Al-w(Cu) 4.5 % alloy and a) its cutting into specimens for the pre-
paration of metallographic samples
Slika 6: a) Shemati~ni prikaz ulitka modelne zlitine Al-w(Cu) 4,5 %
in b) njegov razrez za pripravo metalografskih vzorcev
age cooling rate of this sample was 1220 °C/min (20.3
°C/s), which is rather faster than planned.
Sample 4 – with this sample we tried to simulate very
slow cooling. The cylinder was cooled down from 610
°C over approximately 4 d with an approximate cooling
rate of 0.1 °C/min, (0.002 °C/s) (Figure 9). In this case
the maximum cooling rate was 0.7 °C/min and 0.01 °C/s,
respectively (SDAS  500 μm).
From the above-described experiments we can con-
clude that the experimental work was not completely
successful, as planned in the frame of the project. In
spite of this the obtained results are very interesting and
are published as follows.
2.2 Microstructure investigations of the model alloy
Cast and heat-treated cylinders of the model alloy
Al-w(Cu) 4.5 % were cut up into slices and prepared for
microstructure investigations (Figure 10). The metallo-
graphic specimens were prepared with Struers equipment
(electron saw Accutom 50, automatic press Pronto-
Press-20 and grinding/polishing apparatus Abramin with
MD-system). The microstructure characterization with a
light (LM, Nikon Microphot FXA with 3CCD video
camera Hitachi HV-C20AMP and software AnalySIS
PRO 3.1) and a scanning electron microscope (SEM;
JEOL FE HR JSM-6500F) combined with micro-che-
mical analysis based on a measurement of the dispersed
kinetic energy of X-rays (EDS – Energy Dispersive
X-ray Spectrometer) on polished and etched metallo-
graphic plates 20 mm × 20 mm (specimens) were per-
formed. A systematic non-continuous point profile
(10-points per 20 μm) and surface (mapping) EDS
analyses at different locations (Figure 11) of the samples
were then performed.
3 RESULTS AND DISCUSSION
The prepared model alloys were cooled down with
cooling rates between 0.002 °C/s and 20.3 °C/s. In all
these cases one can expect a dendrite morphology of
B. [U[TAR[I^ et al.: MICROSTRUCTURE CHARACTERISTICS OF THE Al-w(Cu) 4.5 % MODEL ALLOY
Materiali in tehnologije / Materials and technology 48 (2014) 5, 743–752 747
Figure 11: Schematic presentation of EDS analyses’ locations on
metallographic samples of the model alloy Al-w(Cu) 4.5 %
Slika 11: Shemati~ni prikaz analiznih mest SEM/EDS na metalo-
grafskih obruskih modelne zlitine Al-w(Cu) 4,5 %
Figure 9: Diagram of heating and very slow cooling of sample 4 of
model alloy Al-w(Cu) 4.5 % in tube furnace
Slika 9: Diagram segrevanja in zelo po~asnega ohlajanja vzorca 4
modelne zlitine Al-w(Cu) 4,5 % v cevni pe~i
Figure 8: Heating/cooling diagram for model alloy Al-w(Cu) 4.5 % in
the ceramic tube of the furnace, sample 2, ceramic tube with sample
pulled from the furnace chamber
Slika 8: Diagram segrevanja in ohlajanja modelne zlitine Al-w(Cu)
4,5 % v cevi cevne pe~i (po segrevanju je cev potegnjena iz pe~i),
vzorec 2
Figure 10: Schematic presentation of ingot cutting of the model alloy
for the preparation of metallographic samples
Slika 10: Shemati~ni prikaz razreza ingota modelne zlitine Al-w(Cu)
4,5 % za pripravo metalografskih obruskov
solidification with different SDAS in the range between
500 μm and 35 μm, and, actually this was obtained.
Figure 12 shows a typical dendrite morphology of the
solidification formed during the alloy synthesis. Figure
13 shows this microstructure visible under the SEM at
different magnifications. At the highest magnification
one can clearly see the micro-chemical segregation due
to non-equilibrium solidification.
Microchemical point, profile and EDS mapping were
performed at different locations on the samples in order
to obtain information about the local microchemical
composition and the alloy segregation. Figure 14 shows
an example of point (Spectrums 1, 2 and 3) and mapping
(Spectrum 4) EDS analyses. One can clearly see that in
the interdendritic regions a secondary Al2Cu phase, as
well as Al2O3 oxide based is formed. It could be antici-
pated that the oxide inclusions are the nuclei for a sec-
ondary phase precipitation. Figure 15 shows an example
of profile analyses across the dendrite region. It is clear
that the Cu-rich secondary phase has approximately
B. [U[TAR[I^ et al.: MICROSTRUCTURE CHARACTERISTICS OF THE Al-w(Cu) 4.5 % MODEL ALLOY
748 Materiali in tehnologije / Materials and technology 48 (2014) 5, 743–752
Spectrum O Al Ni Cu
Spectrum 1 4.37 67.29 28.34
Spectrum 2 4.01 68.25 27.74
Spectrum 3 5.23 65.72 0.58 28.47
Spectrum 4 99.03 0.97
Figure 14: Area (metal matrix of -Al solid solution) and point
SEM/EDS analyses (secondary phase) of model alloy Al-w(Cu) 4.5 %
Slika 14: Ploskovna (kovinska matrica -Al) in to~kovna (sekundarna
faza) SEM/EDS modelne zlitine Al-w(Cu) 4,5 %
Figure 12: Microstructure of model alloy Al-w(Cu) 4.5 % visible
under LM: a) sample 1, SDAS = 83 μm and b) sample 2, SDAS = 100
μm, magnification 50-times
Slika 12: Mikrostruktura modelne zlitine Al-w(Cu) 4.5 %, vidna pod
LM: a) vzorec 1, SDAS = 83 μm in b) vzorec 2, SDAS = 100 μm,
pove~ava 50-kratna
Figure 13: SE images of microstructure of model alloy Al-w(Cu) 4.5
% at different magnifications; sample 1: a) magnification 100-times,
b) magnification 500-times and c) magnification 2000-times
Slika 13: SE-posnetki mikrostrukture modelne zlitine Al-w(Cu) 4,5
%, vzorec 1: a) pove~ava 100-kratna, b) pove~ava 500-kratna in c)
2000-kratna
Figure 15: EDS profile analysis from Al metal matrix across
secondary phase back to the metal matrix; Al and Cu concentration
distribution
Slika 15: EDS profilna analiza iz kovinske -Al osnove preko sekun-
darne faze in nazaj v kovinsko osnovo; porazdelitev koncentracije Al
in Cu
w(Cu) = 44 %, but the metal matrix, i.e., the -Al solid
solution, has approximately w(Cu) = 5 %.
Non-continuous profile (10-points per 20 μm)
SEM/EDS analyses of metallographic samples cooled
down with an average cooling rate of 10–12 °C/min have
shown that four typical concentration profiles exist:
Generally, the concentration of Al continuously falls
to a minimum crossing the secondary phase in the inter-
dendritic regions (dendrite pockets) and then again the
Al concentration increases back to the Al-based metal
matrix. Simultaneously, the concentration of Cu changes
in the opposite direction (Figure 15). The average local
microchemical composition of the Al matrix is approxi-
mately w = 95 % Al and w = 5 % of Cu ( amount frac-
tions  = 97.7 % Al and 2.3 % of Cu) and the average
local microchemical composition of the secondary phase
is approximately w = 56 % of Al and 44 % Cu ( = 75 %
Al and 25 % of Cu). Theoretically, the metal matrix -Al
solid solution contains w = 100 % of Al at room tempe-
rature and approximately  = 98 % Al and 2 % Cu at
550 °C. The secondary phase Al2Cu contains theoreti-
cally at room temperature  = 66.7 % of Al and 33.3 %
of Cu up to 68 % of Al and 32 % of Cu at 550 °C. The
experimentally determined composition is close to the
equilibrium composition theoretically predicted at appro-
ximately 550 °C. This is a consequence of the non-
equilibrium solidification and the metastable condition
of the model alloy at room temperature. It also has to be
noted that the EDS information comes from a depth of
about 1 μm to 3 μm. The size and shape of this inte-
raction volume is dependent on the primary beam energy
and the sample material (Figure 16). For the surface
analysis EDS is not the proper analytical method. In the
case of very thin secondary phases, inclusions and
surface analyses one must use a surface-sensitive analy-
tical method such as Auger Electron Spectroscopy
(AES).
Theoretical thermodynamic analyses predict the exis-
tence of some intermetallic phases (Al7Cu2Fe, Al7Cu4Ni,
AlFeSi and crystals of Si) because of the presence of Fe,
Si and Ni in the investigated model alloy. Actually, at
some places inside the secondary Al2Cu phase we de-
tected the presence of these elements. In the middle of
the secondary phase, where the concentration of Cu is
the highest, the EDS analyser detected small concentra-
tions of Ni (Figure 17). The intermetallic phase
B. [U[TAR[I^ et al.: MICROSTRUCTURE CHARACTERISTICS OF THE Al-w(Cu) 4.5 % MODEL ALLOY
Materiali in tehnologije / Materials and technology 48 (2014) 5, 743–752 749
Figure 18: AES mapping of a precipitate of secondary phase in the
model alloy Al-w(Cu) 4.5 % (sample 2) showing the distribution of
selected elements Al, O, Ni and Cu and a clearly visible increased
concentration of Ni
Slika 18: AES-mikroskopija porazdelitve izbranih elementov Al, O,
Ni in Cu v izlo~ku sekundarne faze v modelni zlitini Al-w(Cu) 4,5 %
(vzorec 2) z dobro vidnim podro~jem pove~ane koncentracijo Ni
Figure 16: The interaction between an incident electron beam and the
solid sample, showing the analysis volumes for Auger, secondary
electrons, back scattered electrons and X-ray fluorescence
Slika 16: Interakcija vpadnega elektronskega curka s povr{ino trdnega
vzorca prikazuje analizni volumen za Augerjeve elektrone, sekundarne
elektrone, povratno sipane elektrone in rentgensko fluorescenco
Figure 17: EDS profile analysis across -Al dendrite region into the
inter-dendritic region and back into the metal matrix; concentration
distribution of Al, Cu and Ni
Slika 17: EDS profilna analiza iz kovinske -Al osnove preko sekun-
darne faze in nazaj v kovinsko osnovo; porazdelitev koncentracije Al,
Cu in Ni
Al7Cu4Ni is stable from room temperature up to 550 °C
and theoretically contains mass fractions w = 51 % Cu,
38 % Al and 11 % of Ni. The EDS analyses give a lower
Ni concentration (w = 1.5 % to 3.5 %). However, the ra-
tio Cu : Al (58 % Cu : 40 % Al) perhaps confirms that at
these places the Al7Cu4Ni phase or its non-equilibrium
approximate is present. Therefore, we also performed
SEM/AES analyses which suggest that at some places
there is a higher local concentration of Ni (Figure 18).
GIXRD (Grazing incidence X-Ray Diffraction) can help
in the exact identification of the presence of this phase.
However, this investigation was not planned in the frame
of the project. It will be performed later and published
elsewhere.
Aluminium has a high affinity for oxygen; therefore,
it was expected that a slight oxidation of the alloy could
happen during its synthesis. The SEM/EDS microanaly-
ses have shown the presence of oxygen at some locations
at the interdendritic locations. Figure 19 shows the pro-
file EDS analysis across the secondary phase. The pre-
sence of a xSiO2 yAl2O3-based inclusion is clearly visible
at the edge of it.
At some locations only Al2O3 is present. A proof of
this is the simultaneous increased concentration of Al at
these locations (Figure 20).
Finally, EDS analyses have also shown that practi-
cally all the elements of impurities are present at some
locations in the interdendritic regions (Figure 21). This
is important from the mechanical properties point of
view because hard intermetallic phases in the interden-
dritic locations worsen the cohesive strength between the
metal matrix and the interdendritic phase. Additionally,
segregations are more extensive and the SDAS is larger,
respectively, if the cooling rate is lower.
Figure 22 shows the average SDAS measured on
samples 0, 1 and 2. The performed SDAS measurements
show a relatively larger scatter of measured values from
location to location. No evident law or regular change in
SDAS as expected; for example, from the surface to the
middle of the sample. The average measured values are
between 79 μm and 110 μm, which corresponds to the
theoretical initial cooling rate between 1.0 °C/s and
approximately 2.5 °C/s (from 60 °C/min to 150 °C/min.).
These cooling rates are higher than determined during
the experiment (20–30 °C/min) at the thermocouple
location. Besides, the thermocouple was isolated with an
alumina tube against the effects of the melt. Therefore,
further experimental work has to be considerably
improved. In particular, the techniques of cooling-rate
detection during solidification (smaller sample, larger
number of thermocouples at different locations,
B. [U[TAR[I^ et al.: MICROSTRUCTURE CHARACTERISTICS OF THE Al-w(Cu) 4.5 % MODEL ALLOY
750 Materiali in tehnologije / Materials and technology 48 (2014) 5, 743–752
Figure 21: EDS profile analysis across -Al dendrite region into the
inter-dendritic region and back into the metal matrix; concentration
distribution of Al, Cu, Fe, Ni and O
Slika 21: EDS profilna analiza iz kovinske -Al osnove preko sekun-
darne faze in nazaj v kovinsko osnovo; porazdelitev koncentracije Al,
Cu, Fe, Ni in O
Figure 19: EDS profile analysis across -Al dendrite region into the
inter-dendritic region and back into the metal matrix; concentration
distribution of Al, Cu, Si and O
Slika 19: EDS profilna analiza iz kovinske -Al osnove preko sekun-
darne faze in nazaj v kovinsko osnovo; porazdelitev koncentracije Al,
Cu, Si in O
Figure 20: EDS profile analysis across -Al dendrite region into the
inter-dendritic region and back into the metal matrix; concentration
distribution of Al, Cu and O
Slika 20: EDS profilna analiza iz kovinske -Al osnove preko sekun-
darne faze in nazaj v kovinsko osnovo; porazdelitev koncentracije Al,
Cu in O
appropriate melting furnace etc.) must be additionally
improved.
The most important for the formation of the solidi-
fication microstructure is the initial cooling rate and the
cooling rate in the mushy zone, respectively. In the case
of our model alloy the solidification interval is rather
narrow and it is very difficult to control the solidification
process on the relatively large samples that were initially
selected. At the beginning of our experimental work we
did not have enough experiences in the field. However,
we became acquainted with the problems and gained a
lot of experience in the field during the execution of the
present project. In the late phase of the project we
performed much more controlled cooling experiments in
the DSC/TG apparatus (Netzsch STA 449C Jupiter) with
smaller samples ( 4.5 mm × 5 mm, m  0.2 g) and
much more controlled cooling rates in the mushy zone.
However, this equipment does not allow us to perform
the experiments with very high cooling rates. Therefore,
the selected cooling rates were (0.1, 0.5, 1.0, 10 and 25)
°C/min. The DSC cooling curves show a big difference
in solidification and phase precipitation, respectively, if
the cooling rate is changed. A very clear and sharp phase
formation is noticed at the lowest cooling rate (Figure
23). Details of these experiments will be published
elsewhere separately, because of the limited space here.
4 CONCLUSIONS
Samples of the model alloy Al-w(Cu) 4.5 % with a
controlled microstructure obtained at different cooling
rates were synthesized in the frame of the present
investigation. The obtained average cooling rates on
larger samples (Figure 6) in the temperature interval
from 600 °C to 200 °C are estimated to be 0.1 °C/min,
B. [U[TAR[I^ et al.: MICROSTRUCTURE CHARACTERISTICS OF THE Al-w(Cu) 4.5 % MODEL ALLOY
Materiali in tehnologije / Materials and technology 48 (2014) 5, 743–752 751
Figure 23: DSC cooling curves of model alloy Al-w(Cu) 4.5 %, cooled in the temperature region between 700 °C and 500 °C with two different
cooling rates: a) 0.1 °C/min (SDAS > 700 μm) and b) 10 °C/min (SDAS  100 μm), with appertain microstructural characteristics visible by LM.
Slika 23: DSC ohlajevalna krivulja modelne zlitine Al-w(Cu) 4,5 %, ohlajane v temperaturnem obmo~ju med 700 °C in 500 °C s hitrostjo: a) 0,1
°C/min (SDAS > 700 μm) in b) 10 °C/min (SDAS  100 μm), s pripadajo~imi mikrostrukturnimi zna~ilnostmi, vidnimi s svetlobnim
mikroskopom
Sample 0:
0_zg_L 0_zg_D
79.1 94.4 90.6 81.8 98.6 107.9
87.8 89.2 79.0 89.4 91.3 84.0
67.0 90.5 97.8 92.8 90.1 93.3
83.8 82.3 103.9 95.2 81.6 76.3
83.1 87.5 88.6 88.5 79.0 76.4
100.4 102.6 94.5 102.1 108.2 87.7
0_sp_L 0_sp_D
Sample 1:
1_zg_L 1_zg_D
96.8 79.2 98.8 86.9 98.7 84.0
104.7 114.2 85.9 69.4 84.1 94.7
90.5 83.9 72.2 90.8 84.9 95.4
87.6 126.0 79.5 87.3 107.2 76.3
71.8 100.4 87.1 96.1 119.0 98.0
87.1 95.2 105.5 83.5 92.2 97.7
1_sp_L 1_sp_D
Sample 3:
3_zg_L 3_zg_D
106.9 99.7 100.2 89.9 77.5 92.3
98.8 97.7 112.6 94.5 69.9 99.4
92.2 77.1 99.3 84.5 91.4 97.3
94.3 104.7 86.4 90.8 86.5 106.4
109.8 90.3 95.4 104.4 100.8 89.0
100.5 93.2 99.2 110.0 91.0 94.4
3_sp_L 3_sp_D
Figure 22: Average SDAS measured in micrometers, measured on
metallographic LM snapshots on individual samples cooled with
different average cooling rates, in accordance with Figure 11
Slika 22: Povpre~na velikost SDAS v mikrometrih, izmerjena na
metalografskih LM-posnetkih na posameznih vzorcih, ohlajanih z
razli~no povpre~no hitrostjo; vezano na sliko 11
12 °C/min. and 1 220 °C/min. The size of the SDAS is
controlled by the initial cooling rate and the cooling rate
in the mushy zone, respectively. This cooling rate was
estimated to be approximately 0.7 °C to 30 °C/min in the
case of natural cooling and very slow cooling in the
furnace. It is a rather larger cooling rate than the average
measured cooling rate in the middle of the samples. The
average measured SDAS on all the samples are between
approximately 79 μm and 110 μm. This corresponds to
the theoretical initial cooling rate of approximately
between 1.0 °C/s and 2.5 °C/s. The only exception is the
sample with the lowest selected average cooling rate
(approx. 0.002 °C/s). In this case the SDAS was assessed
to be more than 700 μm. But it was difficult to measure
because even the lowest magnification (50-times) of the
LM does not enable exact measurements. For this case
much lower magnifications (10- to 25-times) are
necessary. The performed investigations showed that
more exact measurements of the cooling/solidification
rate, especially in the mushy zone, are necessary. The
experiments performed on smaller samples in the
DSC/TG apparatus enabled better control of the solidifi-
cation, but in a limited range (from 0.1 °C/min to 25
°C/min) of cooling rates. Therefore, new lab equipment
(furnace) with a controllable atmosphere and accurate
temperature control over a wider range of cooling rates
(0.1–100 °C/s) for relatively small samples must be
developed or purchased, respectively.
The next step of the performed investigations was the
preparation of metallographic samples, as well as micro-
structural and microchemical investigations under the
LM and SEM combined with EDS and AES microana-
lyses. The aim of these investigations was a determina-
tion of the segregation at the macro and micro levels,
depending on the cooling rate of the synthesized model
alloy. The microstructural composition of the naturally
cooled model alloy is markedly in the non-equilibrium
state, and similar to those predicted theoretically at 550
°C. Four characteristic states of chemical composition
are detected in the interdendritic regions because of the
presence of impurities (Fe, Si, Ni) in the model alloy, as
well as its slight oxidation during synthesis. For the
synthesis and study of the pure binary Al-w(Cu) 4.5 %,
model alloy it is necessary to select completely pure raw
materials. Impurities can disturb some of the investi-
gations to a certain extent. However, on the other hand, it
can also help in understanding the complexity of
multicomponent alloying systems.
Acknowledgement
The authors wish to thank the Slovenian Research
Agency for financial support in the frame of project no.:
J2-4120 and P2-0132, and to Professor Jo`ef Medved,
Department for Metallurgy and Materials, University of
Ljubljana, for assistance with the DSC/TG experiments.
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3 TAPP – Thermochemical and Physical Properties, A computer data-
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4 E. A. Brandes, G. B. Brook, Smithells Metals Reference Book, 7-th
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B. [U[TAR[I^ et al.: MICROSTRUCTURE CHARACTERISTICS OF THE Al-w(Cu) 4.5 % MODEL ALLOY
752 Materiali in tehnologije / Materials and technology 48 (2014) 5, 743–752
J. BURJA et al.: CHROMITE SPINEL FORMATION IN STEELMAKING SLAGS
CHROMITE SPINEL FORMATION IN STEELMAKING
SLAGS
NASTANEK KROMITNIH SPINELOV V JEKLARSKIH @LINDRAH
Jaka Burja1, Franc Tehovnik1, Jo`ef Medved2, Matja` Godec1, Matja` Knap2
1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2University of Ljubljana, Faculty of Natural Sciences and Engineering, Department of Materials and Metallurgy, A{ker~eva 12,1000 Ljubljana,
Slovenia
jaka.burja@imt.si
Prejem rokopisa – received: 2014-05-13; sprejem za objavo – accepted for publication: 2014-05-21
During the processing of stainless-steel grades in an electric arc furnace (EAF) a considerable amount of chromium can be lost
due to oxidation. These chromium oxides form different phases in the slag, the most stable being the spinel phase. The basicity
of the slag has a major impact on the composition of the chromium oxide phase. A phase analysis revealed two types of
chromium oxide phases, calcium chromites and chromite spinels, which are dependent on the chemistry and the basicity of the
slag. The calcium chromites only form at a high slag basicity, while the chromite spinels form at both high and low basicity. The
effects of ferrosilicon additions were observed as they have a profound effect on both the slag and the chromite spinel chemical
composition.
Keywords: chromite spinel, calcium chromite, stainless steel slag, chromium loss
Med izdelavo nerjavnih jekel v EOP lahko pride do znatnih izgub kroma zaradi oksidacije. Kromovi oksidi v `lindri tvorijo
razli~ne faze, najstabilnej{a je spinelna faza. Bazi~nost `lindre ima velik vpliv na fazno sestavo kromitnih oksidov. Fazna
analiza je pokazala prisotnost dveh vrst kromovih oksidov, odvisnih od kemijske sestave `lindre in bazi~nosti, to sta kalcijev
kromit in kromitni spinel. Kalcijevi kromiti so prisotni ob visoki bazi~nosti, medtem ko so spineli prisotni tako ob visoki kot
nizki bazi~nosti. Opazovani so bili u~inki dodatkov ferosilicija, ki odlo~ilno vplivajo na kemijsko sestavo `lindre in kromitnih
spinelov.
Klju~ne besede: kromitni spinel, kalcijev kromit, nerjavna `lindra, izguba kroma
1 INTRODUCTION
Chromium promotes ferrite1 and is an important
alloying element in steels, especially stainless steels.
