VSEBINA – CONTENTS
Predgovor urednika/Editor’s preface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 639
PREGLEDNI ^LANEK – REVIEW ARTICLE
Effect of the addition of niobium and aluminium on the microstructures and mechanical properties of micro-alloyed PM steels
Vpliv dodatka niobija in aluminija na mikrostrukturo in mehanske lastnosti mikrolegiranih PM jekel
S. Gündüz, M. A. Erden, H. Karabulut, M. Türkmen . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 641
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Characteristics of dye-sensitized solar cells with carbon nanomaterials
Zna~ilnosti na fiksirano barvo ob~utljivih solarnih celic z ogljikovimi nanomateriali
L. A. Dobrzañski, A. Mucha, M. Prokopiuk vel Prokopowicz, M. Szindler, A. Dryga³a, K. Lukaszkowicz . . . . . . . . . . . . . . . . . . . . . . . 649
The effect of the welding parameters and the coupling agent on the welding of composites
Vpliv parametrov varjenja in sredstva za spajanje na varjenje kompozitov
S. E. Erdogan, U. Huner . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 655
Chemical cross-linking of chitosan/polyvinyl alcohol electrospun nanofibers
Kemijsko zamre`enje elektro spredenih nanovlaken iz hitosan/polivinil alkohola
S. Pouranvari, F. Ebrahimi, G. Javadi, B. Maddah . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 663
Investigation of hole profiles in deep micro-hole drilling of AISI 420 stainless steel using powder-mixed dielectric fluids
Preiskava profilov luknje pri globokem vrtanju mikroluknje v AISI 420 nerjavnem jeklu s pomo~jo dielektri~ne teko~ine s prime{anim prahom
V. Yýlmaz . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 667
The phenomenon of reduced plasticity in low-alloyed copper
Pojav zmanj{anja plasti~nosti malo legiranega bakra
W. Ozgowicz, E. Kalinowska-Ozgowicz, B. Grzegorczyk, K. Lenik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 677
The effect of high-speed grinding technology on the properties of fly ash
Vpliv tehnologije hitrega mletja na lastnosti lete~ega pepela
K. Dvoøák, I. Hájková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 683
Investigation of the mechanical properties of electrochemically deposited Au-In alloy films using nano-indentation
Preiskava mehanskih lastnosti elektrokemijsko nane{enega filma zlitine Au-In z nanovtiskovanjem
S. Cherneva, R. Iankov, M. Georgiev, T. Dobrovolska, D. Stoychev . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 689
Growth of K2CO3-doped KDP crystal from an aqueous solution and an investigation of its physical properties
Rast KDP kristalov z dodatkom K2CO3 iz vodne raztopine in preiskava njihovih fizikalnih lastnosti
A. Rousta, H. R. Dizaji . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 695
Surface treatment of heat-treated cast magnesium and aluminium alloys
Obdelava povr{ine toplotno obdelanih magnezijevih in aluminijevih livnih zlitin
T. Tañski, M. Wiœniowski, W. Matysiak, M. Staszuk, R. Szklarek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 699
Analysis of the structural-defect influence on the magnetization process in and above the Rayleigh region
Analiza vpliva strukturnih defektov na proces magnetizacije v in nad Rayleigh podro~jem
K. Gruszka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 707
Effect of sulphide inclusions on the pitting-corrosion behaviour of high-Mn steels in chloride and alkaline solutions
Vpliv sulfidnih vklju~kov na jami~asto korozijo jekel z visoko vsebnostjo Mn v raztopinah kloridov in alkalij
A. Grajcar, A. P³achciñska . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 713
Influence of Na2SiF6 on the surface morphology and corrosion resistance of an AM60 magnesium alloy coated by micro arc
oxidation
Vpliv Na2SiF6 na morfologijo povr{ine in korozijsko odpornost magnezijeve zlitine AM60, prekrite z mikrooblo~no oksidacijo
A. Ayday . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 719
Mechanical properties of polyamide/carbon-fiber-fabric composites
Mehanske lastnosti kompozitne tkanine iz poliamid/ogljikovih vlaken
C.-E. Pelin, G. Pelin, A. ªtefan, E. Andronescu, I. Dincã, A. Ficai, R. Truºcã . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 723
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 50(5)637–834(2016)
MATER.
TEHNOL.
LETNIK
VOLUME 50
[TEV.
NO. 5
STR.
P. 637–834
LJUBLJANA
SLOVENIJA
SEP.–OKT.