Stainless steels contain chromium in excess of the mass
fraction w = 10 %;2 higher amounts of chromium in steel
leads to a higher corrosion resistance.3 Chromium is
added through the melting of stainless-steel scrap and
ferrochromium additions. Ferrochromium is available in
different grades that contain different amounts of chro-
mium and other alloying elements. Carbon plays the
most important role as an alloying element in ferrochro-
mium besides chromium, because oxygen blowing is
needed to remove the carbon from the melt, which can
cause chromium losses.
Chromium-rich slags typically form during the
melting of stainless-steel grades. Studies show that 97 %
of the chromium losses are attributed to the electric arc
furnace (EAF),2 during melting and to a large extent
during the blowing of oxygen into the melt in order to
remove the carbon.4,5 The oxidation of chromium takes
place alongside that of other alloying elements in the
steel, such as aluminum, carbon, silicon, manganese and
a certain amount of iron itself. The reaction of chromium
and oxygen dissolved in steel can be described as:6
2 3 3
1
0
Cr Cr O
1 127 100 250.80 (J)
s 2 s( ) ( )
Δ
+ =
= − +
[O]
G T
(1)
The dominant chromium oxide in the slag at the
melting temperature, without a protective atmosphere, is
Cr2O3.7 Oxides of alloying elements, among them chro-
mium oxides, along with slag-forming oxide additions,
form different phases in the slag. There are two types of
chromium-oxide-based phases that are generally found
in a chromium slag, i.e., chromium-oxide-based spinels,
which are solid solutions, and calcium chromites, which
are stoichiometric compounds.8 Chromium-oxide-based
spinels contain Cr2O3, Al2O3, FeO, MgO and MnO,
while calcium chromites contain only CaO and Cr2O3.
Chromium oxide spinels precipitate in the liquid slag9
and affect the slag’s properties. The slag stiffens and be-
comes inactive, thus preventing the reduction of oxides
and attributing to the loss of alloying elements. When
spinels are formed, the activity of the chromium de-
creases because of the strong bonding in the spinel,
especially in the MgCr2O4.10
Reductive agents such as aluminum, carbon, ferro-
silicon and even calcium carbide are introduced into the
slag in order to minimize the chromium losses.4,5,9,11–15 In
the case of the decarburization of the melt a typical
stainless-steel slag can contain from w = 30 % to 40 %
Cr2O3.15 The aim of this work is to study the phase
composition of chromium-rich slags and the composition
of chromium-oxide-based phases in order to establish the
basic conditions for further studies of the thermodyna-
Materiali in tehnologije / Materials and technology 48 (2014) 5, 753–756 753
UDK 669.187:669.015.8 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(5)753(2014)
mic equilibrium mechanism of the chromium distribu-
tion between the steel and the slag.
2 EXPERIMENTAL
Slag samples were taken with a special spoon during
the steel processing in the EAF. One sample was taken
after the ferrochrome addition and the oxygen blowing
and another after the ferrosilicon addition. The samples
were left to cool; smaller pieces were cut away and put
into the mass for metallographic investigation. They
were grinded and polished and then carbon was evapo-
rated onto the surface to provide electrical conductivity
for the electron microscope. Other pieces of the same
slag samples were crushed in a steel mortar, then in a
steel ball mill, and finally in an agate mortar to achieve
the final fine powder that is needed for the X-ray diffrac-
tion. The slag samples underwent light microscopy
(LM), (Microphot FXA, Nikon), electron microscopy
and electron-dispersive spectroscopy (EDS) analysis
(SEM-EDS, JEOL – JSM6500F) and X-ray diffracto-
metry (XRD, Panalytical XPert Pro PW3040/60).
3 RESULTS AND DISCUSSION
Slag samples, before and after the ferrosilicon addi-
tion, were taken from high and low basicity slag.
The basicity is defined as:
B =
%CaO
%SiO 2
(2)
The samples with high basicity before the ferrosili-
con addition contain both calcium chromites and spinels,
whereas the low-basicity samples after ferrosilicon con-
tain only spinels. The formation of spinels is preferred
over the formation of calcium chromites.16 The spinels
are particles that precipitate at processing temperatures
from liquid slag, and the equations for precipitation
are:17,18
MgO(S) + Cr2O3(S) = MgO ·Cr2O3(S)
ΔG T3
0 30 221 19 945= − − (J). (3)
FeO(S) + Cr2O3(S) = FeO ·Cr2O3(S)
ΔG T4
0 37 706 3846= − + (J). (4)
MgO(S) + Al2O3(S) = MgO ·Al2O3(S)
ΔG T5
0 20 740 1157= − − . (J) (5)
The precipitated spinels are solid solution of all the
oxides mentioned above. The analyzed samples can be
divided into two groups, i.e., those with high basicity and
those with low basicity. The low-basicity samples
contain only chromites and no calcium chromites. The
spinels are of various sizes, from 10 μm to 100 μm, as
shown in Figure 1. The high-basicity slags (Figure 2),
on the other hand, contain large spinels that can reach
J. BURJA et al.: CHROMITE SPINEL FORMATION IN STEELMAKING SLAGS
754 Materiali in tehnologije / Materials and technology 48 (2014) 5, 753–756
Figure 3: Ternary diagram of Mn-Fe-Mg in spinels (mole fraction)
before and after FeSi addition
Slika 3: Ternarni diagram Mn-Fe-Mg v spinelih (molski dele`) pred
dodatkom FeSi in po njem
Figure 1: Microstructure of slag with chromite spinels (B = 0.8, after
FeSi addition)
Slika 1: Mikrostruktura `lindre s kromitnimi spineli (B = 0,8, po
dodatku FeSi)
Figure 2: Microstructure of slag with chromites, and calcium
chromites (B = 2, before FeSi addition)
Slika 2: Mikrostruktura `lindre s kromitnimi spineli in kalcijevimi
kromiti (B = 2, pred dodatkom FeSi)
sizes above 100 μm in diameter. The main difference,
however, is the presence of needle-shaped calcium
chromites. The EDS analysis of the slag samples showed
that the chemical composition of the spinels changes
during the processing. After the ferrosilicon addition, the
content of FeO in the spinels is lowered, as can be seen
in Figure 3.
Iron oxide is the most easily reduced of all the oxides
in the spinels and is therefore the most impacted by the
addition of FeSi. Magnesium oxide, on the other hand, is
very stable in the spinel phase and is not reduced, its
content is also increased by refractory degradation.19,20
FeSi is mainly added in order to reduce the chromium
content in the slag, but a significant part is used to
reduce the iron oxides. The diagram in Figure 3 shows
that FeSi does not reduce the manganese from the spi-
nels very effectively. The composition of spinels follows
the parallel Mn concentration lines in the ternary dia-
grams, showing only slight deviations.
The XRD spectrum of the slag before the FeSi addi-
tion at high basicity (Figure 4) shows the presence of
spinels, calcium chromites and alpha iron.
After the FeSi addition at low basicites there are no
calcium chromites present, only spinels and alpha iron,
as can be seen in Figure 5. There are no more calcium
chromites present.
Alpha iron is the presence of metal droplets that form
due to the entrapment of the melt during the steelmaking
process, either because of mixing during the arc melting
and oxygen blowing or during the reduction of the
oxides to the metal state. The XRD spectra confirm the
observations under both SEM and LM. In fact, the SEM
EDS analysis of the metal droplets revealed that they
contained mostly iron and chromium, hence the ferrite
phase, the composition spanning from w(Cr) = 10 % up
to 20 %.
The examinations of the slags revealed that although
all of the slags contained calcium chromites before the
FeSi additions, none of the samples after the FeSi addi-
tions did. The tendency of calcium oxide to form phases
with silicon oxides is clearly greater than that of forming
them with chromium oxides.
4 CONCLUSION
Chromium slags contain two types of chromium-
based oxides: calcium chromites and chromite spinels.
The composition of the chromites changes during the
processing; the addition of ferrosilicon (FeSi) signifi-
cantly decreases the Fe content in the spinels.
The magnesium contents in spinels increases during
the steelmaking process, both due to the relative increase
of the content (FeO and Cr2O3 reduction) during ferro-
silicon injection and due to MgO dissolution from the
refractory.
Calcium chromites form only in slags with high
basicity, when the silicon content in the slag is low; the
slag is saturated with CaO.
Acknowledgment
The authors would like to thank mag. Alojz Rozman
and Metal Ravne d.o.o. for financing this research
project.
5 REFERENCES
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Microstructure Evolution in SAF 2507 Super Duplex Stainless Steel,
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2 S. Sun, P. Uguccioni, M. Bryant, M. Ackroyd, Chromium control in
the EAF during stainless steelmaking, Proc. of the 55th Electric
Furnace Chicago, 1997, 297–302
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Materiali in tehnologije / Materials and technology 48 (2014) 5, 753–756 755
Figure 5: XRD spectrum of a slag sample after FeSi addition
Slika 5: XRD-spekter vzorca `lindre po dodatku FeSiFigure 4: XRD spectrum of a slag sample before the FeSi addition
Slika 4: XRD-spekter vzorca `lindre pred dodatkom FeSi
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Of The Superaustenitic Stainless Steel AISI 904L, Mater. Tehnol., 48
(2014) 1, 137–140
4 B. Arh, F. Tehovnik, The Oxidation and Reduction of Chromium,
Metalurgija, 50 (2011), 179–182
5 B. Arh, F. Tehovnik, The Oxidation And Reduction of Chromium
During the Elaboration of Stainless Steels in an Electric Arc Furnace,
Mater. Tehnol., 41 (2007) 5, 203–211
6 Y. Xiao, L. Holappa, M. A. Reuter, Oxidation State and Activities of
Chromium Oxides in CaO-SiO2-CrOx Slag System, Met. and Mater.
Trans. B, 33 (2002), 595–603
7 K. Morita, N. Sano, Activity Of Chromium Oxide in CaO-SiO2
Based Slags At 1873 K, Proc. of the VII International Conference on
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8 J. H. Park, I. H. Jung, S. B. Lee, Phase Diagram Study for the CaO-
SiO2-Cr2O3-5 mass.%MgO-10 mass.%MnO System, Met. Mater.
Int., 15 (2009), 677–681
9 S. Mostafaee, A Study of EAF High-Chromium Stainless Steel-
making Slags Characteristics and Foamability, PhD. Thesis, Stock-
holm, 2011
10 G. J. Albertsson, Investigations of Stabilization of Cr in Spinel Phase
in Chromium-Containing Slags, Licentiate Thesis, Stockholm, 2011
11 N. Sano, Reduction Of Chromium Oxide In Stainless Steel Slags,
Proc. of 10th Int. Ferroalloys Congr. INFACON X Transformation
through Technol., Cape Town, 2004, 670–677
12 E. Shibata, S. Egawa, T. Nakamura, Reduction Behavior of Chro-
mium Oxide in Molten Slag Using Aluminum, Ferrosilicon and
Graphite, SIJ Int., 42 (2002), 609–613
13 J. Björkvall, S. Ångström, L. Kallin, Reduction Of Chromium Oxide
Containing Slags Using CaC2, Proc. of the VII Int. Conf. Molten
Slags Fluxes Salts, Johannesbourg, 2004, 663–670
14 X. Hu, H. Wang, L. Teng, S. Seetharaman, Direct Chromium Alloy-
ing by Chromite Ore with the Presence of Metallic Iron, J. Min.
Metall. Sect. B Metall., 49 (2013), 207–215
15 R. I. L. Guthrie, M. Isac, Z. Lin, A Fluid Dynamics Simulation Of
Chromium Recovery From Aod Slags During Reduction With
Ferrosilicon Additions, Proc. of the VII Int. Conf. Molten Slags
Fluxes Salts, Johannesbourg, 2004, 465–472
16 H. Cabrera-Real et al., Effect of MgO and CaO/SiO2 on the Immo-
bilization of Chromium in Synthetic Slags, J. Mater. Cycles Waste
Manag., 14 (2012), 317–324
17 M. Hino, K. Higuchi, T. Nagasaka, S. Ban-Ya, Phase Equilibria and
Activities of the Constituents in FeO·Cr2O3- MgO·Cr2O3 Spinel
Solid Solution Saturated with Cr203, ISIJ Int., 34 (1994), 739–745
18 S. Jo, B. O. Song, S. Kim, Thermodynamics on the Formation of
Spinel (MgO · Al2O3) Inclusion in Liquid Iron Containing Chro-
mium, Met. and Mater. Trans. B, 33 (2002), 703–709
19 M. Guo, S. Parada, S. Smets, P. T. Jones, J. Van Dyck, B. Blanpain,
P. Wollants, Laboratory Study of the Interaction Mechanisms Bet-
ween Magnesia-Chromite Refractories and Al2O3-rich VOD slags,
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20 M. Guo, P. T. Jones, S. Parada, E. Boydens, J. Van Dyck, B. Blan-
pain, P. Wollants, Degradation Mechanisms of Magnesia-Chromite
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Conditions, J. Eur. Ceram. Soc., 26 (2006), 3831–3843
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756 Materiali in tehnologije / Materials and technology 48 (2014) 5, 753–756
K. MOHAJERSHOJAEI, F. DADASHIAN: RECYCLING OF JUTE WASTES USING PULPZYME ENZYME
RECYCLING OF JUTE WASTES USING PULPZYME ENZYME
RECIKLIRANJE ODPADKOV JUTE Z UPORABO ENCIMA PULPZIMA
Khashayar Mohajershojaei1, Fatemeh Dadashian2
1Department of Polymer Engineering and Color Technology, Amirkabir University of Technology, Tehran, Iran
2Department of Textile Engineering, Amirkabir University of Technology, Tehran, Iran
khashayar045@yahoo.com, dadashia@aut.ac.ir
Prejem rokopisa – received: 2012-12-22; sprejem za objavo – accepted for publication: 2013-11-06
In this paper, enzymatic treatment of jute wastes using pulpzyme was studied. The jute wastes from the machine-made
carpet-production factories were used as a model. The effects of several parameters such as enzyme concentration, pH and time
on the recycling process were evaluated. The optimum enzyme concentration, reaction time and pH for the recycling of jute
wastes were 1.5 %, 2 h and 8, respectively.
The results showed that the enzymatic process using pulpzyme was an effective method to hydrolyse cellulosic chains, shorten
cellulosic fibers such as jute and decrease its moisture regain (%). The products obtained from the enzymatic process using
pulpzyme are suitable raw materials for paper-making processes due to their length range between 0 mm to 4 mm.
Keywords: recycling process, jute wastes, pulpzyme enzyme, mass loss, length reduction, fiber shortening
V ~lanku je prikazana encimska obdelava odpadkov jute z uporabo pulpzima. Kot model so bili uporabljeni odpadki jute pri
strojni izdelavi preprog. Ocenjen je bil u~inek ve~ parametrov, kot so koncentracija encima, pH in ~as za postopek recikliranja.
Optimalna koncentracija encima, ~as reakcije in pH pri recikliranju odpadkov jute so bili 1,5 %, 2 h in 8.
Rezultati so pokazali, da je encimski postopek z uporabo pulpzima u~inkovita metoda za hidrolizo celuloznih verig, skraj{anje
celuloznih vlaken jute in zmanj{anje njene ponovne navla`enosti. Dobljeni produkti iz encimskega postopka z uporabo
pulpzima so zaradi njihove dol`ine od 0 mm do 4 mm primerni za izdelavo papirja.
Klju~ne besede: postopek recikliranja, odpadki jute, encim pulpzim, zmanj{anje mase, zmanj{anje dol`ine, skraj{anje vlaken
1 INTRODUCTION
Waste management is an important issue in various
industries. Nowadays, environmental concerns, ecologi-
cal and economic considerations constitute the driving
force for the waste management in the textile industry. In
these cases, managing wastes involves modifying the old
systems, developing new processes to limit, optimize and
process waste materials and finding usage for post-con-
sumer textile wastes.1 Managing wastes in a correct way
will lead to saving the energy and cost, reducing the
landfill usage and solving the present environmental and
ecological problems.2
According to the importance of the recycling proces-
ses in developed countries, different physical, chemical
or biological methods are used to recycle various wastes
from different industries. Cellulosic wastes such as
cotton, viscose, lyocell and jute constitute a major part of
the textile industries such as the machine-made carpet
production. Nowadays, different methods such as chemi-
cal and biological ones are used to recycle these wastes.
Some of the chemical processes proceed slowly; there-
fore, catalysts (especially enzymes) are needed to en-
hance the rate of chemical reactions. In these cases, a
small quantity of an enzyme is able to react with a large
amount of a substrate in a mild condition.3
Hemmpel4 used cellulase enzyme in different pH
conditions to modify woven and knitted cellulose fabrics.
The results showed that cellulose enzyme has a signifi-
cant effect on cellulose fabrics when pH is lower than 5,
by hydrolyzing the cellulose bonds.
Cellulase enzyme was used for bio-polishing of
cotton, viscose and lyocell fabrics by Garret.5 He con-
cluded that cotton, viscose and lyocell fabrics could be
modified using cellulase enzyme due to the surface-fiber
removal.
The effect of an enzymatic treatment on the fine
structure of cellulosic fibers and the properties of cellu-
lose dissolved in aqueous 7.6 % NaOH and ionic liquid
were analyzed by F. Dadashian6 and P. Rosenberg et al.7
It can be concluded from the SEM photographs that
cellulase enzyme has a significant effect on modifying
the fine structure of cellulosic fibers.
G. Buschle-Diller et al.8 concluded that the cellulases
from Trichoderma viride have a significant effect on the
pore structure of different types of bead cellulose.
A literature review showed that the recycling of jute
wastes using pulpzyme was not studied. In this paper,
enzymatic recycling of jute wastes using pulpzyme enzy-
me was studied. The jute wastes from machine-made
carpet-production factories were used as a model. The
effects of several parameters such as enzyme concen-
tration, pH and time on the recycling of the jute wastes
were evaluated with respect to the mass- and length-loss
fractions.
Materiali in tehnologije / Materials and technology 48 (2014) 5, 757–760 757
UDK 658.567.3: 677.13.004.8 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(5)757(2014)
2 EXPERIMENTAL WORK
2.1 Materials
The jute waste and pulpzyme enzyme were obtained
from the Novo Nordisk and Akij (Bangladesh) Compa-
nies, respectively. All the other chemicals were of an
analytical grade and purchased from Merck (Germany).
The main constituents of the jute fiber are repre-
sented in Table 1.9
Table 1: Contents of jute fibers9
Tabela 1: Sestava vlaken jute9
Content Fraction (%)
Ash 3.04
Benzene alcohol 2.72
Lignin 21.61
Acid soluble lignin 0.95
Alpha cellulose 42.91
Beta cellulose 20
Others 8.77
2.2 Jute recycling
Experiments were carried out in a batch-mode reactor
with the total capacity of 250 ml. The recycling of jute
wastes was performed using a 100 mL solution contain-
ing a specified amount of jute wastes (5 g) using pul-
pzyme enzyme. The solution pH was adjusted using a
phosphate buffer.8,9 The samples were withdrawn from
the sample point at certain time intervals and analyzed
for mass and length losses.
The mass- and length-loss fractions were checked
and controlled by measuring the mass and length of the
jute fiber before and after the enzymatic process.
In this study an light projection microscope (400 X)
and a Philips scanning electron microscope (690 X) were
used to measure the fiber length and evaluate the surface
morphology of the enzyme-treated and untreated jute
wastes, respectively.
The effects of the enzyme concentration (0.5–2 %)
on the mass- and length-loss fractions of the samples
were investigated by treating 5 g of jute wastes at the pH
of 8 at 55 °C for 2 h.
The effects of the pH (3–9) on the mass- and length-
loss fractions of the samples were investigated by treat-
ing 5 g of jute wastes with a 1.5 % enzyme concentration
at 55 °C for 2 h.
The effect of the time of treatment (0–3 h) on the
mass- and length-loss fractions of the samples were inve-
stigated by treating 5 g of jute wastes with a 0.5 %
enzyme concentration at the pH of 8 at 55 °C.
The mass-loss fractions of the samples were
measured by weighting jute wastes before and after the
enzymatic treatments.
The mass-loss fraction was defined as follows:
Mass-loss fraction /% = (A2 – A1)/A1 × 100 (1)
where A1 and A2 stand for the dry mass and conditioned
mass of the samples before and after the enzymatic
treatment, respectively.
The moisture regain of the samples was measured
with the gravimetric method. The moisture regain was
defined as follows:
Moisture regain /% = (W2 – W1)/W1 × 100 (2)
where W1 and W2 stand for the dry mass and conditioned
mass of the samples, respectively. The FTIR spectra
were obtained using a Nicolet Magna IR spectro-
photometer equipped with a microscope.
3 RESULTS AND DISCUSSION
3.1 Enzyme concentration
The mass-loss and length-reduction fractions of jute
samples for different enzyme concentrations are shown
in Figure 1. According to the data from Figure 1, with
the increasing enzyme concentration, the mas-loss and
length-reduction fractions increased gradually because of
the existing excess enzyme molecules relative to the
fixed amount of jute. With the 1.5 % enzyme concentra-
tion, the optimum mass-loss (almost 46 %) and length-
reduction (almost 73.45 %) fractions were obtained;
therefore, the optimum enzyme concentration for the
recycling of jute was 1.5 %.
3.2 pH
The mass-loss and length-reduction fractions of jute
samples at different pH values are shown in Figure
2.The results show that the pH significantly influenced
the pulpzyme action during jute recycling. The mass-loss
and length-reduction fractions were found to improve
with an increase in the aqueous-phase pH up to the value
of 8 and, thereafter, an increase in the aqueous-phase pH
from 8.0 to 9.0 caused the efficacy of the enzymatic
recycling process to decrease. The aqueous-phase pH of
8.0 had a significant effect on the rate of the mass-loss
and length-reduction fractions compared to the other pH
K. MOHAJERSHOJAEI, F. DADASHIAN: RECYCLING OF JUTE WASTES USING PULPZYME ENZYME
758 Materiali in tehnologije / Materials and technology 48 (2014) 5, 757–760
Figure 1: Effect of enzyme concentration on: a) mass-loss fraction, b)
length-reduction fraction of the samples
Slika 1: Vpliv koncentracije encima na: a) dele` zmanj{anja mase, b)
dele` zmanj{anja dol`ine vzorcev
conditions. Thus, the aqueous-phase pH plays a signifi-
cant role in enzymatic reactions. Moreover, the pH-acti-
vity relationship of any given enzyme depends on the
acid-base behavior of the enzyme and the substrate (the
jute waste) as well as on many other factors that are
usually difficult to analyze quantitatively.10
3.3 Time of treatment
The mass-loss and length-reduction fractions of the
jute samples at different times are given in Table 2. The
results show that the time of treatment significantly
influenced the pulpzyme action during jute recycling.
The mass-loss and length-reduction fractionswere found
to improve with an increase in the time up to 2 h and,
thereafter, an increase in the time of the enzymatic
treatment from 2 h to 4 h was without any significant
changes.
After the 2 h enzyme reaction, the optimum mass-
loss (almost 46 %) and length-reduction fractions (al-
most 73.45 %) of a jute sample were obtained; therefore,
the optimum time for the enzymatic treatment was 2 h.10
Table 2: Effect of the time of treatment on mass-loss fraction and
length-reduction fractions of the samples for the optimum enzyme
concentration and pH
Tabela 2: Vpliv ~asa obdelave na dele` zmanj{anja mase in dele`
zmanj{anja dol`ine vzorcev pri optimalni koncentraciji encima in pH
Time (h) 1 2 3 4
Mass loss (%) 27.6 46 48.3 50
Length reduction (%) 65.76 73.45 78.5 81.3
3.4 Effect of the enzymatic process on the moisture re-
gain (%)
The moisture regains of the samples before and after
the enzymatic treatment were measured according to
equation (2). According to the data from Table 3, the en-
zyme-treated jute has a lower (22.37 %) moisture regain
than the untreated jute in the optimum condition. The
results show that the enzymatic treatment of jute has a
specific effect on increasing the crystallinity.6,11
Table 3: Moisture regain of jute-fiber wastes after and before the
enzymatic treatment in the optimum conditions
Tabela 3: Ponovno navla`enje vlaken odpadne jute pred obdelavo z
encimi in po njej v optimalnih razmerah
Sample Conditionedmass (g)
Dry mass
(g)
Moisture
regain (%)
Raw jute 0.63 0.54 14.3
Enzyme-treated jute 0.63 0.57 11.1
3.5 Microscopic structure
The effect of pulpzyme enzyme on the macrostruc-
tures of the samples was studied using SEM and light
K. MOHAJERSHOJAEI, F. DADASHIAN: RECYCLING OF JUTE WASTES USING PULPZYME ENZYME
Materiali in tehnologije / Materials and technology 48 (2014) 5, 757–760 759
Figure 2: Effect of pH on: a) mass-loss fraction, b) length-reduction
fraction of the samples
Slika 2: Vpliv pH na: a) dele` zmanj{anja mase, b) dele` zmanj{anja
dol`ine vzorcev
Figure 4: SEM photographs of: a) fiber wastes enzyme treated with
the optimum enzyme concentration and b) untreated jute-fiber wastes
Slika 4: SEM-posnetka povr{ine vlakna odpadne jute: a) obdelanega z
encimom pri optimalni koncentraciji encima in b) neobdelanega
vlakna odpadne jute
Figure 3: Light micrographs of: a) untreated jute-fibre wastes, b)
jute-fibre wastes treated with enzyme in optimum concentration
Slika 3: Svetlobni posnetki a) neobdelanega vlakna odpadne jute, b)
vlakna odpadne jute, obdelanega z encimom pri optimalni koncentra-
ciji encima
microscopic photographs. The light microscopic and
scanning electron microscopic (SEM) photographs of the
samples before and after the enzymatic processes are
shown in Figures 3 and 4. The results show that pulp-
zyme enzyme has a significant effect on the destruction
of the molecular structure of jute in the longitudinal
direction due to the hydrolyzing reaction. This phenome-
non has a significant effect on the shortening of the jute
samples.
4 CONCLUSION
In this study, enzymatic treatments were used to
modify jute wastes. The effect of several parameters such
as enzyme concentration, pH and time on the recycling
of jute wastes was evaluated.
The results showed that the enzymatic treatment has
a significant effect on the mass loss and shortening of the
jute samples by hydrolyzing cellulosic chains.
Moreover, the enzymatic treatment leads to a de-
crease in the moisture regain of the samples due to the
increasing crystallinity. It can be concluded from the
SEM and light microscopic photographs that pulpzyme
enzyme has a significant effect on the destruction of the
molecular structure of jute in the longitudinal direction
due to the hydrolyzing reaction. Finally, it leads to the
shortening of the jute samples.
According to the obtained data, the optimum enzyme
concentration, reaction time and pH for recycling jute
wastes were 1.5 %, 2 h and 8, respectively. The obtained
samples are suitable for a usage in the paper-making
industries due to their longitudinal distribution.
5 REFERENCES
1 M. J. Cornell, Getting to zero Wastes, [cited fall 2007], Available
from World Wide Web: http://www.toenall.org
2 L. Wang, Y. Z. Zhang, H. Yang, Carbohydrate Research, 339 (2004),
819–824
3 Enzyme kinetics, Thomas Jefferson University, Maintained by AISR
Education Services (EdServices@jeffline.tju.edu)
4 W. H. Hemmpel, International Textile Bulletin Dyeing/ Printing/Fini-
shing, 37 (1992) 3, 1–5
5 A. S. Garret, D. M. Cedroni, In AATCC Book of Paper, AATCC
International Conference & Exhibition, Atlanta, 1992
6 F. Dadashian, Changes in the Fine Structure of Lyocell Fibers during
Enzymatic Degradation, Proceedings of the World Textile Con-
ference, 2ndAutex Conference, Bruges, Belgium, 2002, 464–473
7 P. Rosenberg, T. Budtova, M. Rom, P. Fardim, Effect of Enzymatic
Treatment on Solubility of Cellulose in 7.6 % NaOH-Water and Ionic
Liquid, ACS Symposium Series, Chapter 12, 2010, 213–226
8 G. Buschle-Diller, C. Fanter, F. Loth, Cellulose, 2 (1996) 3, 179–203
9 T. C. Ranjan, Handbook of Jute, vol. 3, Oxford and IBH, New Delhi
1973
10 K. P. Katuri, S. V. Mohen, S. Sridhar, B. R. Pati, P. N. Sarma, Water
Resource, 43 (2009), 3647–365
11 K. M. Shojaei, F. Dadashian, M. Montazer, Applied Biochemistry
and Biotechnology, 166 (2011) 3, 744–52
K. MOHAJERSHOJAEI, F. DADASHIAN: RECYCLING OF JUTE WASTES USING PULPZYME ENZYME
760 Materiali in tehnologije / Materials and technology 48 (2014) 5, 757–760
B. R. GNANASUNDARAM, M. NATARAJAN: INFLUENCES OF THE HEAT INPUT ON A 2205 DUPLEX STAINLESS STEEL WELD
INFLUENCES OF THE HEAT INPUT ON A 2205 DUPLEX
STAINLESS STEEL WELD
VPLIV VNOSA TOPLOTE V ZVAR DUPLEKSNEGA NERJAVNEGA
JEKLA 2205
Bansal Rajkumar Gnanasundaram1, Murugan Natarajan2
1United Institute of Technology, Coimbatore, Tamilnadu, India
2Coimbatore Institute of Technology, Coimbatore, Tamilnadu, India
bansalgnanasundaram@gmail.com
Prejem rokopisa – received: 2013-02-05; sprejem za objavo – accepted for publication: 2013-11-19
In arc welding, the energy is transferred from the electrode to the base metal via an electric arc. The energy transferred per unit
length is measured as the heat input. The heat input is an important factor that influences the cooling rate and affects the
metallurgical and mechanical properties of welds. In this study, the flux-cored arc-welding process was used to join duplex
stainless steel. The experiment was conducted on the basis of a three-factor, five-level central composite rotatable design using
the full-replication technique. A mathematical model was developed for the heat input. The effects of the welding-process
parameters on the heat input are discussed. Metallographic techniques were employed to study the microstructure produced at
low, medium, high and optimum heat-input conditions.
Key words: heat input, FCAW, duplex stainless steel, microstructure
Pri oblo~nem varjenju se energija prena{a iz elektrode na osnovno kovino z elektri~nim oblokom. Energija, prenesena na enoto
dol`ine, se meri kot vnos toplote. Vnos toplote je pomemben faktor, ki vpliva na hitrost ohlajanja, na metalur{ke in mehanske
lastnosti zvara. V tej {tudiji je bilo za spajanje dupleksnega nerjavnega jekla uporabljeno varjenje pod pra{kom. Eksperiment je
bil izvr{en na osnovi treh faktorjev in petnivojske centralno vrtljive izvedbe s polno mo`nostjo ponovitve. Za vnos toplote je bil
razvit matemati~ni model. Obravnavan je bil vpliv parametrov varilnega procesa na vnos toplote. Za {tudij mikrostrukture,
nastale pri nizkem, srednjem, velikem in optimalnem vnosu toplote, so bili uporabljeni metalografski postopki.
Klju~ne besede: vnos toplote, FCAW, dupleksno nerjavno jeklo, mikrostruktura
1 INTRODUCTION
The flux-cored arc welding (FCAW) is widely used
by industries because it produces welds with better and
more consistent mechanical properties and fewer weld
defects, it is a high-deposition-rate process and suitable
for stainless steel. Generally, FCAW produces a stronger
weldment than SMAW at room temperature.1 Duplex
stainless steel (DSS) is widely used in industrial applica-
tions due to its excellent corrosion resistance and its
strength that is higher compared to types 316 and 317 of
austenite stainless steel.2 It is a combination of 50 %
austenitic and 50 % ferritic steels. In the present study,
2205 DSS was used as the base metal for the experiment
and a diameter 1.2 mm (E2209T1-4/1) filler wire was
used for depositing the metal. For the shielding purpose,
a combination of 75 % argon plus 25 % CO2 was used as
the welding gas for horizontal-position welding. 2205
DSS was procured from Outokumpu Stainless AB, Swe-
den.
2 EXPERIMENTAL SETUP
The FCAW setup available at the Welding Research
Centre of Coimbatore Institute of Technology (CIT),
Coimbatore, India was used to conduct the experiments.
The welding setup consists of a power source, namely,
INDARC 400 MMR, a unit feeding the filler wire with
the shielding-gas-flow control, a welding gun and a
welding manipulator that helps to deposit the filler metal
on the desired area as shown in Figure 1.
The selection of significant FCAW process
parameters helps us to get the desired weld-bead quality.3
In accordance with the available literature, the welding
current (I), the welding speed (S) and the open-circuit
voltage (OCV) were taken as the significant process
parameters. The trial runs were conducted by varying
Materiali in tehnologije / Materials and technology 48 (2014) 5, 761–763 761
UDK 621.791:669.14.018.8 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(5)761(2014)
Figure 1: FCAW experimental setup
Slika 1: Eksperimentalni sestav FCAW
one process parameter while keeping the other two
parameters constant. The working range was selected, by
visual inspection, on the basis of the appearance of the
weld bead with respect to how smooth and continuous it
was, and with respect to the absence of defects like
porosity, undercut, etc. The welding parameters and their
levels are given in Table 1. The experiment was
conducted as per the design matrix and 20 runs were
made by varying the process variables as shown in Table
2.
Table 1: Welding parameters and their factor levels
Tabela 1: Parametri varjenja in faktorji njihovih nivojev
S. No. Parameter/unit
Factor level
–1.682 –1 0 1 1.682
1 Weldingcurrent (A) 170 182 200 218 230
2 Welding speed(cm/min) 25 27 31 35 37
3 Open-circuitvoltage (V) 28 30 32 34 36
Table 2: Design matrix and the observed heat-input values
Tabela 2: Postavitev matrice in opa`ene vrednosti vnosa toplote
Sp. No. I/A S/(cm/min) OCV/V U/V
HI/
(kJ/cm)
1 –1 –1 –1 28.7 9.87
2 1 –1 –1 24.35 10.03
3 –1 1 –1 27.75 7.36
4 1 1 –1 27.7 8.80
5 –1 –1 1 31.55 10.85
6 1 –1 1 31.35 12.91
7 –1 1 1 30.9 8.19
8 1 1 1 29.35 9.32
9 –1.682 0 0 31.25 8.74
10 1.682 0 0 31.8 12.03
11 0 –1.682 0 32 13.06
12 0 1.682 0 32.55 8.97
13 0 0 –1.682 24.7 8.13
14 0 0 1.682 34.5 11.35
15 0 0 0 32.85 10.81
16 0 0 0 32.25 10.61
17 0 0 0 31.75 10.45
18 0 0 0 29.4 9.67
19 0 0 0 30.6 10.07
20 0 0 0 27.45 9.03
3 DEVELOPMENT OF A REGRESSION MODEL
FOR THE HEAT INPUT
The heat input could not be measured directly; how-
ever, with the help of the welding current, the arc voltage
and the welding speed it could be calculated. The arc
voltage was measured with the voltmeter of the welding
transformer, the welding current was determined on the
basis of the wire feed rate and the welding speed on the
basis of the movement of the table. The heat input was
calculated as the ratio of the power to the velocity of the
heat source:
HI = 60 UI /1000 S (1)
where HI = heat input (kJ/cm), U = arc voltage (V), I =
current (A), S = travel speed (cm/min) and  = arc effi-
ciency accounting for the heat dissipation to the
surrounding as a result of the convection and radiation
(0.85). The observed heat-input values are shown in
Table 2.
The response function representing the parameters is
expressed with the following equation:
Y = f (X1, X2, X3) (2)
where Y = heat input, X1 = current (I), X2 = speed (S)
and X3 = open circuit voltage (V).
The second-order polynomial equation represents the
response for K factors given in equation:4
Y b b X b X X b Xi i
I
K
ij i j
ij
K
ii i
I
i j
K
= + + +
= = =
≠
∑ ∑ ∑0
1 1
2
1
(3)
For the three factors, the above polynomial equation
can be expressed as:
Y = b0 + b1 I + b2 S+ b3 V+ b12 IS + b13 IV+ b23 SV+
+ b11 I²+ b22 S²+ b33 V² (4)
where b0 = free term, coefficients b1, b2, b3 = linear
terms, coefficients b12, b13, b23 = interaction terms and
coefficients b11, b22, b33 = quadratic terms. The deve-
loped regression model in the coded form is given
below; adequacy was checked with the ANOVA method.
Heat input:
HI = 10.213 + 0.756I – 1.235S + 0.778V –
– 0.294V2 – 0.314SV (5)
4 RESULTS
Figure 2 depicts the effects of the process parameters
on the heat input. The welding speed shows a significant
effect on the heat input. The heat input increases up to
12.67 kJ/cm with the lower welding speed of 25 cm/min
and it gradually falls down to 8.52 kJ/cm at the higher
B. R. GNANASUNDARAM, M. NATARAJAN: INFLUENCES OF THE HEAT INPUT ON A 2205 DUPLEX STAINLESS STEEL WELD
762 Materiali in tehnologije / Materials and technology 48 (2014) 5, 761–763
Figure 2: Effects of process parameters on the heat input
Slika 2: Vpliv parametrov procesa na vnos toplote
welding speed of 37 cm/min. At the lower welding
speed, the heat input per unit length of the weld is higher
resulting in large heat-affected zones and causing a se-
vere distortion. The heat input increases from 8.01 kJ/cm
to 11.24 kJ/cm with the increasing welding current and
open-circuit voltage. Obviously, the higher voltage and
the higher current increase the discharge and elevate the
heat input.
In Figure 3a, the ferrite phase is shown in white
color. The presence of a columnar dendrite structure
represents a faster cooling rate due to the low heat-input
condition.5 It is evident from Figure 3b that a vermicular
structure is observed in the weld metal at the magni-
fication of 400-times. The delta ferrite is pale yellow, the
austenitic regions are greenish blue, and the final-soli-
dification regions are bluish. In Figure 3c, the micro-
structure reveals the presence of ferrite (white) and
austenite (dark). The coarse austenite grains found in the
microstructure may be due to the higher heat input
during the welding.
5 CONCLUSION
At the higher welding speed, the heat input per unit
length of the weld decreases.
A higher heat input reduces the weld quality and in-
creases the distortion.
The microstructure study reveals the morphology of
2205 DSS welds at various heat-input conditions.
6 REFERENCES
1 Z. Zhang, A. W. Marshall, G. B. Holloway, Flux cored arc welding:
the high productivity welding process for P91 steels, Proceedings of
the 3rd conference on advances in materials technology for fossil
power plants, The Institute of Materials, University of Wales
Swansea, London, 2001, 267–282
2 I. Alvarez-Armas, Duplex Stainless Steels: Brief History and Some
Recent Alloys, Recent Patents on Mechanical Engineering, 1 (2008),
51–57
3 K. Y. Benyounis, A. G. Olabi, Optimization of different welding
processes using statistical and numerical approaches – A reference
guide, Adva. Engg. Software, 39 (2008), 483–496
4 K. M. Carley, Response surface methodology, CASOS Technical Re-
port, Carnegie Mellon University, 2004
5 J. Elmer, S. M. Allen, T. W. Eagar, Microstructural development
during solidification of stainless steel alloys, Metallurgical Tran-
sactions A, 20A (1989), 2117–2131
B. R. GNANASUNDARAM, M. NATARAJAN: INFLUENCES OF THE HEAT INPUT ON A 2205 DUPLEX STAINLESS STEEL WELD
Materiali in tehnologije / Materials and technology 48 (2014) 5, 761–763 763
Figure 3: Microstructure of the weld metal zone at different conditions: a) low heat input (HI = 7.36 kJ/cm), b) medium heat input (HI = 10.45
kJ/cm), c) high heat input (HI = 13.06 kJ/cm); magnification: 400-times
Slika 3: Mikrostruktura zvara v razli~nih razmerah: a) nizek vnos toplote (HI = 7,36 kJ/cm), b) srednji vnos toplote (HI = 10,45 kJ/cm), c) visok
vnos toplote (HI = 13,06 kJ/cm); pove~ava 400-kratna

I. GUNES: TRIBOLOGICAL BEHAVIOR AND CHARACTERIZATION OF BORIDED COLD-WORK TOOL STEEL
TRIBOLOGICAL BEHAVIOR AND CHARACTERIZATION
OF BORIDED COLD-WORK TOOL STEEL
TRIBOLO[KO VEDENJE IN KARAKTERIZACIJA BORIRANEGA
ORODNEGA JEKLA ZA DELO V HLADNEM
Ibrahim Gunes
Department of Metallurgical and Materials Engineering, Faculty of Technology, Afyon Kocatepe University, 03200 Afyonkarahisar, Turkey
igunes@aku.edu.tr
Prejem rokopisa – received: 2013-08-20; sprejem za objavo – accepted for publication: 2013-11-19
In the present study, tribological and characterization properties of the borides formed on cold-work tool steel were investigated.
Boriding was performed in a solid medium consisting of Ekabor-II powders at 850 °C and 950 °C for 2 h and 6 h. The boride
layer was characterized with light microscopy, X-ray diffraction technique and a micro-Vickers hardness tester. An X-ray
diffraction analysis of the boride layers on the surface of the steel revealed the existence of FeB, Fe2B, CrB, Cr2B and MoB
compounds. Depending on the chemical compositions of the substrates and the boriding time, the boride-layer thickness on the
surface of the steel ranged from 13.14 μm to 120.82 μm. The hardness of the boride compounds formed on the surface of the
steel ranged from 1806 HV0.05 to 2342 HV0.05, whereas the Vickers-hardness value of the untreated steel was 428 HV0.05. The
wear tests were carried out in a ball-disc arrangement under a dry friction condition at room temperature with an applied load of
10 N and with a sliding speed of 0.25 m/s at a sliding distance of 1000 m. The wear surfaces of the steel were analyzed using
SEM microscopy and X-ray energy-dispersive spectroscopy (EDS). It was observed that the wear rate of the unborided and
borided cold-work tool steels ranged from 11.28 mm3/(N m) to 116.54 mm3/(N m).
Keywords: cold-work tool steel, boriding, microhardness, wear rate, friction coefficient
V tej {tudiji so bile preiskovane tribolo{ke in druge zna~ilne lastnosti boridov, nastalih na orodnem jeklu za delo v hladnem.
Boriranje je bilo izvr{eno v trdnem mediju, ki ga je sestavljal prah Ekabor-II pri 850 °C in 950 °C v 2 h in 6 h. Boriran sloj je bil
preiskan s svetlobno mikroskopijo, rentgensko difrakcijo in izmerjena je bila Vickersova mikrotrdota. Rentgenska difrakcija
boriranega sloja na povr{ini jekla je potrdila prisotnost spojin FeB, Fe2B, CrB, Cr2B in MoB. Odvisno od kemijske sestave
podlage in ~asa boriranja je bila debelina boriranega sloja na povr{ini jekla med 13,14 μm in 120,82 μm. Trdota boridov,
nastalih na povr{ini jekla, je bila v razponu od 1806 HV0,05 do 2342 HV0,05, medtem ko je bila Vickersova trdota neobdelanega
jekla 428 HV0,05. Preizkusi obrabe so bili izvr{eni na napravi krogla-disk pri suhem trenju pri sobni temperaturi z uporabljeno
obte`bo 10 N, hitrostjo drsenja 0,25 m/s in pri razdalji drsenja 1000 m. Povr{ina z obrabo je bila analizirana s SEM-mikro-
skopijo in energijsko disperzijsko spektroskopijo rentgenskih `arkov (EDS). Ugotovljeno je bilo, da je hitrost obrabe
neboriranega in boriranega orodnega jekla za delo v hladnem v razponu od 11,28 mm3/(N m) do 116,54 mm3/(N m).
Klju~ne besede: orodno jeklo za delo v hladnem, boriranje, mikrotrdota, hitrost obrabe, koeficient trenja
1 INTRODUCTION
Industrial boriding processes can be applied to a wide
range of steel alloys including carbon steel, low and
high-alloy steel, tool steel and stainless steel. Cold-work
tool steels have a wide range of applications. Therefore,
there has been an extensive research on the development
of surface-treatment processes to improve the wear
resistance, corrosion and oxidation resistance of the
cold-work tool steels for high-temperature and high-
pressure applications in recent years.1–7 Boriding is a
thermochemical surface treatment, in which boron atoms
diffuse into the surface of a workpiece to form borides
with the base material.8,9 The main advantages of this
technique are a high resistance to abrasion wear and a
high oxidation resistance when compared with the other
conventional surface treatments. The thermal diffusion
treatments of boron compounds used to form iron
borides typically require the process temperatures of 700
°C and 1000 °C. The process can be carried out in a
solid, liquid, gaseous or plasma medium.10–14
In recent years, the boriding treatment has been used
for a wide range of applications in industries such as the
manufacture of machine parts for plastics, food pro-
cessing, packaging and tooling, as well as pumps and
hydraulic machine parts, crankshafts, rolls and heavy
gears, motor and car constructions, cold- and hot-work-
ing dies and cutting tools. The wear behavior of borided
steels has been evaluated by a number of investiga-
tors.15–20 However, there is no information about the
friction and wear behavior of the borided cold-work tool
steel. The main objective of this study was to investigate
the friction and wear behavior of the borided cold-work
tool steel. Structural and tribological properties were
investigated using light microscopy, XRD, SEM, EDS,
microhardness tests and a ball-on-disc tribotester.
2 EXPERIMENTAL METHOD
2.1 Boriding and characterization
The high-alloy cold-work tool steel essentially
contained mass fractions 0.90 % C, 0.50 % Mn, 7.80 %
Cr, 2.50 % Mo and 0.50 % V. The test specimens were
cut into the cylinders with the dimensions of ø 25 mm ×
10 mm, ground up to 1200 G and polished using a dia-
Materiali in tehnologije / Materials and technology 48 (2014) 5, 765–769 765
UDK 539.92:669.14:621.793/.795 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(5)765(2014)
mond solution. The boriding heat treatment was carried
out in a solid medium containing an Ekabor-II powder
mixture placed in an electrical-resistance furnace
operated at the temperatures of 850 °C and 950 °C for
2 h and 6 h under atmospheric pressure. Following the
completion of the boriding process, the test specimens
were removed from the sealed stainless-steel container
and allowed to cool down in still air. The microstructures
of the polished and etched cross-sections of the speci-
mens were observed with a Nikon MA100 light micro-
scope. The presence of the borides formed in the coating
layer was confirmed by means of X-ray diffraction
equipment (Shimadzu XRD 6000) using Cu K radia-
tion. The hardness measurements of the boride layer for
each steel type and the unborided steel substrate were
made on the cross-sections using a Shimadzu HMV-2
Vickers indenter with a load 50 g.
2.2 Friction and wear
To perform the friction and wear tests of the borided
samples, a ball-on-disc test device was used. In the wear
tests, WC-Co balls of 8 mm in diameter supplied by H.
C. Starck Ceramics GmbH were used. The errors caused
by a distortion of the surface were eliminated using a
separate abrasion element (WC-Co ball) for each test.
The wear experiments were carried out in a ball-disc
arrangement under a dry friction condition at room
temperature with an applied load of 10 N and with the
sliding speed of 0.25 m/s at a sliding distance of 1000 m.