SEP.–OCT. 2016
Evaluation of the grindability of recycled glass in the production of blended cements
Ocena sposobnosti drobljenja recikliranega stekla pri proizvodnji me{anih cementov
K. Dvoøák, D. Dolák, D. V{ianský, P. Dobrovolný . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 729
Rheological properties of alumina ceramic slurries for ceramic shell-mould fabrication
Reolo{ke lastnosti go{~e iz glinice za izdelavo kerami~nih kalupov
J. Szymañska, P. Wiœniewski, M. Ma³ek, J. Mizera . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 735
Effect of mechanical activation on the synthesis of a magnesium aluminate spinel
Vpliv mehanske aktivacije na sintezo magnezij-aluminatnega {pinela
D. Kýrsever, N. K. Karabulut, N. Canikoðlu, H. Ö. Toplan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 739
Phase and microstructure development of LSCM perovskite materials for SOFC anodes prepared by the
carbonate-coprecipitation method
Razvoj kristalnih faz in mikrostrukture LSCM perovskitnih materialov za SOFC anode, pripravljenih s karbonatno metodo koprecipitacije
K. Zupan, M. Marin{ek, T. Skalar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 743
Artificial aggregate from sintered coal ash
Umetni agregat iz sintranega pepela premoga
V. Cerny, R. Drochytka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 749
Investigation studies involving wear-resistant ALD/PVD hybrid coatings on sintered tool substrates
Preiskave obrabne odpornosti hibridnega nanosa ALD/PVD na sintranem orodju
M. Staszuk, D. Paku³a, T. Tañski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 755
Dissimilar spot welding of DQSK/DP600 steels: the weld-nugget growth
To~kasto varjenje jekel DQSK/DP600: rast jedra zvara
S. P. Hoveida Marashi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 761
Armour plates from Kozlov rob – analyses of two unusual finds
Oklepni plo{~i s Kozlovega roba – analize dveh nenavadnih najdb
T. Lazar, P. Mrvar, M. Lamut, P. Fajfar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 767
Numerical and experimental investigation of the effect of hydrostatic pressure on the residual stress in boiler-tube welds
Numeri~na in eksperimentalna preiskava vpliva hidrostati~nega tlaka na zaostale napetosti v zvaru na kotlovski cevi
D. Danyali, E. Ranjbarnodeh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 775
Effect of direct cooling conditions on the microstructure and properties of hot-forged HSLA steels for mining applications
Vpliv pogojev ohlajanja na mikrostrukturo in lastnosti vro~e kovanih HSLA jekel za uporabo v rudarstvu
P. Skubisz, £. Lisiecki, T. Skowronek, A. ¯ak, W. Zalecki . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 783
Influence of the tool rotational speed on the microstructure and joint strength of friction-stir spot-welded pure copper
Vpliv hitrosti vrtenja orodja na mikrostrukturo in trdnost torno vrtilno to~kasto zvarjenega spoja ~istega bakra
I. Dinaharan, E. T. Akinlabi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 791
Measurement of bio-impedance on an isolated rat sciatic nerve obtained with specific current stimulating pulses
Meritev bioimpedance na izoliranem `ivcu Ischiadicus pri podgani, vzbujenem s posebnimi tokovnimi stimulacijskimi impulzi
J. Rozman, M. C. @u`ek, R. Frange`, S. Ribari~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 797
Influence of different production processes on the biodegradability of an FeMn17 alloy
Vpliv razli~nih procesov izdelave na biorazgradljivost zlitine FeMn17
A. Kocijan, I. Paulin, ^. Donik, M. Ho~evar, K. Zeli~, M. Godec. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 805
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Effect of a combination of fly ash and shrinkage-reducing additives on the properties of alkali-activated slag-based mortars
Vpliv kombinacije lete~ega pepela in dodatka za zmanj{anje kr~enja na lastnosti malte iz z alkalijami aktivirane `lindre
V. Bílek, L. Kalina, J. Koplík, M. Mon~eková, R. Novotný . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 813
Cutting-tool performance in the end milling of carbon-fiber-reinforced plastics
Zmogljivost rezilnega orodja pri rezkanju plastike, oja~ane z ogljikovimi vlakni
O. Bílek, S. Rusnáková, M. @aludek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 819
Influence of solidification speed on the structure and magnetic properties of Nd10Fe81Zr1B6 in the as-cast state
Vpliv hitrosti strjevanja na strukturo in magnetne lastnosti zlitine Nd10Fe81Zr1B6 v litem stanju
M. Doœpia³, M. Nabia³ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 823
Metalografska preiskava in korozijska odpornost zvarov feritnega nerjavnega jekla
Metallographic investigation and corrosion resistance of welds of ferritic stainless steels
M. Torkar, A. Kocijan, R. Celin, J. Burja, B. Podgornik. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 829
EDITOR’S PREFACE
After 4 years of great service in the
job as Editor-in-Chief of Materials &
Technology, Matja` Torkar has taken
retirement. He saw the journal
through some difficult times and
established it as an important inter-
national record of scientific
achievements. He will be a tough act
to follow.