Before and after each wear test, each sample and the
abrasion element were cleaned with alcohol. After the
tests, the wear volumes of the samples were quantified
by multiplying the cross-sectional areas of the wear by
the width of the wear track obtained from a Taylor-
Hobson Rugosimeter Surtronic 25 device. The wear rate
was calculated with the following formula:
Wk /(mm
3/(N m)) = Wv /M · S (1)
where Wk is the wear rate, Wv is the wear volume, M is
the applied load and S is the sliding distance. The fric-
tion coefficients depending on the sliding distance were
obtained through a friction-coefficient program. The
surface profiles of the wear tracks on the samples and
the surface roughness were measured with the Taylor-
Hobson Rugosimeter Surtronic 25. The worn surfaces
were investigated with scanning electron microscopy
(SEM) and energy-dispersive X-ray spectroscopy
(EDS).
3 RESULTS AND DISCUSSION
3.1 Characterization of boride coatings
Light micrographs of the cross-sections of the
cold-work tool steel borided at the temperatures of 850
°C and 950 °C for 2 h and 6 h are shown in Figure 1. As
can be seen, the borides formed on the cold-work
tool-steel substrate have a smooth morphology due to the
high-alloy content. It was found that the coating/matrix
interface and the matrix could be significantly
distinguished and that the boride layer had a columnar
structure. Depending on the chemical compositions of
the substrates and the boriding time, the boride-layer
thickness on the surface of the steel ranged from 13.14
μm to 120.82 μm.
In this study, the presence of borides was identified
using an XRD analysis (Figure 2). The XRD patterns
show that the boride layer consists of the borides such as
AB and A2B (A = metal: Fe, Cr). The XRD results
showed that the boride layers formed on the steel con-
tained the FeB, Fe2B, CrB, Cr2B and MoB compounds
(Figure 2). The microhardness measurements were
carried out along a line between the surface and the
interior in order to see the variations in the boride-layer
hardness, the transition zone and the matrix, respectively.
The microhardness of the boride layers was measured at
10 different locations at the same distance from the
surface, and the average value was taken as the hardness
result. The microhardness measurements were carried
out on the cross-sections, along a line between the
surface and the interior (Figure 3). The hardness of the
boride compounds formed on the surface of the steel
ranged from 1806 HV0.05 to 2342 HV0.05, whereas the
Vickers-hardness value of the untreated steel was 428
HV0.05. When the hardness of the boride layer is
compared with the matrix, the boride-layer hardness is
approximately four times greater than that of the matrix.
3.2 Friction and wear behavior
Table 1 shows the surface-roughness values of the
borided and unborided cold-work tool steels. The sur-
face-roughness values of the unborided and borided
cold-work tool steels varied from 0.12 to 0.54, as can be
seen in Table 1. It was observed for the cold-work tool
I. GUNES: TRIBOLOGICAL BEHAVIOR AND CHARACTERIZATION OF BORIDED COLD-WORK TOOL STEEL
766 Materiali in tehnologije / Materials and technology 48 (2014) 5, 765–769
Figure 1: Cross-sections of borided cold-work tool steel: a) 850 °C, 2
h, b) 850 °C, 6 h, c) 950 °C, 2 h, d) 950 °C, 6 h
Slika 1: Prerez boriranega orodnega jekla za delo v hladnem: a) 850
°C, 2 h, b) 850 °C, 6 h, c) 950 °C, 2 h, d) 950 °C, 6 h
steel that the surface-roughness values increased with the
boriding treatment and time. Gunes18 studied plasma-
paste-borided AISI 8620 steel and reported that the
surface-roughness values increased with an increase in
the boriding temperature. On the other hand, the friction
coefficients of the unborided and borided cold-work tool
steels varied from 0.38 to 0.65, as can be seen in Table
2. With the boriding treatment, a slight reduction was
observed in the friction coefficients of the borided steels.
I. GUNES: TRIBOLOGICAL BEHAVIOR AND CHARACTERIZATION OF BORIDED COLD-WORK TOOL STEEL
Materiali in tehnologije / Materials and technology 48 (2014) 5, 765–769 767
Figure 2: X-ray diffraction patterns of borided cold-work tool steel: a)
850 °C, 2 h, b) 850 °C, 6 h, c) 950 °C, 2 h, d) 950 °C, 6 h
Slika 2: Rentgenska difrakcija boriranega orodnega jekla za delo v
hladnem: a) 850 °C, 2 h, b) 850 °C, 6 h, c) 950 °C, 2 h, d) 950 °C, 6 h
Figure 4: Wear rate of unborided and borided cold-work tool steels
Slika 4: Obraba neboriranega in boriranega orodnega jekla za delo v
hladnem
Figure 3: Variation of the hardness depth for borided steel
Slika 3: Spreminjanje trdote po globini boriranega jekla
Table 1: Surface roughness values for unborided and borided steels
Tabela 1: Vrednosti za hrapavost povr{ine neboriranega in boriranega
jekla
Unborided Borided
850 °C,2 h 850 °C,6 h 950 °C,2 h 950 °C,6 h
0.12 0.32 0.37 0.43 0.54
Table 2: Friction coefficients for unborided and borided steels
Tabela 2: Koeficient trenja za neborirano in borirano jeklo
Unborided Borided
850 °C,2 h 850 °C,6 h 950 °C,2 h 950 °C, 6 h
0.65 0.38 0.41 0.48 0.56
I. GUNES: TRIBOLOGICAL BEHAVIOR AND CHARACTERIZATION OF BORIDED COLD-WORK TOOL STEEL
768 Materiali in tehnologije / Materials and technology 48 (2014) 5, 765–769
Figure 6: SEM micrographs and EDS analysis of the wear surfaces of borided steel: a) 850 °C, 2 h, b) 850 °C, 6 h, c) 950 °C, 2 h, d) 950 °C, 6 h,
e) EDS
Slika 6: SEM-posnetki in EDS-analiza obrabljene povr{ine na boriranem jeklu: a) 850 °C, 2 h, b) 850 °C, 6 h, c) 950 °C, 2 h, d) 950 °C, 6 h, e)
EDS
Figure 5: a) SEM micrograph and b) EDS analysis of the wear surfaces of unborided steel
Slika 5: a) SEM-posnetek in b) EDS-analiza obrabljenih povr{in na neboriranem jeklu
Figure 4 shows the wear rate of the unborided and
borided cold-work tool steels. The reductions in the wear
rates of the borided steels were observed with respect to
the unborided steels. Due to the toughness of the FeB,
Fe2B, CrB, Cr2B and MoB phases, these steels showed a
larger resistance to wear. The lowest wear rate was
obtained for the steel borided at 950 °C for 6 h, while the
highest wear rate was obtained for the unborided steel.
The wear test results indicated that the wear resistance of
borided steels increased considerably with the boriding
treatment. It is well known that the hardness of the bo-
ride layer plays an important role in the improvement of
the wear resistance. As shown in Figures 3 and 4, the
relationship between the surface microhardness and the
wear resistance of borided samples also confirms that the
wear resistance was improved with an increased hard-
ness. This is in agreement with the reports of the pre-
vious studies.20–23 When the wear rate of the borided steel
is compared with the unborided steel, the wear rate of
the borided steel (950 °C, 6 h) is approximately nine
times lower than that of the unborided steel.
SEM micrographs of the worn surfaces of the unbo-
rided and borided cold-work tool steels are in Figures 5
and 6. Figure 5a shows a SEM micrograph of the wear
surface of the unborided steel. In the wear region of the
unborided steel, deeper and wider wear scars, debris,
surface grooves and cracks can be observed on its
surface. Figure 5b shows an EDS analysis obtained from
Figure 5a (point A). Fe-based oxide layers formed as a
result of the wear test.
Figure 6 shows SEM micrographs of the wear
surfaces of the borided cold-work tool steel. In Figure
6a, the worn surface of the borided steel is rougher, and
coarser wear debris is present. There are microcracks, a
delamination layer and wear, grooves, cracks and
abrasive particles on the worn surfaces of the boride
coatings (Figures 6a to 6d). In the wear region of the
borided cold-work tool steel there are cavities, probably
formed as a result of the layer fatigue (Figure 6) and
cracks formed due to the delamination wear. Figure 6e
shows an EDS analysis obtained from Figure 6c (point
B). The Fe-based oxide layers formed as a result of the
wear test. The spallation of the oxide layers in the sliding
direction and their orientation along the wear track were
identified. When the SEM image of the worn surface of
the unborided sample in Figure 5a is examined, it can be
seen that the wear marks are larger and deeper.
4 CONCLUSIONS
In this study, the wear behavior and some of the
mechanical properties of the borides on the surface of the
borided cold-work tool steel were investigated. Some of
the conclusions can be drawn as follows:
• Boride types formed on the surface of the cold-work
tool steel have a smooth morphology.
• The boride-layer thickness on the surface of the
cold-work tool steel was between 13.14–120.82 μm,
depending on the chemical compositions of the
substrates.
• The multiphase boride coatings that were thermoche-
mically grown on the cold-work tool steel were con-
stituted of the FeB, Fe2B, CrB, Cr2B and MoB
phases.
• The surface hardness for the borided steel was in the
range of 1806–2342 HV0.05, while for the untreated
steel substrate, it was 428 HV0.05.
• The lowest wear rate was obtained for the steel bo-
rided at 950 °C for 6 h, while the highest wear rate
was obtained for the unborided steel.
• The wear rate of the borided steel was found to be
approximately nine times lower than the wear rate of
the unborided steel.
5 REFERENCES
1 A. K. Sinha, B. P. Division, Boriding (Boronising) of Steels, ASM
Handbook, Vol. 4, J. Heat Treating, ASM International, 1991,
437–447
2 C. Bindal, A. H. Ucisik, Surface and Coatings Technology, 122
(1999), 208
3 A. G. von Matuschka, Boronizing, Heyden and Son Inc., Philadel-
phia, USA 1980, 11
4 N. Ucar, O. B. Aytar, A. Calik, Mater. Tehnol., 46 (2012) 6, 621–625
5 I. Ozbek, S. Sen, M. Ipek, C. Bindal, S. Zeytin, A. H. Ucisik, Va-
cuum, 73 (2004), 643
6 I. Gunes, Sadhana, 38 (2013), 527
7 S. Ulker, I. Gunes, S. Taktak, Indian Journal Eng. Mater. Sci., 18
(2011), 370
8 Y. Sun, T. Bell, G. Wood, Wear, 178 (1994), 131
9 K. Genel, Vacuum, 80 (2006), 451
10 I. Celikyurek, B. Baksan, O. Torun, R. Gurler, Intermetallics, 14
(2006), 136
11 I. Uslu, H. Comert, M. Ipek, O. Ozdemir, C. Bindal, Materials and
Design, 28 (2007), 55
12 J. H. Yoon, Y. K. Jee, S. Y. Lee, Surface and Coatings Technology,
112 (1999), 71
13 M. Hudáková, M. Kusý, V. Sedlická, P. Grga~, Mater. Tehnol., 41
(2007) 2, 81–84
14 I. Gunes, S. Ulker, S. Taktak, Materials and Design, 32 (2011), 2380
15 C. Bindal, A. H. Ucisik, Vacuum, 82 (2008), 90
16 O. Ozdemir, M. A. Omar, M. Usta, S. Zeytin, C. Bindal, A. H. Uci-
sik, Vacuum, 83 (2009), 175
17 T. Eyre, Wear, 34 (1975), 383
18 I. Gunes, Journal of Materials Science and Technology, 29 (2013),
662
19 S. Taktak, Surface and Coatings Technology, 201 (2006), 2230
20 M. Ulutan, M. M. Yildirim, O. N. Celik, S. Buytoz, Tribology Let-
ters, 38 (2010), 231
21 E. Atýk, U. Yunker, C. Meric, Tribology International, 36 (2003), 155
22 C. Li, B. Shen, G. Li, C. Yang, Surface and Coatings Technology,
202 (2008), 5882
23 I. Ozbek, C. Bindal, Vacuum, 86 (2011), 391
I. GUNES: TRIBOLOGICAL BEHAVIOR AND CHARACTERIZATION OF BORIDED COLD-WORK TOOL STEEL
Materiali in tehnologije / Materials and technology 48 (2014) 5, 765–769 769

M. VODÌROVÁ et al.: MICROSTRUCTURES OF THE Al-Fe-Cu-X ALLOYS PREPARED AT VARIOUS ...
MICROSTRUCTURES OF THE Al-Fe-Cu-X ALLOYS
PREPARED AT VARIOUS SOLIDIFICATION RATES
MIKROSTRUKTURA ZLITIN Al-Fe-Cu-X PO RAZLI^NIH
HITROSTIH STRJEVANJA
Milena Vodìrová, Pavel Novák, Filip Prù{a, Dalibor Vojtìch
Institute of Chemical Technology, Department of Metals and Corrosion Engineering, Technická 5, 166 28 Prague 6, Czech Republic
voderovm@vscht.cz
Prejem rokopisa – received: 2013-08-26; sprejem za objavo – accepted for publication: 2013-11-08
Aluminium alloys are usually prepared with conventional casting, but rapid-solidification methods lead to the alloys with better
mechanical properties or thermal stability. When an improved thermal stability is required, aluminium is alloyed with one or
more of the elements from the group of transition metals (TM), for example, Ni, Fe, Cr or Ti. These elements are characterized
by a low diffusivity and solubility in aluminium even at elevated temperatures, while Cu in an alloy forms a CuAl2 phase that
enables precipitation hardening. In this work, the microstructures of the Al-7Fe-4Cu, Al-4Fe-4Cu-3Ni and Al-7Fe-4Cu-3Cr
(mass fraction, w/%) alloys prepared with various solidification processes were investigated. The aim of this work was to
determine the changes in the microstructure caused by the increasing solidification rate and to determine the influence of the
copper present in each alloy. All the samples were prepared with single-roll melt spinning, water quenching of the melt and
conventional casting. The microstructures of the alloys were studied with light and scanning electron microscopy (SEM). The
phase composition was determined with X- ray diffraction (XRD). Vickers hardness (HV 5) and microhardness (HV 0.005)
were measured to compare the mechanical properties of the alloys. The microstructure and the hardness of the alloys strongly
depended on the solidification rate. The fine microstructure and high microhardness values obtained with melt spinning are
promising for the use of these alloys in special applications at elevated temperatures.
Keywords: aluminium alloy, rapid solidification, melt spinning, transition metals, microstructure
Zlitine aluminija se najpogosteje izdelujejo z navadnim ulivanjem, vendar pa metode hitrega strjevanja povzro~ijo nastanek
zlitin z bolj{imi mehanskimi lastnostmi ali toplotno stabilnostjo. Kadar se zahteva toplotna stabilnost, se aluminij legira z enim
ali dvema elementoma iz skupine prehodnih kovin (TM), na primer: Ni, Fe, Cr ali Ti. Zna~ilno za te elemente je majhna difuziv-
nost in topnost v aluminiju celo pri povi{anih temperaturah, medtem ko Cu v zlitini tvori fazo CuAl2, ki omogo~a izlo~evalno
utrjanje. V tem delu so preiskovane mikrostrukture zlitin Al-7Fe-4Cu, Al-4Fe-4Cu-3Ni in Al-7Fe-4Cu-3Cr (masni dele`i, w/%),
pripravljenih z razli~nimi postopki strjevanja. Namen tega dela je bil opredeliti razlike v mikrostrukturi, ki jih povzro~i
pove~anje hitrosti strjevanja, in opredeliti vpliv bakra v vsaki od navedenih zlitin. Vsi vzorci so bili pripravljeni z ulivanjem
tankega traku na bakren valj, z ohlajanjem v vodi in z navadnim ulivanjem. Mikrostruktura zlitin je bila pregledana s svetlobnim
mikroskopom in z vrsti~nim elektronskim mikroskopom (SEM). Sestava faz je bila dolo~ena z rentgensko difrakcijo (XRD).
Trdota HV 5 in mikrotrdota HV 0,005 sta bili izmerjeni za primerjavo z mehanskimi lastnostmi zlitin. Mikrostruktura in trdota
zlitin sta mo~no odvisni od hitrosti strjevanja. Drobnozrnata mikrostruktura in velika mikrotrdota, dobljeni z ulivanjem na valj
iz bakra, sta obetajo~i za uporabo teh zlitin v posebnih primerih pri povi{anih temperaturah.
Klju~ne besede: zlitina aluminija, hitro strjevanje, ulivanje na bakreni valj, prehodne kovine, mikrostruktura
1 INTRODUCTION
Aluminium alloys processed with the conventional
technologies, such as casting and forming, are widely
used in many technical branches such as the aerospace
and automotive industries. The main advantages of alu-
minium alloys are price, good strength-to-weight ratio,
good castability, formability or the ability of precipita-
tion hardening. However, the mechanical properties of
traditional alloys made of Zn, Mg or Cu strongly degrade
at elevated temperatures, which means that their appli-
cation is then limited to 150–200 °C. One way of
improving the thermal stability of aluminium alloys is to
use the elements from the transition metals group (TM).
Transition metals, such as Ni, Fe, Cr or Mo, are charac-
terized by a low diffusivity and solubility in aluminium
even at elevated temperatures and they are able to stabi-
lize the materials properties up to relatively high tempe-
ratures (about 400 °C). Cu is used as an alloying element
to increase both the strength and the hardness due to the
CuAl2 phase that allows precipitation hardening of the
material.1 Conventional casting processes produce the
alloys containing coarse particles of hard and brittle
Al-TM intermetallic phases, degrading the mechanical
properties.2 Therefore, it is desirable to keep these alloy-
ing elements dissolved in the matrix or in the finely
dispersed intermetallic particles. A fine microstructure
can be obtained by increasing the solidification rate, e.g.,
by atomisation or melt spinning.3,4
The alloying elements mentioned above are often the
contaminants of Al scrap. In recent years, the consump-
tion of aluminium alloys in engineering has been rising,
causing the problems of recycling and waste disposal. Al
scrap is never only pure aluminium, but it is mixed with
steel, cast iron, copper alloys, etc. Parts of ferromagnetic
iron-based alloys can be separated using magnetic sepa-
ration. The other way is to dilute the melt with pure alu-
minium, but this technique increases the cost of recycled
aluminium alloys. In general, the transition elements
Materiali in tehnologije / Materials and technology 48 (2014) 5, 771–775 771
UDK 669.715:621.74.046 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(5)771(2014)
included, e.g., in the austenitic-stainless-steel admixtures
in Al scrap are very difficult and costly to remove. We
would like to develop a new way of preparing the alloys
with interesting mechanical properties and a thermal sta-
bility using this contaminated Al scrap. The manufac-
tured alloys would have better properties such as hard-
ness, thermal stability and ductility associated with the
low density. Due to the mentioned properties these alloys
could be able to replace titanium alloys in some appli-
cations, while their price and density would be lower.
This work describes the microstructure and properties
of the Al-Fe-Cu-X alloys prepared at various cooling
rates. These alloys simulate the real alloys originating
from melting the contaminated Al scrap. The alloys with
the mentioned chemical composition have not been stu-
died yet; there are only a few studies dealing with the
microstructures of the rapidly solidified ternary alloys or
systems of the quasicrystal chemical compositions.1,5–8
2 MELT-SPINNING PRINCIPLES
Atomization of a melt with an inert gas or water
produces the powders that solidify with the rate ranging
from 102–104 K s–1. Melt spinning allows even higher
cooling rates (104–106 K s–1). In this process, a molten
alloy is ejected on a high-speed rotating metallic wheel.
The alloy solidifies rapidly in contact with the wheel.
This method produces thin ribbons, whose thickness
varies in the order of ten micrometres. Due to rapid-soli-
dification processes transition metals can be added to
aluminium even above their equilibrium-solubility limits.
Increased solidification rates lead to the formation of
supersaturated solid solutions and fine particles of
metastable and stable intermetallic phases. The amounts
of intermetallic phases are reduced and the shape is
usually spherical. The slow decomposition of a super-
saturated solution containing transition metals at higher
temperatures can lead to the precipitation strengthening
of the material. The layout of the melt-spinning process
is shown in Figure 1.3
To obtain a bulk material, the rapid-solidification
process is associated with the compaction of the RS
product. At first, the powder has to be milled, e.g., by
cryogenic milling, to obtain a metallic powder with a
well-preserved fine microstructure The most suitable
technology is hot extrusion or e.g., hot isostatic pressing
(HIP) or spark plasma sintering (SPS).9–11
3 EXPERIMENTAL WORK
Alloys with the chemical compositions of
Al-7Fe-4Cu, Al-4Fe-4Cu-3Cr and Al-4Fe-4Cu-3Ni were
prepared by melting the master alloy Al-11Fe (w/%)
with the additions of pure Cu, Cr and Ni in an elec-
tric-resistance furnace in a graphite crucible and then
poured into a brass mould. The second series was pre-
pared by remelting the alloy and subsequent water
quenching. The third series was prepared with single-roll
melt spinning. The melting was carried out under an
argon protective atmosphere and the temperature of the
melt was 950 °C. The material was melted in a quartz-
glass nozzle and then poured onto a copper wheel using
overpressured argon. The circumferential speed of the
wheel was 30 m s–1. The process yielded aluminium-
alloy ribbons approximately 40 μm thick. The metallo-
graphic cuts of the investigated alloys were etched in
Kroll’s reagent (10 mL HF, 5 mL HNO3 and 85 mL H2O)
and investigated with an Olympus PME3 light micro-
scope and a TESCAN VEGA 3 LMU scanning electron
microscope (SEM) equipped with an Oxford Instruments
INCA 350 EDS analyser. The phase composition was
determined with X- ray diffraction (XRD, PANalytical
X’Pert Pro). The mechanical properties of the investi-
gated alloys were examined with Vickers-hardness
measurements with the 5 kg (HV 5) and 0.005 kg (HV
0.005) loads. The microhardness was measured using a
Neophot 2 light microscope equipped with a Hanemann
microhardness tester.
4 RESULTS AND DISCUSSION
4.1 Microstructure
The microstructures of the aluminium alloys prepared
by conventional casting into a brass mould are shown in
Figures 2 to 4. It is obvious that the microstructure
obtained after the conventional casting is composed of an
inhomogeneous material with large amounts of coarse
and brittle binary intermetallic phases Al13Fe4 and CuAl2
and ternary phase Al23CuFe4 in the solid solution of the
alloying elements in aluminium. Moreover, the
nickel-alloyed material contains Al4Ni3, Al7Cu4Ni and
Al75Ni10Fe15 as well.
Figures 5 to 7 show the microstructures of the alloys
prepared by melting at 1000 °C and then water quenched.