The next issue, 6/2016, will see
some changes. After many years of
being both open access and free to
publish, we will have to
acknowledge the economic
realities of the situation and
start to charge for each
published paper. In line with
other Slovenian journals we will
introduce a publication fee of 300, payable on accept-
ance of the manuscript, with a reduced fee of 150 for
those people presenting at the annual International
Conference on Materials & Technology, held annually in
Portoro`, Slovenia. This fee will be applied to
manuscripts submitted after 30 Sept. 2016.
A second change, which we feel is in line with the
progress the journal has been making over the past de-
cade, is to publish original and review scientific articles.
From now on we will not have a separate section for pro-
fessional articles. As the number of submissions in the
past years has soared, from 120 in 1995 to near 400 in
2015, our selection criteria have had to become stricter.
So, in order to help authors get their articles published
we have re-written the Instructions for Authors, intro-
duced a template for manuscript submission and now
require a mandatory check-list to accompany all new
manuscripts. For the time being we will still be accepting
submissions by e-mail to mit@imt.si, where they will be
handled first by the journal’s new Technical Editor, Erika
Nared, who will be checking each manuscript’s eligi-
bility, before they are assessed by myself or members of
the Editorial Board. However, we expect to be switching
to a web-based submission procedure in next few years.
The journal is becoming increasingly international,
with manuscripts being submitted from all over the
world. It is very gratifying to see that Materials & Tech-
nology is now such a familiar and respected journal in
places outside of Slovenia, but we would still like to
receive more submissions from researchers from Slo-
venia.
PREDGOVOR UREDNIKA
Po 4 letih delovanja kot glavni in odgovorni
urednik revije je dr. Matja` Torkar zaradi
upokojitve prenehal z urednikovanjem.
Skozi izzivov polno obdobje je vodil
revijo, da je postala to, kar predstavlja
danes: pomembno mednarodno uveljav-
ljeno revijo, ki objavlja mnoge znan-
stvene dose`ke.
Z naslednjo {tevilko (6/2016), ki bo
iz{la v decembru 2016, bomo uvedli
nekaj sprememb. Po mnogih letih,
tako odprtega dostopa revije kot
brezpla~ne objave ~lankov, bo-
mo zaradi ekonomske situacije
za~eli zara~unavati objavo ~lan-
kov v na{i reviji. Tako kot to
po~ne `e prenekatera slovenska
znanstvena revija. Za obi~ajne
~lanke bo cena 300 , za tiste ~lanke, ki bodo predstav-
ljeni na letni Mednarodni konferenci o materialih in teh-
nologijah, ki poteka vsako leto v Portoro`u, pa 150 . Za
~lanke, ki jih bomo prejeli od 30. septembra 2016 dalje,
bo potrebno pla~ilo za objavo.
Druga sprememba, za katero menimo, da je dobro-
do{la in nujna za nadaljnji razvoj revije, je objavljanje
izvirnih znanstvenih in preglednih ~lankov. Od sedaj
naprej tako ne bo ve~ sekcije za strokovne ~lanke. Ker je
v zadnjem ~asu {tevilo oddaje ~lankov izredno naraslo,
od nekaj 120 v letu 1995 na blizu 400 v letu 2015, smo
mnenja, da je potrebno kriterije za izbiro ~lankov za-
ostriti. Tako smo, v pomo~ avtorjem, posodobili Navo-
dila za avtorje, predstavili bomo vzorec oz. predlogo,
kako naj bo ~lanek napisan in poleg oddaje ~lanka bomo
zahtevali obvezno oddajo kontrolnega seznama, ki ga bo
moral izpolniti avtor. Oddaja ~lankov bo tako {e vedno
potekala preko e-po{te: mit@imt.si, kjer jih bo najprej
obravnavala nova tehni~na urednica, Erika Nared, preden
bodo posredovani v pregled meni ali ~lanom uredni{kega
odbora. Nadalje, v bodo~e na~rtujemo prehod na spletno
oddajanje ~lankov za objavo v reviji.