Intermetallic phases Al13Fe4 and CuAl2 become finer due
M. VODÌROVÁ et al.: MICROSTRUCTURES OF THE Al-Fe-Cu-X ALLOYS PREPARED AT VARIOUS ...
772 Materiali in tehnologije / Materials and technology 48 (2014) 5, 771–775
Figure 1: Melt-spinning principles
Slika 1: Shematski prikaz ulivanja na bakreni valj
M. VODÌROVÁ et al.: MICROSTRUCTURES OF THE Al-Fe-Cu-X ALLOYS PREPARED AT VARIOUS ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 771–775 773
Figure 7: Microstructure of the water-quenched Al-4Fe-4Cu-3Ni
(SEM)
Slika 7: Mikrostruktura Al-4Fe-4Cu-3Ni po ohlajanju v vodi (SEM)
Figure 4: Microstructure of the as-cast Al-4Fe-4Cu-3Ni (SEM)
Slika 4: Strjevalna struktura Al-4Fe-4Cu-3Ni (SEM)
Figure 5: Microstructure of the water-quenched Al-7Fe-4Cu (SEM)
Slika 5: Mikrostruktura Al-7Fe-4Cu po ohlajanju v vodi (SEM)
Figure 2: Microstructure of the as-cast Al-7Fe-4Cu (SEM)
Slika 2: Strjevalna struktura Al-7Fe-4Cu (SEM)
Figure 3: Microstructure of the as-cast Al-4Fe-4Cu-3Cr (SEM)
Slika 3: Strjevalna struktura Al-4Fe-4Cu-3Cr (SEM)
Figure 6: Microstructure of the water-quenched Al-4Fe-4Cu-3Cr
(SEM)
Slika 6: Mikrostruktura Al-4Fe-4Cu-3Cr po ohlajanju v vodi (SEM)
to a more intensive cooling and the amount of the
intermetallics in the microstructure is decreasing. In
Al-7Fe-4Cu, quasicrystalline Al65Cu20Fe15 is formed
instead of a stable Al23CuFe4. On the other hand, no
quasicrystalline phases were detected in the
water-quenched Al-4Fe-4Cu-3Cr, but Al13Cr2 occurred in
the microstructure. No differences in the phase compo-
sition of the alloy containing nickel in the as-cast and
water-quenched states were detected.
The microstructures of the rapidly solidified alloys in
the longitudinal cuts are documented in Figures 8 to 10.
It is evident that the microstructure of a prepared ribbon
is strongly dependent on the distance from the cooling
wheel. On the wheel side, which is cooled more inten-
sely, a supersaturated solid solution with nanocrystalline
intermetallics is formed. On the free side, fine spherical
intermetallic particles are formed. The saturation of the
solution decreases when moving the ribbon from the
wheel side to the free side. The amounts of the Al13Fe4
and CuAl2 phases are negligible; instead of them,
metastable phases Al4Ni3, Al75Ni10Fe15, Al23CuFe4 or
Al7Cu4Ni and quasicrystalline phase Al65Cu20Fe15, are
formed.12–14 The phase compositions of all the samples
M. VODÌROVÁ et al.: MICROSTRUCTURES OF THE Al-Fe-Cu-X ALLOYS PREPARED AT VARIOUS ...
774 Materiali in tehnologije / Materials and technology 48 (2014) 5, 771–775
Figure 12: XRD patterns of the Al-4Fe-4Cu-3Cr prepared with diffe-
rent methods
Slika 12: XRD-posnetki Al-4Fe-4Cu-3Cr, izdelane po razli~nih me-
todah
Figure 10: Microstructure of the rapidly solidified Al-4Fe-4Cu-3Ni
(SEM)
Slika 10: Mikrostruktura hitro strjenega traku iz Al-4Fe-4Cu-3Ni
(SEM)
Figure 11: XRD patterns of the Al-7Fe-4Cu prepared with different
methods
Slika 11: XRD-posnetki Al-7Fe-4Cu, izdelane po razli~nih metodah
Figure 9: Microstructure of the rapidly solidified Al-4Fe-4Cu-3Cr
(SEM)
Slika 9: Mikrostruktura hitro strjenega traku iz Al-4Fe-4Cu-3Cr
(SEM)
Figure 8: Microstructure of the rapidly solidified Al-7Fe-4Cu (SEM)
Slika 8: Mikrostruktura hitro strjenega traku iz Al-7Fe-4Cu (SEM)
are summarized in Figures 11 to 13. The phase composi-
tion, the amounts of intermetallics and the particle size
are inevitably dependent on the solidification rate.
4.2 Hardness measurement
The Vickers hardness of the as-cast and water-
quenched samples was measured with a 5 kg load. The
microhardness of the rapidly solidified alloys was
measured with a 5 g load because of the low thickness of
the produced ribbons. The microhardness of the rapidly
solidified alloys was measured in the centre of a ribbon
to avoid the influence of the epoxy resin surrounding the
sample. The measurement results are shown in Figure
14. It is obvious that the hardness increases with the
increasing solidification rate. The as-cast alloys consist
of large sharp-edged particles of the intermetallics that
have a negative effect on the hardness. A decrease in the
particle size and the strengthening caused by the pre-
sence of the supersaturated solutions are the main expla-
nations of the increased hardness.
5 CONCLUSIONS
This work focused on a comparison of the micro-
structures of the Al-Cu-Fe-X alloys prepared at various
solidification rates. The microstructures of the alloys
prepared with traditional casting and water quenching
are considerably inhomogeneous. There are large
amounts of coarse Al13Fe4 and CuAl2 intermetallic
phases in the aluminium matrix. The amount of interme-
tallics decreases and the particles become finer, if the
solidification rate increases. In the rapidly solidified
alloys, the amounts of Al13Fe4 and CuAl2 are limited
because these phases are replaced by metastable and
quasicrystalline intermetallics. The microstructures of
RS alloys consist of aluminium supersaturated with
transition metals and spherical intermetallics. The hard-
ness of the investigated materials is hardly dependent on
the cooling rate; higher values were reached for very fine
materials.
Acknowledgement
This research was financially supported by the Czech
Science Foundation, within project No. P108/12/G043.
6 REFERENCES
1 S. J. Andersen, Materials Science and Engineering A, 179–180
(1994), 665–668
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(1992), 45–65
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Processing Technology, 87 (1999), 1–27
11 L. Wang, J. Zhang, W. Jiang, Int. Journal of Refractory Metals and
Hard Materials, 39 (2013), 103–112
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Urban, Materials Science and Engineering A, 226–228 (1997),
976–980
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(2004), 150–174
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tallics, 16 (2008), 847–853
M. VODÌROVÁ et al.: MICROSTRUCTURES OF THE Al-Fe-Cu-X ALLOYS PREPARED AT VARIOUS ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 771–775 775
Figure 14: Hardness and microhardness measurements
Slika 14: Izmerjene trdote in mikrotrdote
Figure 13: XRD patterns of the Al-4Fe-4Cu-3Ni prepared with diffe-
rent methods
Slika 13: XRD-posnetki Al-4Fe-4Cu-3Ni, izdelane po razli~nih me-
todah

D. KYTÝØ et al.: ASSESSMENT OF THE POST-IMPACT DAMAGE PROPAGATION IN ...
ASSESSMENT OF THE POST-IMPACT DAMAGE
PROPAGATION IN A CARBON-FIBRE COMPOSITE
UNDER CYCLIC LOADING
OCENA NAPREDOVANJA PO[KODBE PO UDARCU PRI
PONAVLJAJO^IH SE OBREMENITVAH KOMPOZITA Z
OGLJIKOVIMI VLAKNI
Daniel Kytýø1, Tomá{ Fíla2, Jan [leichrt2, Tomá{ Doktor1, Martin [perl1
1Institute of Theoretical and Applied Mechanics, v.v.i., Academy of Sciences of the Czech Republic, Prosecká 76, 190 00 Prague 9, Czech
Republic
2Czech Technical University in Prague, Faculty of Transportation Sciences, Department of Mechanics and Materials, Konviktská 20, 110 00
Prague 1, Czech Republic
kytyr@itam.cas.cz
Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2013-11-11
Carbon fibre in polyphenylene sulfide composites (C/PPS) became a popular material in the aircraft industry but its fragility and
low impact resistance limits its application in primary aircraft structures. This study is focused on damage propagation in the
laminated composites reinforced with carbon fibres. The damage may be inflicted during the ground maintenance, by an inflight
bird strike or during a flight in severe meteorological conditions (heavy storms). The initial damage was created by a
drop-weight out-of-plane impact using a spherical indenter. The response of the material was analysed by monitoring the
impacted zones and their propagation history. The influenced area and specimen thickness in the centres of indents were chosen
as the degradation parameters. The post-impact damage propagation induced by cyclic loading was assessed using a
custom-designed computer-controlled laser-profilometery device. Both the upper and lower profiles of the specimen were
scanned during the interruptions of the fatigue test. Global deformation was described with an analytically determined
centroidal-axis curve. Local topography changes were obtained with a subtraction of this curve. Surface-deformation maps were
created and used for a demonstration of the damage propagation in the specimen.
Keywords: carbon-fibre composites, post-impact damage, laser profilometry
Ogljikova vlakna v kompozitih iz polifenilen sulfida (C/PPS) so postala priljubljen material v letalski industriji, toda njihova
krhkost in slaba odpornost proti udarcem omejujeta njihovo uporabo v primarnih letalskih konstrukcijah. Ta raziskava se
osredinja na napredovanje po{kodbe na laminiranih kompozitih, oja~anih z ogljikovimi vlakni. Po{kodba lahko nastane med
vzdr`evanjem na tleh, pri tr~enju s ptico med letom ali med letom v hudih vremenskih razmerah (huda nevihta). Za~etna
po{kodba je bila narejena z udarno pregibnim preizkusom s kroglastim vtiskovalcem. Odziv materiala je bil analiziran z
opazovanjem obmo~ij udarca in potekom napredovanja. Prizadeto obmo~je in debelina vzorca v podro~ju vtiska sta bila izbrana
kot parametra degradacije. Napredovanje po{kodbe po cikli~nem obremenjevanju po udarcu je bilo ocenjeno s po meri
oblikovane ra~unalni{ko vodene naprave za lasersko profilometrijo. Zgornji in spodnji profil vzorca sta bila skenirana med
prekinitvami preizku{anja utrujenosti. Celotna deformacija je bila opisana z analiti~no dolo~eno krivuljo te`i{~nice. Lokalne
spremembe topografije so bile dobljene z od{tetjem te krivulje. Ustvarjeni videzi deformacije povr{ine so bili uporabljeni za
prikaz napredovanja po{kodbe na vzorcu.
Klju~ne besede: kompoziti z ogljikovimi vlakni, po{kodba po udarcu, laserska profilometrija
1 INTRODUCTION
The design and safe operation of lightweight struc-
tures, especially in the aviation industry, is particularly
important and challenging due to the inauspicious load
spectra composed of a large number of low-amplitude
cycles and sudden impacts1. Low-amplitude cycles are
caused by aerodynamic loads and engine vibrations.
Wayward strikes may be inflicted during the ground
maintenance, by inflight collisions (bird strikes) or
severe meteorological conditions (heavy storms).
The damage-tolerance approach commonly used in
aerospace engineering requires a comprehensive know-
ledge of the material-degradation process and a reliable
prediction of a structure safe life2. The thermoplastic
composites commonly used for these purposes allow an
application of an optimised manufacturing technology3,4.
An application of a polymeric matrix lowered the
tendency towards brittle behaviour (common for car-
bon-fibre composites) and exhibited the advantages of
high chemical resistivity, insensitivity to moisture, good
fatigue performance5,6 and recyclability.
Micromechanical modelling of the composites with
imperfections7 sufficiently describes the degradation
process. However, the material models based on the
X-ray computed tomography of the specimen represent-
ing the material at the macroscopic level including a
complex microstructure could not be evaluated using the
finite-element simulations with the plasticity applied due
to the computational complexity and enormous memory
requirements8. The presented work aimed to extend the
range of non-destructive testing (NDT) techniques com-
prising the lock-in thermography9 or the modified-im-
pulse excitation technique10.
Materiali in tehnologije / Materials and technology 48 (2014) 5, 777–780 777
UDK 677.4:66.017 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(5)777(2014)
2 MATERIALS AND METHODS
2.1 Specimen description
The base material, a carbon-fibre/polyphenylene
sulphide (C/PPS) composite manufactured by Letov le-
tecká výroba, s. r. o., was delivered as plates with a
thickness of (2.5 ± 0.05) mm. The material consists of
quasi-isotropic 8-ply carbon fabric with its volume
fraction higher than 90 %, bonded with a thermoplastic
matrix. The surface is covered with a thin glass-fibre
cloth protecting the core against mechanical and che-
mical influences. The final specimens with a rectangular
shape with the dimensions of 250 mm × 25 mm were cut
from the plates using a water-jet cutter.
2.2 Initial damage
The first step of the experimental procedure was to
inflict the initial damage to the specimens under con-
trolled conditions. A drop tower designed within project
SGS12/163/OHK2/2T/16 with the maximum impact
energy of 50 J was used. The strike was carried out using
a spherical indenter with a diameter of 20 mm and the
energies of (10, 20 and 15) J on (30, 50 and 70) % of the
length of the samples. The imprints of the diameter in
the range of millimetres and the depth in the range of
tens of micrometers then occurred.
2.3 Fatigue loading
For a life-cycle assessment the specimens were cycli-
cally loaded using a Mikrotron (Russenberger Prüfma-
schinen, AG) resonant testing machine (Figure 1). To
ensure the loading at the chosen stress level (33 % of the
tensile strength) the mean loading-force value of 6 kN
and the amplitude of 5 kN were set. A sinusoidal force
was applied in the force-driven experiments.
Due to a relatively high testing frequency (approxi-
mately 75 Hz), the experiment was monitored with a
thermal imaging camera SC7600 (FLIR Systems, Inc.).
To prevent exceeding 50 % of the glass transition tempe-
rature the specimen temperature was held at maximally
60 °C. At the same time, the lower frequency limit was
set in order to avoid a specimen rupture11. The fatigue
experiment was interrupted six times at the predefined
numbers of cycles to perform profile scanning.
2.4 Profile measurement
To obtain the information about damage propagation
during the life cycle, a set of profilometery experiments
was performed. A custom-designed scanning device
equipped with laser scanner ScanControl LLT2600-25
(Micro-Epsilon Messtechnik) depicted in Figure 2 was
used for this purpose. The device allowed us to measure
the line profiles with the length of 20–40 mm, defined by
1024 measured points. The altitude resolution of the
scans was 4 μm. The scanner was mounted on a motor-
ised computer-controlled single-axis linear stage with the
minimum incremental motion of 10 μm and the on-axis
accuracy of ± 0.5 μm. One scanning sequence took
approximately 15 minutes.
2.5 Damage-propagation assessment
The changes in the impact depth, the sample thick-
ness and the area of influenced zones were chosen as the
degradation parameters. The automatic procedure for a
surface reconstruction (Figure 3) and profile-change
assessment was carried out using the tools developed in
the MATLAB (Mathworks, Inc.) computational environ-
ment. The variable position of the samples in the scann-
ing area required the use of the corner detection
D. KYTÝØ et al.: ASSESSMENT OF THE POST-IMPACT DAMAGE PROPAGATION IN ...
778 Materiali in tehnologije / Materials and technology 48 (2014) 5, 777–780
Figure 2: Custom-designed computer-controlled profilometery device
equipped with a ScanControl LLT2600-25 laser scanner
Slika 2: Po meri oblikovana ra~unalni{ko vodena naprava za profilo-
metrijo, opremljena z laserskim opti~nim bralnikom ScanControl
LLT2600-25
Figure 1: Experimental device for dynamic loading
Slika 1: Eksperimentalna naprava za dinami~no obremenjevanje
algorithm based on the altitude threshold to detect the
specimen boundaries in the captured data. Transforma-
tion functions were obtained and the objects were
transformed into a unitary coordination system.
Divergence of the laser beam was taken into account
for the real-altitude matrix estimation and the blur of the
edges caused by the same effect was reduced with
gradient filters. The curvature of the surfaces was not
caused only by the local impact zones but also by the
overall bending of the samples due to a combination of
the initial impact damage and cyclic loading. A piece-
wise continuous second-order curve (the centroidal axis)
was fitted and set as a new reference level. Then the
altitude matrices were updated. On the straightened
surfaces, the local impacted zones were quantified (area,
maximum depth) using the data-registration procedure.
From the subtraction of the upper and lower profile, the
change in the sample thickness was obtained.
3 RESULTS
Based on the reconstructed profiles from the laser
measurements, the influenced zones were identified on
the basis of thresholding. In the areas of interest, the
impact depression depth and the local thickness were
assessed. Propagation of the chosen degradation para-
meters on two selected samples for several distinct
impact levels is depicted in Figure 4.
Damage propagation exhibits similar evolution on
different tested samples. The most significant parameter
was the maximum depth of the impact on the impacted
side. The initial depth corresponds to the strike energy,
while later the depth decreases with the increasing
number of the loading cycles. The area of influenced
zones grows with the number of the loading cycles but,
surprisingly, the initial areas were not proportional to the
strike energy. The area of damaged zones inflicted by
lower energy impacts also showed a faster increase. The
changes in the thickness of the samples due to the
D. KYTÝØ et al.: ASSESSMENT OF THE POST-IMPACT DAMAGE PROPAGATION IN ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 777–780 779
Figure 3: Reconstruction of the sample surface based on laser triangulation
Slika 3: Rekonstrukcija povr{ine vzorca, ki temelji na laserski triangulaciji
Figure 4: a) Increase in the influenced zones, b) the maximum depth
of the impact depression and c) the thickness of the sample plotted
against the number of loading cycles
Slika 4: a) pove~anje obsega prizadetega obmo~ja, b) maksimalne
globine udrtine ter c) debelina vzorca glede na {tevilo ciklov obreme-
nitve
influenced zones were negligible as the differences in the
thickness were only two or three times higher than the
noise.
4 CONCLUSIONS
The presented study describes the possibility of a
time-lapse profilometery measurement for an evaluation
of the post-impact damage propagation in a C/PPS
composite under cyclic loading. The chosen parameters
(the area of impacted zone, the maximum depth and the
sample thickness) provide the information about damage
accumulation in the material. Generally, laser profilo-
metery is a suitable method for the NDT testing and
evaluation of the surface damage. The described modi-
fied method is applicable to bigger components and
structures. With respect to our measured data, the
reliability of the method was reduced by the resolution
of the available laser scanner.
Acknowledgements
The research was supported by Technology Agency
of the Czech Republic (grant No. TA03010209), Grant
Agency of the Czech Technical University in Prague
(grant No. SGS12/205/OHK2/3T/16), research plan of
the Ministry of Education, Youth and Sports
MSM6840770043 and by institutional support
RVO: 68378297.
5 REFERENCES
1 R. Aoki, J. Heyduck, An Experimental Study of Impact-Damaged
Panels under Compression Fatigue Loading, In: J. Füller et al. (eds.),
Developments in the Science and Technology of Composite Mate-
rials, Elsevier Science Publishers Ltd., 1990, 633–642 1990,
633–642
2 J. P. Gallagher, USAF damage tolerant design handbook, Flight
Dynamics Laboratory, Air Force Wright Aeronautical Laboratories,
1984
3 Z. Padovec, M. Rù`i~ka, Mechanics of Composite Materials, 49
(2013) 2, 221–230
4 I. C. Finegan, R. F. Gibson, Composite Structures, 44 (1999) 2–3,
89–98
5 J. Minster, O. Bláhová, J. Hristova, J. Luke{, J. Nìme~ek, M. [perl,
Journal of Applied Polymer Science, 123 (2012) 4, 2090–2094
6 J. Minster, O. Blahová, J. Luke{, J. Nìme~ek, Mechanics of Time-
Dependent Materials, 14 (2010) 3, 243–251
7 M. [ejnoha, J. Zeman, International Journal of Engineering Science,
46 (2008) 6, 513–526
8 J. Vorel, J. Zeman, M [ejnoha, International Journal for Multiscale
Computational, 11 (2013) 5, 443–462
9 R. Montanini, Infrared Physics & Technology, 53 (2010) 5, 363–371
10 D. Kytyr, T. Fila, J. Valach, M. [perl, UPB Scientific Bulletin, Series
D: Mechanical, 75 (2013) 2, 157–164
11 M. Pirner, S. Urushadze, International Applied Mechanics, 40 (2004)
5, 487–505
D. KYTÝØ et al.: ASSESSMENT OF THE POST-IMPACT DAMAGE PROPAGATION IN ...
780 Materiali in tehnologije / Materials and technology 48 (2014) 5, 777–780
M. KORBÁ[ et al.: POSSIBILITIES FOR INCREASING THE PURITY OF STEEL IN PRODUCTION ...
POSSIBILITIES FOR INCREASING THE PURITY OF
STEEL IN PRODUCTION USING
SECONDARY-METALLURGY EQUIPMENT
MO@NOSTI POVE^ANJA ^ISTOSTI JEKLA PRI PROIZVODNJI Z
UPORABO OPREME ZA SEKUNDARNO METALURGIJO
Martin Korbá{1, Libor ^amek2, Milan Raclavský3
1Vítkovice Heavy Machinery a.s., Ruská 2887/101, Vítkovice, Ostrava, Czech Republic
2V[B-Technical University of Ostrava, Faculty of Metallurgy and Material Engineering, Department of Metallurgy and Foundry, 17. listopadu
15/2172, Ostrava – Poruba, Czech Republic
3ECOFER s.r.o., Ka{tanová 182, 739 61 Tøinec, Dolní Lí{tná, Czech Republic
martin.kor@centrum.cz
Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2013-12-09
The possibilities for increasing the purity of steel during the production of the liquid phase using secondary metallurgy mainly
relate to affecting the number of emerging occlusions, their size, morphology and chemical composition. The metallographic
purity of steel during production in an electric-arc furnace (EAF), in a ladle furnace (LF) and during processing with VD
caisson technology was assessed. The steel samples were processed by means of an electron microscope and were
simultaneously tested using the single-spark-evaluation (SSE) method.
The aim of this investigation was to find the possibility for an operative steel-quality control and to obtain the desired
mechanical properties already while processing the liquid phase.
Keywords: steel purity, secondary metallurgy, occlusions, electron microscope, single-spark evaluation
Mo`nosti pove~anja ~istosti staljenega jekla med proizvodnjo z uporabo sekundarne metalurgije so predvsem glede na koli~ino
vklju~kov, njihove velikosti, vrsto morfologije in kemijsko sestavo. Metalografsko je bila ugotavljana ~istost jekla med
proizvodnjo v elektrooblo~ni pe~i (EAF), nato v ponov~ni pe~i (LF) in po obdelavi v VD. Vzorci so bili pregledani z
elektronskim mikroskopom in so bili preizku{eni z metodo plamenskega od`iganja (SSE).
Namen je bil poiskati u~inkovito kontrolo jekla in vplive na `elene mehanske lastnosti `e med obdelavo taline.
Klju~ne besede: ~istost jekla, sekundarna metalurgija, zapore, elektronska mikroskopija, ocena s plamenskim od`iganjem
1 INTRODUCTION
The continuously increasing requirements for
achieved levels of mechanical values like ductility,
toughness and fatigue properties for steel are increasing
the demands put on manufacturers when it comes to
searching for new technological processes. One of the
main possibilities for achieving higher levels of mecha-
nical properties for steel increases with its metallurgical
purity. This can be achieved with the technological
processes utilised in secondary metallurgy that can affect
the size, the quantity of inclusions and their morpho-
logy.1 Within the conditions existing at the steelworks of
Vítkovice Heavy Machinery a.s. (VHM), there are
facilities for processing steel in a ladle furnace (LF) and
for vacuum steel processing in a caisson (VD).
The objective of the study was to assess the conti-
nuous metallographic purity in the course of steel pro-
duction in an electric-arc furnace (EAF), in a ladle
furnace (LF) and during vacuum processing in a caisson
using the VD technology. The study was based on the
idea that the implemented technological processes (TS
No. 1 and TS No. 2) can assist in obtaining metallo-
graphic purity.
Metallic disk-pin samples were collected during the
steel’s production from secondary-metallurgical aggre-
gates of LF and VD. Consequently, they were processed
with an optical emission spectrometer, which assesses
the analytical signal, while differentiating each indivi-
dual spark with the single-spark-evaluation (SSE)
method. A part of the analytical signal corresponds with
the average content of an element in the collected sample
(classically in the course of the steel’s production
process), while the second signal part corresponds with
the elemental concentration in the inclusion. The single-
spark-evaluation (SSE) method helps us assess infor-
mation about the amount, the type and the composition
of the inclusion. A big advantage of the SSE method is
the fact that the sample preparation for the analysis with
a spectrometer is identical to the disk-pin sample
preparation. Analyses of the standard control sample of
the steel’s chemical composition were extended by only
20 s. The method took place online. The software for the
optical emission spectrometer was set up for an elemen-
tal analysis in relation to the chemical composition of the
produced steel.2
In order to express certain opinions about the values
measured with the SSE method, identical samples,
collected in the course of the steel’s production, were
consequently also processed with a roentgen spectral
analysis (using an electron microscope) and subjected to
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a metallographic assessment, which was generally
recognised within the metallurgical practice. The ana-
lyses of the inclusions with the assistance of the roentgen
spectral analysis were performed on the basis of the
Pirelli Norm No. 18. V. 008 rev. 7 metallographic test of
the microstructure and of the defects of the rod wires.
The test was slightly modified for the disk-pin measure-
ments. The device scans the image in the BSE mode
(reflected electrons); the image size is 1024 × 1024
pixels at a magnification of 500-times; the field of vision
has a size of 0.25 mm × 0.25 mm; and the scanning time
for a single pixel is 2800 ns. It looks for inclusions based
on the set brightness threshold and then it analyses them.
The threshold is set in such a way that all the oxides are
recorded, while in the case of sulphides only the partly
complex ones (with an oxidative core) are recorded. The
analysis is performed for 5 s at about 25 % of the
detector’s dead time. This corresponds to an intensity of
about 4500 impulses per second. The electron beam
focuses on the particle’s centre of gravity during the
analysis. The set of results from the analysed particles
passes through a filter that ensures the exclusion of the
non-oxidative included particles and, as a result, records
them in a ternary diagram with the silicates in one cor-
ner, aluminates in the second corner and the remaining
oxides (Mn, Ca, and Mg) in the third corner. Particles
larger than 1 μm are measured. The total assessed area of
a sample consisted of 80 fields, i.e., 5 mm2, which
corresponded to the size of the metallic-sample areas
analysed with the SSE method. The measured areas were
fully comparable for both measurement techniques.
2 EXPERIMENTAL
The samples were manually collected from the fur-
nace with the assistance of submersible disk-pin sam-
plers and they were identical for both methods of steel
purity assessment. The first sample, marked as LF1,
represented the beginning of the steel-processing tech-
nology in the ladle furnace and, at the same time, the
result of the preliminary steel de-oxygenation in the ladle
furnace (LP) after the completed tapping, without slag,
of the electric-arc furnace. In order to ensure standardi-
sation, a temperature of 1580 °C was selected for the LF
at the beginning of the collection of each LF1 sample.
Considering a certain time delay, related to the necessary
heating of the tapped and de-oxygenated melt to that
temperature, we assumed a good melt-concentration
homogeneity was achieved, including a good liquid, and
mixed the newly occurring basic slag. The sample LF2
represents the melt situation at the end of the melt
processing in the ladle furnace, where the main stress
was put on the creation of the melt slag mode. The LP
with the melt was prepared for the subsequent steel
vacuuming in the caisson furnace using the VD process.
A detailed technological prescription TS 1, prepared by
VHM, established the steel temperature before the
vacuuming, the steel’s chemical composition, the slag
parameters, and the oxygen activity. The samples VD1
and VD2 were collected at the beginning and at the end
of the steel’s vacuum processing in the furnace. The 15
min of homogenisation using porous argon followed
after the sample’s VD2 collection. The samples were
collected from melts processed by following two techno-
logical processes. The sample collections from melts No.
K58494, K58500, K58507, K58508, K58509, K58555
were performed following the technological process TS
No. 1, while the sample collections from melts No.
K58489 and No. K58495 followed the technological
process TS No. 2 prepared at VHM.
The technological instructions differ, especially in
terms of the method of the preliminary steel de-oxygena-
tion in the ladle furnace after the tapping of the electric
arc furnace. Other principles of the melt management in
the ladle furnace or in the caisson are almost identical.2
TS No. 1
• 2 kg/t of CaC2 into the ladle furnace (LP) before the
tapping,
• FeSi to 0.15 % (steel in the furnace without Si after
the oxygenation),
• Calcinated anthracite in the course of carbonaceous
steels, e.g., C45,
• Al 0.3–0.6 kg/t into the flow in the course of the first
tapping third, depending on the produced steel and
according to the melted C content
• Pieces of CaO + synthetic slag + FeMn follow (up to
2/3 of the production composition),
• Possibly FeCr (none in the case of melts monitored
by us)
• 0.3–0.6 kg/t of Al, again in the course of the last
tapping third
TS No. 2
• 3.0 kg/t CaC2 + min. of 3.0 kg/t of the calcinated
anthracite into the ladle furnace (LP)
• FeMn (up to 2/3 of the production composition) +
CaO + synthetic slag during the tapping, before the
end of the tapping
• 0.5–0.7 kg/t of Al, according to the melted C
3 DISCUSSION
When assessing the research results, we must
consider that the performed analysis could not be
identical to the resulting composition of inclusions in the
final product. The difference is given especially by the
different speeds of cooling for the relatively small
metallic sample, when compared with, for example, an
ingot of many tons. The analysis does not present the
inclusion composition and their quantity in the final
product, but the actual inclusion types and the purity of
the steel, which are analysed in the course of the
technological steps LF and VD.
The following set of figures (Figures 1 to 6) sum-
marises the results of the analysis of the number of
inclusions performed with a microprobe.
M. KORBÁ[ et al.: POSSIBILITIES FOR INCREASING THE PURITY OF STEEL IN PRODUCTION ...
782 Materiali in tehnologije / Materials and technology 48 (2014) 5, 781–786
The above-presented figures clearly show that the
inclusion density in the individual melts progressively
decreased in the course of the technological process. The
fundamental decrease in the inclusion density took place
in the course of the melt processing in the ladle furnace.
This trend is clear for inclusions up to 4 μm in size. The
density of the smallest inclusions then changed only a
little, or not at all, in the subsequent vacuum steel
processing because of their low concentration. In the
case of some melts, even a small increase was noticed.
Inclusions larger than 4 μm occur rarely in the samples
and their occurrence is thus accidental. No clear trend
was noticed. However, we can assume that their
occurrence at the end of the technological process in VD
is close to zero. The melts 58489, 58495 and 58527,
which went through a different de-oxygenation process,
differ to some extent from the set. The density of
inclusions is lower at the beginning. These melts were
produced by TS No. 2.
M. KORBÁ[ et al.: POSSIBILITIES FOR INCREASING THE PURITY OF STEEL IN PRODUCTION ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 781–786 783
Figure 5: Numbers of inclusions in individual areas of the ternary
diagram for the sample VD1 collection
Slika 5: [tevilo vklju~kov v posameznih podro~jih ternarnega
diagrama za vzorce iz VD1
Figure 4: Numbers of inclusions in individual areas of the ternary
diagram for the sample LF2 collection
Slika 4: [tevilo vklju~kov v posameznih podro~jih ternarnega
diagrama za vzorce iz LF2
Figure 3: Numbers of inclusions in individual areas of the ternary
diagram for the sample LF1 collection
Slika 3: [tevilo vklju~kov v posameznih podro~jih ternarnega dia-
grama za vzorce iz LF1
Figure 2: Summary of 2-μm inclusions’ density during individual
technological operations, in all the monitored melts
Slika 2: Se{tevek vklju~kov velikih 2 μm med posameznimi tehno-
lo{kimi operacijami v vseh preiskovanih talinah
Figure 1: Summary of 1-μm inclusions’ density during individual
technological operations, in all the monitored melts
Slika 1: Se{tevek vklju~kov velikih 1 μm med posameznimi tehno-
lo{kimi operacijami v vseh preiskovanih talinah
Figure 6: Numbers of inclusions in individual areas of the ternary
diagram for the sample VD2 collection
Slika 6: [tevilo vklju~kov v posameznih podro~jih ternarnega diagra-
ma za vzorce iz VD2
The inclusions were divided into three groups using
the following scheme (Figure 7) to monitor their
chemical composition.
The inclusion compositions in individual melts and
stages were similar. We can claim that:
• Inclusions from Zone A in the ternary diagram are
practically absent in LF
• A small number of inclusions from Zone A occur in
VD
• The dominant inclusion group consists of the
inclusions from Zone B
• The main parts of the individual inclusions are Al,
Ca, and O
The measured and presented results suggest that their
effects, from the point of view of the number of
inclusions in both technologies, TS No. 01 and TS
No. 02, are similar. The result of the use of TS No. 02 is
only the smaller number of inclusions at the beginning of
the technological process in LF. When using TS No. 01
M. KORBÁ[ et al.: POSSIBILITIES FOR INCREASING THE PURITY OF STEEL IN PRODUCTION ...
784 Materiali in tehnologije / Materials and technology 48 (2014) 5, 781–786
Figure 8: The average number of impulses for the individual elements
and melts LF1
Slika 8: Povpre~no {tevilo impulzov za posamezne elemente iz taline
LF1
Figure 7: Scheme of the inclusions’ division into individual classes
Slika 7: Shemati~en prikaz razporeditve vklju~kov v posamezne
razrede
Figure 9: The average number of impulses for the individual elements
and melts LF2
Slika 9: Povpre~no {tevilo impulzov za posamezne elemente iz taline
LF2
Figure 12: The number of impulses in the course of the steel pro-
cessing LF1-VD2
Slika 12: [tevilo impulzov med izdelavo jekla (LF1-VD2)
Figure 11: The average number of impulses for the individual
elements and melts VD2
Slika 11: Povpre~no {tevilo impulzov za posamezne elemente iz taline
VD2
Figure 10: The average number of impulses for the individual
elements and melts VD1
Slika 10: Povpre~no {tevilo impulzov za posamezne elemente iz taline
VD1
and TS No. 02, the inclusion number and their size
become almost equal at the end of the processing in LF.
Inclusions larger than 4 μm occur very rarely and
accidentally.
The fundamental decrease in the number of inclu-
sions occurs only by 3/4 in the course of steel processing
in the LF. The processing in VD does not present a more
pronounced change in the number of smaller inclusions,
up to 2 μm. The occurrence of inclusions larger than
3 μm is negligible.
From the chemical inclusion composition in the LF
point of view, the inclusions from Zone A are not present
in any of the tested melts. A slight increase occurs only
in the course of the VD technology. The inclusions from
Zone B are the dominant ones. Their number from LF1
to VD 2 decreases in accordance with the above-pre-
sented conclusions. The Zone C is basically a half of
Zone B, where the inclusions of the type CaO.Al2O3
prevail. Only very fine inclusions of the CaO.Al2O3 type
and pure Al2O3, but also CaO, remain at the end of the
VD technology.
The results obtained with the SSE method were
based on an analysis of 2000 sparks. If any particles of a
different composition than those in the basic metallic
matrix, occurred in a given place, they are in the peaks of
the present elements and, then, their bonds could be
analysed. Each analysis of the relevant sample (disk-pin)
was performed in three places, always in such a way that
the relevant axis part of the sample was not affected and
the measurements were averaged.
The measured values of the analysed samples are
summarised graphically in Figures 8 to 11. The number
of sparks for the melt and the technological step are on
the x-axis (2000 for each measurement). The individual
colours mark inclusions based on the mentioned
elements. The number of impulses is on the y axis.3
The individual analyses and melt graphics, inclusion
types and technological steps of the production material
flow suggest that the mutual connections in the inclusion
volume and number, between individual samples for
melting or for individual technological steps LF1, LF2,
VD1 and VD2, seem difficult to define. Inclusions based
on elements like aluminium and oxygen (AlO) occur
more frequently than those based on aluminium, calcium
and oxygen (AlCaO). Inclusions based on manganese
and oxygen (MnO) or on manganese and sulphur (MnS)
were not detected because the optical spectrometer
Spectro Lab M10 was not equipped with a photo-
multiplier for manganese. MnS inclusions were thus
derived by calculation from the value S corr. There is a
relation MnS = S corr – CaS, where (S corr) is the total
content of inclusions incorporating sulphur and (CaS)
are the inclusions based on calcium and sulphur. The
analysis suggests that the number of inclusions (CaS)
means a significant amount. We assume that the
remaining number of inclusions based on sulphur has got
the form of MnS inclusions. The occurrence and the
number of inclusions based on aluminium and nitrogen
(AlN) and those based on titanium and nitrogen (TiN) is
not important. The occurrence of inclusions based on
aluminium and oxygen (CaO) and the sulphur content
show mostly the decreasing character. The relatively
high and unbalanced contents of inclusions based on
silicon and oxygen (SiO) are an oddity.
The graphical presentation of the development in the
total number of recorded impulses, in the course of
individual technological operations, in all the assessed
melts is in Figure 12. These data, in the sum of recorded
impulses, correspond to the number or, possibly, to the
volumes of steel inclusions. The assumption that the
curve of recorded impulses would decrease in correspon-
dence to the inclusion analysis made with the assistance
of roentgen spectral analysis and the process of steel
processing LF1-VD2 has not been quite proved.
Thanks to the graphic assessment in Figures 8 to 12,
which describe the number of impulses or the relative
frequency of the individual elements and melts, we can
say in summary that:
• It is very difficult to find any similarity in the number
of inclusions in individual samples in a melt and,
also, in the technological step
• The similarities of the trends’ decreases or increases
in individual inclusion types could be followed to a
melt or to a technological step
• The occurrence of AlCaO inclusions is not the
dominant detected bond
• The occurrence of AlO inclusions is more frequent,
when compared with AlCaO inclusions, but the trend
after VD2 is relatively balanced in most melts
• The number of AlO and AlCaO inclusions is lower,
when compared with VD1<VD2
• The occurrence of CaO inclusions is a dominant part
in all the technological steps and it is lower in
VD1<VD2
• The number of SiO inclusions is a dominant part in
all the technological steps
• When the occurrence of AlO inclusions exceeds the
value 0.1, the number of SiO inclusions is lower than
AlO
• The number of SiO inclusions shows a lower value in
VD1<VD2
• The analysis of the occurrence of the inclusions with
the bond MnS or MnO has not been performed
because the item for manganese was not purchased
for the device and we could not thus analyse it by the
SSE method
• The number of S corr inclusions mostly decreases in
VD1>VD2; the number of CaS inclusions logically
decreases
• The AlN and TiN inclusions do not practically occur
and are not created in the liquid stage
• The CaS inclusions make the dominant sulphur bond,
while the remaining sulphur amount was not
identified and it will possibly be connected with
inclusions of the MnS type
M. KORBÁ[ et al.: POSSIBILITIES FOR INCREASING THE PURITY OF STEEL IN PRODUCTION ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 781–786 785
• The number of CaO inclusions occurs in a relatively
large number
• The number of TiO inclusions is not negligible and it
is lower in VD1<VD2
• Some inclusions show in a relatively larger number
after vacuuming
The comparison of the two methods assessing some
steel purity factors with the assistance of roentgen
spectral analysis and the SSE method is suggested in
Table 1. In spite of the fact that the testing performance
identity of both methods was made as much as possible
for individual melts and samples, we have not achieved
comparable results. It is probable that the volume of the
analysed sample part makes an important difference
because the roentgen spectral analysis provides for the
relative point character of the detection in an area, while
the SSE method examines a larger volume of metal. The
comparison of the two methods shows that the difference
in the number of Al2O3-CaO inclusions, which prevail,
and the minimum number of CaO inclusions, detected by
the roentgen spectral analysis, is the most important. The
occurrence of these inclusions is more probable in every
technological step, when compared with the smaller
number of AlCaO inclusions and the large number of
occurring CaO inclusions detected by the SSE method.
4 CONCLUSIONS
The new method Single Spark Evaluation, or its
variant Spark – Data, which are available for modern
optical emission spectrometers, performs an analysis of
the chemical composition of steel. At the same time, they
may identify the inclusion compositions within condi-
tions of the steel production at Vítkovice Heavy Machi-
nery a.s. The advantage of this method is operability,
because the inclusion analysis takes place at the same
time as the analysis of the steel’s chemical composition.
The method has already been implemented in the
steelworks of VHM, but the correct interpretation of the
measured values and their utilisation within the operative
influencing of technological process of melts’ processing
is still discussed. The conclusions of the performed study
partly suggested a further verification process and the
practice, which might be utilised when interpreting the
measured values.
We recommend not comparing the SSE method with
the method of the roentgen spectral analysis. The
research of the utilisation of the SSE method should
continue in various technological production processes,
for example, in the condensing or diffusing de-oxyge-
nation technologies. The used de-oxygenation deter-
mines the type of an inclusion and its majority share in
steel. Similarly, the same should be applied in various
technological production processes. For example, in an
EAF and in the technology processing melts by facilities
of secondary metallurgy with the assistance of the
chemical heating by aluminium (VOD), or in production
processes in an EAF and when processing in a LF with
the shortened heating and in VD, possibly within other
tested combinations.
The current idea considers the fact that in such a way
the prepared technological processes would require the
creation of etalons - the steel status for the specific kind
of inclusions and the technological indicator of conta-
minated steel or pure steel. The number and the probable
size of inclusions will be determined by 3 Sigma, 7
Sigma and 9 Sigma. For example, the differential way
for the likely occurrence of fine inclusions, classified for
an assessment in accordance with the given technolo-
gical process, should be tested in a similar way and there
should be kinds of inclusions specified and determined
for the assessment in accordance with the given
technological process.
The study results suggest a chance for the Single
Spark Evaluation method’s utilisation in the steel work-
ing practice when determining the inclusions in an online
steel production process with secondary metallurgical
facilities LF or VD at Vítkovice Heavy Machinery a.s.
5 REFERENCES
1 G. Stolte, Secondary Metallurgy, Verlag Stahleisen GmbH, 2002, p.
216
2 M. Korbá{, M. Raclavský, L. ^amek, Change of quantity, sizes and
chemical composition of inclusions during treatment of steel in
EAF-LF-VD equipment, METAL 2012, Brno, 2012
3 M. Korbá{, L. ^amek, Evaluation of the results of steel purity
achieved in steel production using SSE method, In 22nd International
Conference on Metallurgy and Materials, METAL 2013, Brno, 2013
M. KORBÁ[ et al.: POSSIBILITIES FOR INCREASING THE PURITY OF STEEL IN PRODUCTION ...
786 Materiali in tehnologije / Materials and technology 48 (2014) 5, 781–786
Table 1: Comparison of some results measured by the Roentgen spec-
tral analysis and the SSE method
Tabela 1: Primerjava med rezultati, izmerjenimi z rentgensko spek-
tralno analizo in SSE-metodo
Roentgen spectral
analysis SSE method
Number of
inclusions
It decreases in the
course of processing
in the order
LF1-LF2-VD1-VD2
The number of
impulses does not
depend on the
technological step;
usually VD1<VD2
Mn Assessed Not assessed
MnS
The assessed Mn and
S, they occur at the
same time
The calculated
MnS = S corr – CaS
AlN, TiN Not assessed Assessed
AlCaO
The dominant
occurrence of Al2O3
and CaO at the same
time
Not the dominant
detected bond
CaO It does notpractically occur Relatively frequent
Ti
Titanium oxide
makes the dominant
element in some
inclusions
TiO occur practically
in all samples
A. AYDAY, M. DURMAN: EFFECT OF DIFFERENT SURFACE-HEAT-TREATMENT METHODS ...
EFFECT OF DIFFERENT SURFACE-HEAT-TREATMENT
METHODS ON THE SURFACE PROPERTIES OF AISI
4140 STEEL
VPLIV RAZLI^NIH TOPLOTNIH OBDELAV POVR[INE NA
LASTNOSTI POVR[INE JEKLA AISI 4140
Aysun Ayday, Mehmet Durman
Sakarya University, Faculty of Engineering, Department of Metallurgical and Materials Engineering, Sakarya 54187, Turkey
aayday@sakarya.edu.tr
Prejem rokopisa – received: 2013-10-01; sprejem za objavo – accepted for publication: 2013-11-19
In this study, the effect of different heat treatments on the hardness and wear properties of AISI 4140 steel was investigated.
Sample surfaces of AISI 4140 steel were modified with traditional induction heat treatment and the new treatment of
electrolytic-plasma hardening. The microstructural characteristics of surface-treated steel samples were examined with scanning
electron microscopy (SEM). The mechanical properties of the samples including the surface microhardness and modified-layer
thickness were also evaluated. The test results indicate that the wear resistance increases with the increasing modified-layer
hardness, due to the transformation of the martensitic structure. The electrolytic-plasma hardening is as suitable a technique for
improving the mechanical properties of AISI 4140 steel as the traditional hardening method.
Keywords: electrolytic-plasma treatment (EPT), induction, hardness, wear
V tej {tudiji je bil preiskovan vpliv razli~nih toplotnih obdelav na trdoto in obrabne lastnosti jekla AISI 4140. Povr{ina vzorcev
jekla AISI 4140 je bila obdelana z navadnim indukcijskim kaljenjem in z novim postopkom elektrolitskega utrjevanja s plazmo.
Zna~ilnosti mikrostrukture povr{insko obdelanih vzorcev jekla so bile preiskane z vrsti~nim elektronskim mikroskopom (SEM).
Ovrednotene so bile mehanske lastnosti vzorcev, vklju~no z mikrotrdoto povr{ine in debelino spremenjenega sloja. Rezultati
preizkusov ka`ejo, da se pove~uje odpornost proti obrabi s pove~anjem trdote modificiranega sloja zaradi pretvorbe v
martenzitno mikrostrukturo. Elektrolitsko utrjevanje s plazmo je enako primerna tehnika za izbolj{anje mehanskih lastnosti
jekla AISI 4140, kot obi~ajna metoda utrjevanja povr{ine.
Klju~ne besede: elektrolitska obdelava s plazmo (EPT), indukcija, trdota, obraba
1 INTRODUCTION
AISI 4140 steel (DIN 42CrMo4) is a widely used
material in industrial applications, such as automotive,
aerospace and manufacturing industries. However, the
parts required for mechanical working must have not
only high strength but also toughness, abrasion resi-
stance and corrosion resistance. AISI 4140 steel includes
Cr and Mo alloying elements that provide high harden-
ability and toughness after the surface treatments.1,2
Induction hardening (IH) is one of the most widely
employed treatments to improve component durability. It
determines, in a workpiece, a tough core with tensile
residual stresses and a hard surface layer with compres-
sive stresses, which have proved to be very effective in
extending a component’s fatigue life and wear resi-
stance.3,4 On the other hand, the conventional induction-
hardening treatment has some disadvantages: distortion,
change of coil, surface deformation, etc.5–7
Nowadays, the novel electrolytic-plasma treatment
(EPT) is one of the most promising surface-hardening
methods as an alternative to the induction hardening.
EPT has been successfully used to improve the hardness,
the wear resistance and the corrosion resistance of mate-
rials.8 When hardening is not necessary for the whole
surface or bulk of a material, EPT is a suitable method
for treating a specific location on the surface, like the
induction hardening.9 The standard mechanisms of the
two hardening treatments include two main steps: the
austenitization, during which the material is heated
above the critical temperature (but below the melting
point) to achieve the austenite formation and the quench-
ing or cooling down, when austenite is transformed into
martensite.10,11 Also, with EPT, the surface-treated zones
have an average depth of 0.1–10 mm, exhibiting an
increased hardness, improved resistance to wear, better
corrosion resistance and higher fatigue strength (e.g.,
piston ring groove, valve seat, compressor screws,
diesel-engine cylinder liners, gears, etc.).8,12
The present study aims at determining the wear resi-
stance of AISI 4140 steel treated with different surface-
hardening procedures: the induction hardening and the
electrolytic-plasma hardening. The modified samples
were characterized before and after the wear tests with
metallographic, SEM microscopy and microhardness
techniques.