Revija postaja vse bolj mednarodna, saj so ~lanki
poslani iz razli~nih delov sveta. V ~ast nam je, da je
revija Materiali in tehnologije danes tako znana in
ugledna tudi izven Slovenije, vendar si `elimo, da bi v
na{i reviji objavljalo ve~je {tevilo strokovnjakov in
raziskovalcev iz Slovenije.
Na{e osnovno prizadevanje ostaja {e naprej: da je
revija Materiali in tehnologije znanstvena revija, ki je
posve~ena izvirnim znanstvenim ~lankom in ~lankom, ki
Our basic remit remains the same: we are a scientific
journal, devoted to original scientific papers and articles
relating to fundamental and applied science and techno-
logy. Topics of particular interest to the journal include
metallic materials, inorganic materials, polymers, va-
cuum techniques and nanomaterials. In response to
requests by authors we will be making a lot of effort over
the next year to reduce the waiting period for publication
to 4–6 months, so as to provide a better service to both
authors and readers. In addition, we welcome construc-
tive comments from anyone with an interest in the
journal and are happy to discuss new ideas.
I would like to thank you for your interest in Mate-
rials & Technology journal and look forward to working
with you in the future.
Paul McGuiness,
Editor-in-Chief
obravnavajo tematiko temeljne in aplikativne znanosti ter
tehnologije. Tematike, ki so posebnega pomena za revijo
vklju~ujejo kovinske materiale, anorganske materiale,
polimere, vakuumsko tehniko in nanomateriale. Glede na
`elje in zahteve mnogih avtorjev, se bomo v naslednjem
letu trudili skraj{ati obdobje ~akanja na objavo ~lanka
(na 4–6 mesecev) in tako stopiti naproti tako avtorjem,
kot tudi bralcem. Veseli bomo vsake va{e konstruktivne
kritike, predlogov ali novih idej, ki nam bodo pomagale
narediti revijo {e bolj{o.
Zahvaljujem se vam za va{e zanimanje in prebiranje
revije Materiali in tehnologije ter se veselim prihodnjega
sodelovanja z vami.
Paul McGuiness,
Glavni in odgovorni urednik
S. GÜNDÜZ et al.: EFFECT OF THE ADDITION OF NIOBIUM AND ALUMINIUM ON THE MICROSTRUCTURES ...
641–648
EFFECT OF THE ADDITION OF NIOBIUM AND ALUMINIUM ON
THE MICROSTRUCTURES AND MECHANICAL PROPERTIES OF
MICRO-ALLOYED PM STEELS
VPLIV DODATKA NIOBIJA IN ALUMINIJA NA
MIKROSTRUKTURO IN MEHANSKE LASTNOSTI
MIKROLEGIRANIH PM JEKEL
Süleyman Gündüz1, Mehmet Akif Erden2, Hasan Karabulut3, Mustafa Türkmen4
1Karabük University, Faculty of Technology, Department of Manufacturing Engineering, 78050 Karabük, Turkey
2Karabük University, Institute of Science and Technology, Department of Manufacturing Engineering, 78050 Karabük, Turkey
3Karabük University, Karabük Vocational School, 78050 Karabük, Turkey
4Kocaeli University, Hereke Vocational School, Department of Metallurgy, Kocaeli, Turkey
hasankarabulut@karabuk.edu.tr
Prejem rokopisa – received: 2015-08-08; sprejem za objavo – accepted for publication: 2015-09-04
doi:10.17222/mit.2015.248
In this work, the effects of the addition of Nb and Al on the microstructures and tensile behaviours of micro-alloyed powder
metallurgy (PM) steels were investigated. The microstructure of the micro-alloyed PM steels was examined by light microscope,
SEM, XRD, XRF and EDS. The results indicated that the addition of (0.1, 0.15 or 0.2) % of Nb-Al increases the yield strength
(YS) and the ultimate tensile strength (UTS) of the PM sintered steels. Elongation also tends to improve with an increasing Nb
and Al content. In addition, the Nb and Al limit the grain growth during austenitization.