2 EXPERIMENTAL PROCEDURE
AISI 4140 steel was selected for the present study
and Table 1 shows its chemical composition. The dia-
meter of cylindrical samples was 20 mm and the height
Materiali in tehnologije / Materials and technology 48 (2014) 5, 787–790 787
UDK 621.78:620.178 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(5)787(2014)
was 10 mm. All the samples were modified with EPT
and induction hardening. The EPT voltage, heating and
cooling times were 310–260 V, 3 s and 3 s, respectively,
and a constant thermal cycle was selected. The
specimens were induction hardened with 30 kW and 300
kHz for 6 s and quenched in water. The sample codes
and EPT parameters are listed in Table 2.
The morphology of the modified layers was inve-
stigated with a scanning electron microscope (SEM Jeol,
JSM 6060-LU). The hardness measurements were
conducted on the cross-sections of the samples with a
Vickers microhardness tester. The test load was 100 g for
the hardness measurements at the cross-sections. The
hardness value was obtained by averaging the results of
three measurements. The temperature distribution of the
samples from the plasma-treated side to the internal side
was investigated via thermocouples during the process.
The surface temperature data were collected from the
system with the aid of a computer-data-acquisition sys-
tem. The volume fractions of different phases were
measured with an image analyzer on the metallographic
sections.
Table 1: Chemical composition (w/%) of AISI 4140 steel
Tabela 1: Kemijska sestava (w/%) jekla AISI 4140
C Si Mn Cr Mo P S Fe
0.41 0.21 0.8 0.95 0.18 0.025 0.027 Balance
Table 2: Treatment parameters
Tabela 2: Parametri obdelave
Parame-
ter code
Electro-
lytic
solution
Heating
(V)
Cooling
(V)
Total
time
Thermal
cycle
Untreated Original AISI 4140
EPT–6 Na2CO3;12 % 310 260
(3 s and 3 s)
× 6 = 36 s 6
IH 30 kW, 300 kHz, 6 s
The wear tests were performed on the original AISI
4140 and on the modified specimens to determine the
best process and parameters. All the wear tests were
carried out under dry sliding conditions at room tem-
perature using a ball-on-disc (CSM tribometer), friction-
and wear-test machine. The counterpart was an Al2O3
ball (Ø = 6 mm) used in accordance with DIN 50 324
and ASTM G 99-95a. The tests were performed with a
nominal load of 3 N and a sliding speed of 0.10 m/s for
the total sliding distance of 200 m.
3 RESULTS AND DISCUSSION
Figure 1 shows the features of the modified layers,
observed on the cross-sections of the EPT and induction
materials. During both modification treatments, austenite
transforms completely or partially into martensite on the
surface called the hardened zone (HZ). An amount of the
retained austenite may be present in this region.13 In the
neighborhood of the HZ with the base material, a narrow
heat-affected zone was observed, consisting of marten-
site, bainite and some traces of the initial pearlitic
structure. These are the most probable structures accord-
ing to13,14. The base-material microstructure is composed
of ferrite and pearlite. All the experimental data showing
the results of the mechanical tests are given in Table 3.
The maximum layer thicknesses of the modified surfaces
were 5.2 mm for the EPT and 2.1 mm for the IH (Table
3). An effective increase in the hardness was observed
after the surface heat treatments. The EPT-sample
hardness was higher than that of the induction-hardened
samples, as shown in Table 3. The hardness of the modi-
fied layers depends on the martensitic transformation,
and its high hardness is usually explained with the
martensite volume fraction (MVF).15,16 The volume frac-
tions of martensite are given in Table 3. The martensite
volume fraction of the EPT and induction hardening
A. AYDAY, M. DURMAN: EFFECT OF DIFFERENT SURFACE-HEAT-TREATMENT METHODS ...
788 Materiali in tehnologije / Materials and technology 48 (2014) 5, 787–790
Figure 1: Cross-sectional appearance of modified samples: a) EPT, b)
IH
Slika 1: Prerez modificiranega vzorca: a) z EPT, b) z IH
Table 3: General results for modified samples
Tabela 3: Rezultati na modificiranih vzorcih
Thick-
ness
(mm)
Hard-
ness
(HV0.1)
Surface
tempe-
rature
(°C)
MVF
(%)
Friction
coeffi-
cient
Wear rate
(mm3/
(N m))
Untreated – 200 – – 0.2 9.06 E–05
EPT 5.2 930 835 47.4 0.203 4.66 E–05
IH 2.1 810 800 39.4 0.377 4.96 E–05
were 47.4 % and 39.4 %, respectively. As shown in
Table 3, the EPT volume fraction of martensite is higher
than in the case of induction because of the surface
temperature.
Figure 2 shows the longitudinal hardness profiles for
both specimen types. A significant surface hardening
was observed on all the modified samples. The hardness
of the AISI 4140 steel was 10 HRC. The surface hard-
ness of the specimens increased considerably after the
surface-hardening treatments. The surface-hardening
behavior of the steel depended on its microstructure
constituents, i.e., the martensite (or bainite) morphology
and volume fraction with reference to12,15,17. The EPT-
modified sample shows the maximum hardness in the
center, corresponding to the martensitic structure. The
hardness along the EPT track decreases rapidly for about
20 mm from its center, while the induction-hardened
sample has a single hardness value. This fact can be
explained by examining the shape and the intensity
profile of the EPT. In addition, the exposure (heating)
time is lower at the border of the EPT tracks due to the
elliptical nozzle shape. The combination of these two
effects results in significantly lower temperatures
reached at the external limits of the modified zone, at
which the austenization does not take place, so no hard-
quenched martensite can be achieved.
Figure 3 shows the wear surfaces of the original
AISI 4140, EPT and induction-modified samples, tested
at a load of 3 N. The modified samples have quite
smooth surfaces with shallow abrasive-wear scars due to
the high hardness of the samples. The original AISI 4140
steel was tested under a similar wear-test condition. In
this case, the plastic deformation caused due to a low
surface hardness is obvious, as shown in Figure 3. The
lowest wear rate was obtained for the EPT-modified sam-
ple followed by the induction-hardened and untreated
AISI 4140 samples (Table 3). This is due to the micro-
structure of the sample, consisting of martensite (47.4 %
MVF) and having the highest hardness among the tested
samples.
4 CONCLUSIONS
In this study, the following results were observed:
AISI 4140 steel surfaces were successfully hardened
with the electrolytic-plasma technique and the traditional
induction-hardening method. The maximum surface
hardness was obtained with the EPT. The maximum
microhardness value increased from 200 HV0.1 to 930
HV0.1 during the experiments. After the wear experi-
ments, the EPT-modified samples had the lowest wear
rate, showing almost the same wear resistance as ob-
served for the induction-hardened samples. The un-
treated samples, as expected, had the highest wear rate.
The hardness of the surface is a very important factor
with respect to the wear rate. It is a result of the micro-
structure of the modified samples, which consisted of
martensite.
5 REFERENCES
1 S. Choo, S. Lee, M. G. Golkovski, Materials Science and Engi-
neering A, 293 (2000), 56–70
2 T. Miokovic, V. Schulze, O. Vohringer, D. Lohe, Acta Materialia, 55
(2007), 589–599
3 Y. Totik, R. Sadeler, H. Altun, M. Gavgali, Materials and Design, 24
(2003), 25–30
4 J. Yi, M. Gharghouri, P. Bocher, M. Medraj, Chemistry and Physics,
142 (2013), 248–258
A. AYDAY, M. DURMAN: EFFECT OF DIFFERENT SURFACE-HEAT-TREATMENT METHODS ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 787–790 789
Figure 2: Surface-hardness profiles of untreated AISI 4140, EPT and
IH samples
Slika 2: Profil trdote neobdelanega AISI 4140, EPT in IH vzorca
Figure 3: SEM micrographs of worn surfaces of untreated AISI 4140, EPT and IH samples
Slika 3: SEM-posnetki obrabljene povr{ine neobdelanega AISI 4140, EPT in IH vzorca
5 W. M. Sten, J. Mazumder, Laser Material Processing, 4th ed., Sprin-
ger, Berlin, 2010, 297
6 ASM Handbook, Surface Engineering, vol. 5, 10th ed., ASM Inter-
national, 1994
7 A. K. Sinha, Physical Metallurgy Handbook, vol. 16-2, McGraw-
Hill, 2003
8 Yu. N. Tyurin, A. D. Pogrebnyak, Surface and Coatings Technology,
142–144 (2001), 293–299
9 C. Ye, S. Suslov, B. J. Kim, E. A. Stach, G. J. Cheng, Acta Mate-
rialia, 59 (2011), 1014–1025
10 S. A. Jenabali Jahromi, A. Khajeh, B. Mahmoudi, Materials and De-
sign, 34 (2012), 857–862
11 N. S. Bailey, W. Tan, Y. C. Shin, Surface & Coatings Technology,
203 (2009), 2003–2012
12 A. Ayday, Surface Properties Improvement of Nodular Cast Iron Mo-
dified by Electrolytic Plasma Technology, PhD thesis, Sakarya Uni-
versity, 2013
13 C. Soriano, J. Leunda, J. Lambarri, V. G. Navas, C. Sanz, Applied
Surface Science, 257 (2011), 7101–7106
14 C. T. Kwok, K. I. Leong, F. T. Cheng, H. C. Man, Materials Science
and Engineering A, 357 (2003), 94–103
15 Y. Sahin, M. Erdogan, M. Cerah, Wear, 265 (2008), 196–202
16 A. Ayday, M. Durman, Mater. Tehnol., 47 (2013) 2, 177–180
17 A. Zare, A. Ekrami, Materials Science and Engineering A, 530
(2011), 440–445
A. AYDAY, M. DURMAN: EFFECT OF DIFFERENT SURFACE-HEAT-TREATMENT METHODS ...
790 Materiali in tehnologije / Materials and technology 48 (2014) 5, 787–790
Y. YILDIZ et al.: PREPARATION AND APPLICATION OF POLYMER INCLUSION MEMBRANES (PIMs) ...
PREPARATION AND APPLICATION OF POLYMER
INCLUSION MEMBRANES (PIMs) INCLUDING
ALAMINE 336 FOR THE EXTRACTION OF METALS
FROM AN AQUEOUS SOLUTION
PRIPRAVA IN UPORABA MEMBRANE IZ POLIMERA (PIM) IN
ALAMINA 336 ZA LO^ENJE KOVIN IZ VODNIH RAZTOPIN
Yasemin Yildiz1, Aynur Manzak1, Büºra Aydýn1, Osman Tutkun2
1Department of Chemistry, Sakarya University, Sakarya, Turkey
2Beykent University, Department of Chemical Engineering, Engineering and Architecture Faculty, Istanbul, Turkey
manzak@sakarya.edu.tr
Prejem rokopisa – received: 2013-10-12; sprejem za objavo – accepted for publication: 2013-11-12
Polymer inclusion membranes (PIMs) present an attractive approach for the separation of metals from an aqueous solution. The
present study is about the application of Alamine 336 as an ion carrier in PIMs. The separation of copper (II), cobalt (II), nickel
(II) and cadmium (II) from aqueous solutions with polymer inclusion membranes was investigated. PIMs are formed by casting
a solution containing a carrier (extractant), a plasticizer and a base polymer, such as cellulose tri-acetate (CTA) or poly(vinyl
chloride) (PVC), to form a thin, flexible and stable film. Several important transport parameters such as the type and amount of
the plasticizer, the type of the stripping solution, the thickness of the membrane, the pH of the acid in the donor phase and the
concentration of the base in the acceptor phase are discussed. The membrane was characterized to obtain information regarding
its composition using AFM, FT-IR and SEM.
Keywords: polymer inclusion membranes, plasticizer, extractant, thickness of membrane
Membrane, ki vsebujejo polimere (PIM), so zanimive za lo~enje kovin iz vodnih raztopin. Prikazana je {tudija uporabe Alamina
336 kot nosilca ionov v PIM. Preiskovano je bilo lo~enje bakra (II), kobalta (II), niklja (II) in kadmija (II) iz vodnih raztopin z
membrano s polimeri. PIM je bila izdelana z ulivanjem raztopine z nosilcem (ekstraktantom), z meh~alcem in osnovo iz poli-
mera, kot je celuloza-tri-acetat (CTA) ali polivinil klorid (PVC), da je nastala tanka, gibljiva plast. Razlo`enih je ve~ pomemb-
nih transportnih parametrov, kot so dele` meh~alca, vrsta raztopine za snemanje, debelina membrane, pH kisline v donorski fazi
in koncentracija baze v aceptorski fazi. Izvr{ene so bile preiskave z AFM, FT-IR in SEM, da bi dobili podatke o sestavi
membrane.
Klju~ne besede: membrane s polimerom, meh~alec, ekstraktant, debelina membrane
1 INTRODUCTION
The separation of metals from sulphate and chloride
media has been of practical interest to the researchers.
Solvent extraction is a well-established technology used
for the production of metals from a relatively concen-
trated feed. However, industrial diluent effluents pose an
important challenge as the solvent-extraction technique
is not cost effective for the separation of metals from a
dilute solution1.
Recently, the supported liquid membrane (SLM)
extraction has been emerging as an alternative to the
conventional solvent extraction due to its advantages
such as high selectivity, operational simplicity, low
solvent inventory, low energy consumption, zero effluent
discharge, and a combination of extraction and stripping
in a single unit2,3. Currently, considerable attention is fo-
cused upon polymer inclusion membranes (PIMs)4. Their
specific advantages are an effective carrier immobili-
zation, easy preparation, versatility, stability, good che-
mical resistance and better mechanical properties than in
the case of SLM5. The large surface-area-to-volume ratio
exhibited by PIMs gives them the potential to be used in
nuclear and harmful-metal waste remediation on an
industrial scale. They consist of a polymer providing the
mechanical strength, a carrier molecule that effectively
binds and transports the ions across the membrane, and a
plasticizer that provides elasticity and acts as the solvent,
in which the carrier molecule can diffuse. PIMs are
formed by casting a solution containing a carrier (extrac-
tant), a plasticizer and a base polymer, such as cellulose
tri-acetate (CTA) or poly(vinyl chloride) (PVC), to form
a thin, flexible and stable film6.
The choice of different constituents of the membrane
is crucial to ensure its separation efficiency, so it is
important to investigate the effect of different compo-
nents on the extraction and transport of the target
species. Among the polymers used to form a gel-like
network that entraps the carrier and plasticizer/modifier,
poly(vinyl chloride) (PVC) and cellulose triacetate
(CTA) are most frequently encountered7.
Examples of such membranes are those containing
only PVC and Aliquate 336 that have been used success-
fully for the transport of both metallic (e.g., Cd (II) and
Cu (II)8 and non-metallic (e.g., thiocyanate)9 ionic spe-
cies. Moreover, Konczyk et al.10 have used Aliquate 336
Materiali in tehnologije / Materials and technology 48 (2014) 5, 791–796 791
UDK 678.7 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(5)791(2014)
as a plasticizer in a PIM system containing D2EHPA as
the carrier for the removal of Cr (III).
The present study focuses on the application of Ala-
mine 336 as an ion carrier in PIMs and deals with the
selective separation of Co, Cd, Ni and Cu ions from an
acidic media into an NH4SCN aqueous solution. Amines
were used to extract the metal ions. The amine extraction
chemistry of thiocyanate complexes was investigated by
Sanuki et al.11
2 EXPERIMENTAL WORK
2.1 Materials
All the reagents used were of analytical grade.
Cellulose triacetate (CTA), 2-nitrophenyl pentyl ether
(NPPE) and 2-nitrophenyl octyl ether (NPOE) were
obtained from Fluka. Tributyl phosphate (TBP), dichlo-
romethane, CoCl2 · 6H2O, NiSO4 · 6H2O, 3CdSO4 ·
8H2O, CuSO4 · 5H2O, acetic acid, NaOH, ammonium,
triethanolamine, NH4SCN and Alamine 336 were of
analytical grade (Merck) and all the stock solutions were
prepared by dissolving the salts in distilled water.
2.2 Preparation of PIMs
PIMs were prepared in accordance with the casting
solution. CTA (480 mg) was dissolved in 70 mL of
dichloromethane at room temperature. In the following
step 0.1–0.5 mL of 2-NPPE was added into the solution.
After stirring, the carrier (Alamine 336 and TBP) was
added and the solution was stirred for 6 h to obtain a
homogenous solution. The solvent of this mixed solution
was allowed to slowly evaporate in a square glass con-
tainer (24 cm × 24 cm). The organic solvent was allowed
to evaporate overnight at room temperature. After the
evaporation of the solvent, a few drops of cold and
distilled water swirled on the top of the polymer film.
Afterwards, the membrane was peeled out of the con-
tainer. The average thickness of the membrane was
determined as 25 μm with a digital micrometer (Salu
Tron Combi-D3).
2.3 PIM transport experiment
The prepared polymeric film was sandwiched bet-
ween two glass cells. The transport of metal ions across
the PIM from the aqueous solutions was studied by using
a two-compartment permeation cell made from Pyrex
glass, having flat-sheet membranes with the 12.56 cm2
area (A), as shown schematically in Figure 1.
The volumes of both the aqueous feed and the strip
phases were 250 mL. The feed solutions were prepared
by adding cobalt, nickel, cadmium and copper salts to
study the effect of the feed composition. Ammonium
thiocyanate (NH4SCN) was added into the feed mixture
to increase the selectivity of cobalt against nickel. 1 M
acetic acid/1 M sodium acetate buffer was used to main-
tain the desired feed pH. A stripping solution containing
1 M NH3 + 1 M TEA was selected as the stripping-phase
mixture. The feed and stripping phases were mechani-
cally stirred at the desired mixing speed of (20 ± 1) °C to
avoid the concentration polarization conditions at the
membrane interfaces and in the bulk of the solutions.
During the PIM-transport experiments, the samples of
the feed and strip phases (about 1 mL) were periodically
removed for a determination of the metal concentration
with ICP-OES.
3 RESULTS AND DISCUSSION
3.1 Plasticizer type and concentration
The nature of the plasticizer used to form the mem-
brane is also a key parameter to consider. Plasticizers are
organic compounds incorporating a hydrophobic alkyl
backbone and one or several highly solvating polar
groups. They are added to hard, stiff plastics to make
them softer and more flexible. The softening action of
the plasticizers, plasticization, is usually attributed to
their ability to reduce the intermolecular attractive forces
between the polymer chains. For this reason, it is
anticipated that in PIMs the presence of these com-
pounds may also influence the mobility of membrane
components, the degree of interaction between different
constituents of the membrane and the characteristics of
the polymeric medium7. A low plasticizer concentration
may cause more rigid and brittle membranes. So, it is not
preferred4. The minimum plasticizer concentration varies
widely depending on both the plasticizer and the base
polymer. The influence of the plasticizer nature on the
Cd2+, Co2+, Ni2+, and Cu2+ transport through PIMs with
different plasticizers, i.e., 2-Nitrophenyl octyl ether
(NPOE) and 2-nitrophenyl pentyl ether (NPPE) was
tested. Copper was precipitated in the feed phase. Nickel
was not transferred to the stripping solution. The cobalt
and cadmium ions in the acidic feed solutions reacted
with the excess NH4SCN, whereas in the case of nickel
ions, they hardly formed a thiocyanate complex12,13.
The results obtained for the Cd2+ and Co2+ ion
transport with different concentrations of the plasticizers
Y. YILDIZ et al.: PREPARATION AND APPLICATION OF POLYMER INCLUSION MEMBRANES (PIMs) ...
792 Materiali in tehnologije / Materials and technology 48 (2014) 5, 791–796
Figure 1: Schematic diagram of the experimental apparatus
Slika 1: Shema naprave za preizkuse
in the PIMs are shown in Figures 2 and 3. For NPPE,
this concentration can be in the range of up to 0.2 mL (w
= 27 %) (Figure 2). Above this upper limit the mass
transport diminishes.
The results obtained for the Cd2+ and Co2+ ion trans-
port with different types of plasticizers in the PIMs are
shown in Figures 4 and 5. 2-Nitrophenyl pentyl ether
(NPPE) is the most frequently used plasticizer in PIMs
due to its high dielectric constant that enhances the
membrane permeability. A decrease in the permeability,
together with an increase in the plasticizer content, is
probably related to a reduction in the viscosity of the
medium4.
The recovery factor (RF) of the metal ions from the
feed phase into the stripping phase is given by:
RF
C C
C
I
I
=
−
⋅100% (1)
where C is the metal-ion concentration in the feed phase
at some given time and Ci is the initial metal-ion con-
centration in the feed phase. Recovery factors (RF) for
different plasticizers are shown in Table 1.
Table 1: Effect of the plasticizer type on the cobalt and cadmium
transport
Tabela 1: Vpliv vrste meh~alca na prenos kobalta in kadmija
Plasticizer type RF (Co) RF (Cd)
Y. YILDIZ et al.: PREPARATION AND APPLICATION OF POLYMER INCLUSION MEMBRANES (PIMs) ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 791–796 793
Figure 4: Effect of the plasticizer type on the cadmium transport (feed
phase: 100 mg/L Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100 mg/L
Cu2+; feed stirring speed: 1200 r/min; strip-phase stirring speed: 1200
r/min; strip solution: 1 M NH3 + 1 M TEA; complex reagent
(NH4SCN): 0.5 mol/l; temp.: 20 °C; feed solution pH: 4)
Slika 4: Vpliv vrste meh~alca na prenos kadmija (raztopina: 100 mg/L
Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100 mg/L Cu2+; hitrost me{anja
raztopine: 1200 r/min; hitrost me{anja v fazi traku: 1200 r/min; raz-
topina traku: 1 M NH3 + 1 M TEA; kompleksni reagent (NH4SCN):
0.5 mol/l; Temp.: 20 °C; raztopina pH: 4)
Figure 2: Effect of the NPPE concentration on the cadmium extrac-
tion (feed phase: 100 mg/L Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100
mg/L Cu2+; feed stirring speed: 1200 r/min; strip-phase stirring speed:
1200 r/min; strip solution: 1 M NH3 + 1 M TEA; complex reagent
(NH4SCN): 0.5 mol/l; temp.: 20 °C; feed solution pH: 4)
Slika 2: Vpliv koncentracije NPPE na ekstrakcijo kadmija (iz
raztopine: 100 mg/L Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100 mg/L
Cu2+; hitrost me{anja raztopine: 1200 r/min; hitrost me{anja v fazi
traku: 1200 r/min; raztopina traku: 1 M NH3 + 1 M TEA; kompleksni
reagent (NH4SCN): 0.5 mol/l; Temp.: 20 °C; raztopina pH: 4)
Figure 5: Effect of the plasticizer type on the cobalt transport (feed
phase: 100 mg/L Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100 mg/L
Cu2+; feed stirring speed: 1200 r/min; strip-phase stirring speed: 1200
r/min; strip solution: 1 M NH3 + 1 M TEA; complex reagent
(NH4SCN): 0.5 mol/l; temp.: 20 °C; feed solution pH: 4)
Slika 5: Vpliv vrste meh~alca na prenos kobalta (raztopina: 100 mg/L
Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100 mg/L Cu2+; hitrost me{anja
raztopine: 1200 r/min; hitrost me{anja v fazi traku: 1200 r/min; raz-
topina traku: 1 M NH3 + 1 M TEA; kompleksni reagent (NH4SCN):
0.5 mol/l; temp.: 20 °C; raztopina pH: 4)
Figure 3: Effect of the NPPE concentration on the cobalt extraction
(feed phase: 100 mg/L Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100
mg/L Cu2+; feed stirring speed: 1200 r/min; strip-phase stirring speed:
1200 r/min; strip solution: 1 M NH3 + 1 M TEA; complex reagent
(NH4SCN): 0.5 mol/l; temp.: 20 °C; feed solution pH: 4)
Slika 3: Vpliv koncentracije NPPE na ekstrakcijo kobalta (raztopina:
100 mg/L Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100 mg/L Cu2+;
hitrost me{anja raztopine: 1200 r/min; hitrost me{anja v fazi traku:
1200 r/min; raztopina traku: 1 M NH3 + 1 M TEA; kompleksni
reagent (NH4SCN): 0.5 mol/l; temp.: 20 °C; raztopina pH: 4)
NPOE 59 25
NPPE 81 46
3.2 Effect of the stripping-solution type
In general, metallic ions extracted by amines can be
stripped from the protonated amine with the removal of a
proton using neutral or alkaline solutions. 1 M ammonia
and 1 M triethanol amine solution mixtures were used as
the reagents to strip and separate the cobalt and cadmium
from the membrane phase to the aqueous phase.
3.3 Membrane characteristics
One important aspect of PIMs is the microstructure
of the membrane materials, which determines the distri-
bution of the carriers in the polymer matrix and ultima-
tely affects the membrane transport efficiency. Conse-
quently, a considerable research effort was devoted to
clarifying this issue. While a variety of surface-characte-
rization techniques were employed in these studies,
scanning electron microscopy (SEM) and atomic force
microscopy (AFM) were most frequently used. The
results obtained from the SEM and AFM studies consi-
stently indicate a remarkable influence of the polymeric
composition on the membrane morphology.
The membrane was characterized to obtain informa-
tion regarding its composition using AFM (Figure 6),
SEM (Figure 7) and FT-IR (Figure 8).
The AFM technique was used to characterize the sur-
face morphology of the prepared membranes. The AFM
picture of the PIM formed with CTA + NPPE + TBP +
Alamine 336 is shown in Figure 6. The surface morpho-
logy of the membrane shows a rough surface. These
regions may have occurred because of either a different
speed of the solvent vaporization14, 15 or the membrane
having a porous structure where the pores were filled by
NPPE or NPPE + Alamine 336 + TBP16,17.
Although both SEM and AFM techniques are versa-
tile and can provide a good image of the membrane
surface and, to some degree, of the membrane interior
structure, to date, the studies employing these techniques
have not been able to clearly elucidate the distribution of
the carrier and the plasticizer within the membrane. Con-
sequently, more advanced material-characterization tech-
niques have been attempted.4
In order to investigate the absorption bands of the
constituents of the membranes containing CTA + NPPE+
Alamine 336 + TBP, FTIR was performed as shown in
Figure 8. The bands at 2986 cm–1 and 2936 cm–1 were
attributed to the stretching vibration of C–H in -CH2 and
-CH3. The absorption at 1750 cm–1 was assigned to the
stretching vibration of C=O in CTA.
The fingerprint region of the spectra becomes com-
plicated because of the P–O, C–O and C–N vibrations.
These three vibrations are absorbed in the same region.
For example, the peaks in the 1250 cm–1 and 1100 cm –1
region appear in both CTA and TBP and they overlap
completely. The expected peaks in the spectra appear in
almost the same region as in the case of pure compo-
nents like CTA, Alamine 336 and TBP. This indicates
that these four compounds do not form any new covalent
interactions, but only secondary interactions like hydro-
gen bonding or electrostatic interactions18.
Consequently, the analysis and comparison of the
obtained spectra revealed that all the membrane consti-
tuents remained as pure components inside the membra-
ne17,18. The surface of the films shows a good uniformity
and the absence of cracks indicates a good regularity of
the membranes as shown in Figure 7.
3.4 Membrane thickness
The investigated membrane thickness was 20 μm to
45 μm, shown in Figures 9 and 10. The best recovery
factor (RF) was obtained with a thickness of 25 μm, with
81 % in the feed phase over 5 h as shown in Table 2. The
Y. YILDIZ et al.: PREPARATION AND APPLICATION OF POLYMER INCLUSION MEMBRANES (PIMs) ...
794 Materiali in tehnologije / Materials and technology 48 (2014) 5, 791–796
Figure 7: CTA + NPPE + Alamine 336 + TBP, SEM image
Slika 7: SEM-posnetek CTA + NPPE + Alamin 336 + TBP
Figure 6: CTA + NPPE + Alamine 336 + TBP, AFM image
Slika 6: AFM-posnetek CTA + NPPE + Alamin 336 + TBP
Figure 8: CTA + NPPE + Alamine 336 + TBP, FT-IR
Slika 8: FT-IR-posnetek CTA + NPPE + Alamin 336 + TBP
optimum membrane thickness was 25 μm. As the mem-
brane thickness increased, the extraction would decrease.
As shown in19 thinner membranes exhibiting high
permeability are formed. However, the thinnest mem-
branes only partly allow high permeability due to a
decrease in the extractant content limiting the transport
efficiency.
As shown in reference20 the flux decreased linearly
with the membrane thickness. This is unambiguous evi-
dence that the slow step in the transport process repre-
sents the migration through the membrane and not a
decomplexation from the carrier.
4 CONCLUSIONS
With the use of Alamine 336 and TBP as the carriers,
the competitive transport of metal ions shows the prefe-
rential selectivity order: Co (II) > Cd (II). The transport
facilitated through the polymer inclusion membranes
containing Alamine 336 and TBP was found to be an
effective method for separation and recovery of cobalt
(II) and cadmium (II) from aqueous solutions. Copper
was precipitated in the feed phase. Nickel was not
transferred to the stripping solution. The recovery factor
for the cobalt ions was over 87 % over a period 6 h.
Acknowledgement
The financial support of this work, provided by the
scientific research commission of Sakarya University
(BAPK), Project No: 2010-02-04-025, is gratefully
acknowledged.
5 REFERENCES
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Y. YILDIZ et al.: PREPARATION AND APPLICATION OF POLYMER INCLUSION MEMBRANES (PIMs) ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 791–796 795
Figure 10: Effect of the membrane thickness on the cobalt transport
(feed phase: 100 mg/L Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100
mg/l Cu2+; feed stirring speed: 1 200 r/min; strip-phase stirring speed:
1200 r/min; strip solution: 1 M NH3 + 1 M TEA; complex reagent
(NH4SCN): 0.5 mol/l; temp.: 20 °C; feed solution pH: 4)
Slika 10: Vpliv debeline membrane na prenos kobalta (raztopina: 100
mg/l Co2+, 100 mg/l Ni2+, 100 mg/l Cd2+, 100 mg/l Cu2+; hitrost
me{anja raztopine: 1 200 r/min; hitrost me{anja v fazi traku: 1 200
r/min; raztopina traku: 1 M NH3 + 1 M TEA; kompleksni reagent
(NH4SCN): 0.5 mol/l; temp.: 20 °C; raztopina pH: 4)
Figure 9: Effect of the membrane thickness on the cadmium transport
(feed phase: 100 mg/L Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100
mg/L Cu2+; feed stirring speed: 1200 r/min; strip-phase stirring speed:
1200 r/min; strip solution: 1 M NH3 + 1 M TEA; complex reagent
(NH4SCN): 0.5 mol/l; temp.: 20 °C; feed-solution pH: 4)
Slika 9: Vpliv debeline membrane na prenos kadmija (raztopina: 100
mg/L Co2+, 100 mg/L Ni2+, 100 mg/L Cd2+, 100 mg/L Cu2+; hitrost
me{anja raztopine: 1200 r/min; hitrost me{anja v fazi traku: 1200
r/min; raztopina traku: 1 M NH3 + 1 M TEA; kompleksni reagent
(NH4SCN): 0.5 mol/l; temp.: 20 °C; raztopina pH: 4)
Table 2: Effect of the membrane thickness on the cobalt and cadmium
transport
Tabela 2: Vpliv debeline membrane na prenos kobalta in kadmija
Membrane thickness
(μm) RF (Co) RF (Cd)
20 39 15
25 81 46
45 54 15
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796 Materiali in tehnologije / Materials and technology 48 (2014) 5, 791–796
D. HAUSEROVA et al.: ACCELERATED CARBIDE SPHEROIDISATION AND REFINEMENT (ASR) ...
ACCELERATED CARBIDE SPHEROIDISATION AND
REFINEMENT (ASR) OF THE C45 STEEL DURING
CONTROLLED ROLLING
POSPE[ENA SFEROIDIZACIJA IN UDROBNJENJE KARBIDOV
(ASR) PRI KONTROLIRANEM VALJANJU JEKLA C45
Daniela Hauserova, Jaromir Dlouhy, Zbysek Novy
COMTES FHT, Prumyslova 995, 334 41 Dobrany, Czech Republic
daniela.hauserova@comtesfht.cz
Prejem rokopisa – received: 2013-10-14; sprejem za objavo – accepted for publication: 2013-11-06
Current industry trends include the search for cost- and energy-saving procedures and technologies. A new process has been
discovered recently, which allows a significant refinement of ferrite grains and a carbide spheroidisation in a shorter time than in
the case of conventional heat-treatment techniques. During this newly-developed ASR-based (accelerated spheroidisation and
refinement) plastic deformation an accelerated spheroidisation and a refinement due to the heat treatment in the vicinity of the
A1 temperature occur.
Controlled rolling enables a production of the materials with a fine microstructure and better mechanical properties than
conventional production processes. Accelerated carbide spheroidisation and refinement (ASR) is aimed to produce steel
workpieces with a microstructure consisting of a fine-grained ferrite matrix and globular carbide particles. In carbon steels, this
microstructure has higher yield strength and toughness than the conventional ferritic-pearlitic microstructure.
The presented paper describes the effect of the ASR process on the C45 steel. The pearlite morphology was influenced by
forming it at the temperatures around critical temperature A1 and an accelerated carbide-particle spheroidisation was achieved.
The deformation increases the dislocation density and enhances the diffusion rate. Cementite globules form rapidly, within
seconds or minutes at the most.
Keywords: accelerated spheroidisation, refinement, rolling, C45 steel
Sedanje usmeritve industrije vklju~ujejo tudi iskanje stro{kovno in energijsko ugodnej{ih postopkov in tehnologij. Razvit je bil
nov postopek, ki omogo~a znatno udrobnjenje zrn ferita in sferoidizacijo karbidov v kraj{em ~asu v primerjavi s konven-
cionalnimi tehnikami toplotne obdelave. Pri tej novo razviti plasti~ni deformaciji, na kateri temelji ASR (pospe{ena
sferoidizacija in udrobnjenje), se pri plasti~ni deformaciji pojavi pospe{ena sferoidizacija in udrobnjenje med toplotno obdelavo
v bli`ini temperature A1.
Kontrolirano valjanje omogo~a izdelavo materialov z drobno mikrostrukturo in z bolj{imi mehanskimi lastnostmi kot pri
navadnih proizvodnih procesih. Namen pospe{ene sferoidizacije karbidov in udrobnjenja zrn (ASR) je izdelava jekla z mikro-
strukturo iz drobnih zrn ferita in globularnih karbidnih zrn. Pri ogljikovih jeklih ima ta mikrostruktura vi{jo mejo te~enja in
ve~jo `ilavost kot navadna feritno-perlitna mikrostruktura.
^lanek opisuje u~inek ASR-procesa na jeklo C45. Na morfologijo perlita se vpliva s preoblikovanjem pri temperaturah okrog
kriti~ne temperature A1, in s tem je dose`ena sferoidizacija karbidnih delcev. Deformacija pove~uje gostoto dislokacij in pove~a
hitrost difuzije. Globularni cementit najve~krat nastane v nekaj sekundah ali minutah.
Klju~ne besede: pospe{ena sferoidizacija, udrobnjenje, valjanje, jeklo C45
1 INTRODUCTION
The current processes leading to a carbide-particle
spheroidisation rely on diffusion of carbon in a work-
piece heated to a temperature close to or slightly below
Ac1.1 Diffusion-based processes of this type are usually
time-consuming and the times of up to tens of hours2
make this type of annealing a very expensive heat-treat-
ment process. During annealing, softening processes
occur in the microstructure and, in some cases, a reco-
very and a recrystallization also take place.3 The strength
and hardness of the steel workpiece decline, whereas its
ductility and plastic-deformation capability are in-
creased. The newly-designed and patented thermo-
mechanical process brings a several-fold reduction in the
processing time and cost.4,5
The present paper describes an investigation of the
influence of the plastic-deformation intensity and strain
applied at various stages of transformation on the steel
microstructure and mechanical properties. A significant
acceleration of the process is due to the steel heating at a
temperature just below transformation temperature Ac1
and the plastic strain.6
2 EXPERIMENTAL WORK
2.1 Material and thermomechanical treatment
The experimental work was performed using the
carbon steel C45 with the chemical composition listed in
Table 1. The initial microstructure consisted of ferrite
and lamellar pearlite with pronounced banding along the
bar axis (Figure 1). The hardness of the as-received
material was 180 HV, the 0.2 proof stress was 345 MPa,
the ultimate tensile strength was 629 MPa, the elongation
was A5 = 29 % and the impact toughness was KCV =
29 J/cm2.
Materiali in tehnologije / Materials and technology 48 (2014) 5, 797–800 797
UDK 621.77:669.111.3 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(5)797(2014)
The thermomechanical treatment was carried out in a
universal rolling mill that can be configured as either a
four-high rolling mill or a two-high mill. The two-high
configuration is used for hot rolling. The working roll
diameter is 550 mm. The maximum width of the rolled
plate is 400 mm and the thickness may range from 100
mm to 5 mm. The maximum rolling speed is 1.5 m/s.
During rolling the induction heating system, situated on
both sides of the rolling mill, can be used and a water-
spray facility for quenching is provided on the mill. The
rolling mill also includes hydraulic shears. The initial
dimensions of the specimens for thermomechanical treat-
ment were 330 mm × 50 mm × 30 mm.
Thermomechanical-treatment schedules (Table 2)
were proposed for investigating the impact of the strain
magnitude and the strains applied at various stages of the
pearlitic transformation on the microstructure and
mechanical properties. The main focus was the degree of
carbide spheroidisation and ferrite-grain refinement.
Austenitizing at 850 °C was followed by a thickness
reduction with the isothermal strain of  = 0.4. Before
the second and third deformations, the specimens were
air cooled. The second and third deformation steps were
applied at various stages of the austenite-to-pearlite
transformation (I, II, III – Figure 2). The strain magni-
tudes applied at these lower temperatures and at various
stages of the pearlitic transformation were 0.5 (one pass)
or 1 (two passes). When two passes were used, they
immediately followed each other and took place at a
virtually equal temperature. The temperature at point I
was approximately 675 °C. At point II, it was 685 °C
and at point III it was 675 °C. In all the cases, the
deformation speed was 1.5 m/s. After the last pass, the
specimens were either air cooled or quenched in water.
The quenching immediately followed the last pass to
allow a later determination of the austenite content.
3 RESULTS AND DISCUSSION
3.1 Metallographic observation
Specimens 1w, 2w, 3w were processed using the
schedules with a deformation at 850 °C applied at
various stages of the transformation (I, II, III) and
quenched in water. The metallographic examination
clearly revealed varying amounts of martensite in the
microstructure. In the course of the last deformation, the
expected austenite amounts in the specimens in sche-
dules 1w, 2w and 3w were approximately 50 %, 25 %
and 10 % max., respectively. These fractions correspond
D. HAUSEROVA et al.: ACCELERATED CARBIDE SPHEROIDISATION AND REFINEMENT (ASR) ...
798 Materiali in tehnologije / Materials and technology 48 (2014) 5, 797–800
Figure 3: SEM micrograph of the 4a specimen
Slika 3: SEM-posnetek mikrostrukture vzorca 4a
Table 2: List of schedules
Tabela 2: Seznam poteka preizkusov
Schedule
Temperature
of first
deformation
 =0.4
Temperature
of second
deformation
 = 0.5
Temperature
of third
deformation
 = 0.5
Cooling
1a 850 °C I - Air
1w 850 °C I - Water
2a 850 °C II - Air
2w 850 °C II - Water
3a 850 °C III - Air
3w 850 °C III - Water
4a 850 °C I I Air
4w 850 °C I I Water
5a 850 °C II II Air
5w 850 °C II II Water
6a 850 °C III III Air
6w 850 °C III III Water
Figure 1: SEM micrograph of the initial state
Slika 1: SEM-posnetek za~etne mikrostrukture
Table 1: Chemical composition of the C45 steel (mass fractions, w/%)
Tabela 1: Kemijska sestava jekla C45 (masni dele`i, w/%)
C Si Mn S P Cr Ni Cu Mo W
0.42 0.24 0.69 0.019 0.016 0.12 0.16 0.12 0.02 0.01
Figure 2: Transformation stages during deformation
Slika 2: Faze transformacije pri deformaciji
to the fractions of lamellar pearlite in the air-cooled
specimens (1a, 2a, 3a). The lamellar pearlite exhibits no
sign of spheroidisation or lamellae fragmentation and it
is assumed that the pearlite formed after the deformation.
However, in the pearlite already present in the micro-
structure during the last deformation, the lamellae were
transformed into elongated particles or, less frequently,
to globules and only a small fraction of the initial
pearlite lamellae spheroidised completely.
The 4a, 5a and 6a schedules comprised a deformation
at 850 °C, consisting of two deformation steps at various
stages of transformation, and the final air cooling. The
specimens contained two pearlite morphologies, as in the
1a, 2a and 3a specimens, with parts of lamellar pearlite
and, by regions, with globules and rod-like cementite
particles. The fraction of lamellar pearlite decreases from
the 4a to the 6a schedule (Figures 3 and 4), i.e., with the
progress of transformation of the austenite present
during the plastic deformation. The amount of austenite
was found by mapping the martensite fraction in water-
quenched specimens 4w, 5w and 6w, with the decreasing
martensite proportion in this order.
The mechanical deformation of austenite at transfor-
mation stages I or II (Figure 2) caused it to transform to
lamellar pearlite, as in the transformation of the austenite
unaffected by deformation. The deformation of lamellar
pearlite with the strain magnitude  = 0.5 at stages I, II
or III led to a fragmentation of the lamellae and to a for-
mation of predominantly elongated cementite particles.
Strain magnitude 1 (i.e., two passes with the strains of
0.5) caused a partial spheroidisation of pearlite lamellae
at all the stages of the pearlitic transformation, producing
cementite in the form of globules and rod-like particles.
3.2 Ferrite Grains
The characteristics of the ferrite grains in the micro-
structure depend strongly on the transformation stage (I,
II or III), at which deformation was applied (Figure 2).
The strain of 0.5 applied at stage I led to a 90 %
recrystallization of ferrite. The resulting grain size was
less than 8 μm (Figure 5). The strain of 0.5 applied at
stages II and III caused a recrystallization of only a small
fraction of ferrite grains. Approximately 80 % of the
ferrite grains exhibited a deformation substructure with
elongated grains and deformation-induced subgrains
(Figure 6). The size of the elongated ferrite grains was
approximately 20 μm.
In the case of the strain of magnitude 1, the differen-
ces between the ferrite grains after the schedules
involving the deformation at stages I and III were
smaller. The larger strain caused a recrystallization of
approximately 50 % of the ferrite grains even after the
deformation applied at stage III (Figures 4 and 7). Upon
schedules 4a, 5a and 6a, the size of the recrystallized
grains was 4 μm. Scarcely recrystallized grains with the
size of approximately 10 μm were also observed. With
an EBSD analysis the grain size and the volume fractions
of deformed and recrystallized grains could be assessed.
3.3 Mechanical properties
Tensile tests were performed on the flat specimens
with the dimensions of 35 mm × 6 mm × 4 mm. The
V-notch impact-toughness test-specimen size was 3 mm
× 4 mm with the notch depth of 1 mm. The HV10 hard-
ness was measured as well.
D. HAUSEROVA et al.: ACCELERATED CARBIDE SPHEROIDISATION AND REFINEMENT (ASR) ...
Materiali in tehnologije / Materials and technology 48 (2014) 5, 797–800 799
Figure 5: SEM micrograph of the 1a specimen
Slika 5: SEM-posnetek mikrostrukture vzorca 1a
Figure 4: SEM micrograph of the 6a specimen
Slika 4: SEM-posnetek mikrostrukture vzorca 6a
Figure 6: SEM micrograph of the 3a specimen
Slika 6: SEM-posnetek mikrostrukture vzorca 3a
Different proof stresses and ultimate strengths were
measured for three specimens upon these schedules
(Table 3). One of the three specimens showed approxi-
mately 50 MPa higher proof stress and ultimate tensile
strength than the others. The figures listed in Table 3 are
the average values of three tests.
Table 3: Mechanical properties
Tabela 3: Mehanske lastnosti
Schedule PS/MPa
UTS/
MPa
A5/
%
KCV/
(J cm–2) HV 10
Initial
condition 345 629 29 29 180
1a 489 702 26 38 204
2a 511 710 25 36 213
3a 542* 688* 21 35 201
4a 547 709 28 36 212
5a 550 708 27 35 210
6a 534* 668* 25 35 199
Different properties will be the subject of a further
investigation. All the thermomechanical treatment sche-
dules led to higher proof stress, ultimate tensile strength,
hardness and impact toughness (Table 3). The final
elongation was slightly lower than, or equal to, the initial
elongation. Upon the schedules with smaller strain mag-
nitudes (1a, 2a and 3a) the following trends were
observed for the specimens: If deformation was applied
at an early stage of the pearlitic transformation [I], the
resulting proof stress was the lowest of those measured.
When deformation was applied at a later stage of the
transformation [III], the resulting proof stress was
higher. In the specimens under the schedules with higher
strains (4a, 5a, 6a), this trend was not observed and their
strengths were virtually equal. The schedules with the
higher strains applied at the same transformation stage as
in the other schedules lead to a higher proof stress. This
observation does not apply to specimen 6a, as the
strength levels achieved are very similar (Table 3). In
terms of proof stress, less strain is sufficient ( = 0.5),
provided that it is applied at the final stage of the
pearlitic transformation.
4 CONCLUSION
The purpose of the investigation was to improve the
mechanical properties, promote the carbide
spheroidisation and refine the ferrite grains in the
medium-carbon C45 steel using controlled rolling. The
final deformation was applied in the intercritical range at
various stages of the transformation of austenite into
ferrite and carbide particles. In the specimens quenched
in water after the final deformation, no carbide
spheroidisation or fragmentation was observed. On the
contrary, in the specimens cooled in air spheroidised
carbides were observed. A higher proportion of
spheroidised particles was found in the specimens after a
larger final deformation with  = 1. The fraction of
spheroidised carbides increased with the applied strain
magnitude at the advanced stages of transformation. The
microstructure showed that after the deformation applied
at the final stages of transformation, recrystallisation also
took place in the ferrite grains. The proof stress and
ultimate tensile strength of the processed material were
higher by 200 MPa and almost 100 MPa, respectively,
than the corresponding characteristics of the feedstock.
The elongation levels were identical, whereas the
toughness of the final material was slightly higher than
that of the feedstock. The experimental rolling, thus,
improved the mechanical properties, facilitated a partial
spheroidisation of the carbides and required less time in
comparison with the conventional carbide-spheroidising
methods.
Acknowledgment
The results were achieved within the project
Thermo-chemical treatment of steels using fluidised bed
with thermoactive micropowders no. LF13032 co-funded
by the Ministry of Education, Youth and Sports of the
Czech Republic.
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800 Materiali in tehnologije / Materials and technology 48 (2014) 5, 797–800
Figure 7: SEM micrograph of the 6a specimen
Slika 7: SEM-posnetek mikrostrukture vzorca 6a