Keywords: powder metallurgy, micro-alloyed steels, microstructure
V delu je bil preiskovan vpliv dodatka Nb in Al na mikrostrukturo in na natezno trdnost mikrolegiranih PM jekel, izdelanih iz
prahov. Mikrostruktura mikrolegiranih PM jekel je bila preiskovana s svetlobnim mikroskopom ter s SEM, XRD, XRF in EDS.
Rezultati so pokazali, da dodatek (0,1, 0,15 ali 0,2) % Nb-Al pove~a mejo plasti~nosti (YS) in natezno trdnost (UTS) sintranih
PM jekel. Tudi raztezek se izbolj{a s pove~ano vsebnostjo Nb in Al. Ugotovljeno je {e, da dodatek Nb in Al omejuje rast zrn
med avstenitizacijo.
Klju~ne besede: metalurgija prahov, mikrolegirana jekla, mikrostruktura
1 INTRODUCTION
Steels with a minimal strength high toughness and
excellent weldability are required in a wide range of
applications. This combination of properties is achieved
by optimizing the chemical composition and by thermo-
mechanical processing (TMP). The addition of micro-
alloying elements such as Nb, V, Ti and Al contribute to
an increase in strength, both directly, through micro-
structural refinement and precipitation strengthening,
and indirectly, through enhanced hardenability and an
associated modification of the transformation micro-
structure.1–3
Niobium forms nitrides and carbides, but it is the
carbide that is the most important. In steels it precipitates
at a temperature just below 1000 °C and prevents auste-
nite recrystallization. Niobium carbide particles are very
effective in preventing austenite recrystallization and the
formation of "pancake" grains that transform to fine
ferrite grains.4–6 Thus, the precipitation of niobium car-
bide particles plays a major role in controlling the final
microstructure and the product properties. Besides in-
creasing the non-recrystallization temperature, the pre-
sence of precipitates also increases the austenite grain-
coarsening temperature7, which is important through
controlled reheating.8 Niobium is widely used in this
way for the production of fined-grained pipeline and
other structural steels.
A large volume of work was spent investigating the
effect of niobium on the recrystallization and growth of
austenite grains,5,6,9,10 and it was found that a minor
addition of niobium to the steel was sufficient to inhibit
the static recrystallization of austenite and to achieve the
final microstructure.11 Similar effects were observed in
titanium, and vanadium steels were found generally less
marked.12
Aluminium only forms nitride precipitates, which are
stable at temperatures above 900 °C. It forms during
reheating to heat-treatment temperatures and at the ex-
pense of the vanadium nitride, if present. At normalising
temperatures it is stable, pins grain boundaries and is
effective in refining the grain size.12 Aluminium nitride
forms only with a hexagonal crystal lattice and it was not
found in any substantial solid solution in the face-cen-
tred-cubic micro-alloy carbonitrides.4
AlN precipitation occurs at both grain boundaries and
dislocations. It has been shown that the precipitation
kinetics depends on the content of nitrogen and alumi-
Materiali in tehnologije / Materials and technology 50 (2016) 5, 641–648 641
UDK 621.762:669.293:669.71:67.017 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 50(5)641(2016)
nium in the steel and also on the grain size and the
annealing temperature.13 In the austenite region, the pre-
cipitation occurs predominantly at the grain boundaries
because of the considerable volumetric misfit of the AlN
precipitates and the steel matrix,14 and the increased
diffusion rate of both elements at the grain boundaries,
as compared to the grains.15
The rising costs and disposition volatility of metals
has led to the development of new PM steels. In
particular, PM steels with the addition of copper, nickel
and molybdenum to compete with wrought grades.
Recently, due to cost constraints and availability, PM
steels have also included chromium and manganese as
the alloying elements. These material systems are cate-
gorized as alloy steels, since significant levels of these
elements are required to achieve changes in the mecha-
nical properties.16 The demand for cheaper but high-
strength structural steels forced powder metallurgists to
seek newer compositions with wider applications, and
such steels found applications in the automotive industry,
aerospace, and power tools.17
Micro-alloyed steels contain small amounts of nio-
bium, vanadium or titanium, generally at levels between
0.02 and 0.2 % of mass fractions. Despite the low alloy
content, micro-alloying can lead to a major increase in
the strength and toughness as a result of carbonitride par-
ticles, which led to precipitation strengthening and grain
refinement.16 The production of micro-alloyed steels is
estimated to be around 12 % of the total world steel pro-
duction and are used in every major steel market sector.
In many parts of the world their development has played
an important role in the expansion of some key indu-
stries, such as oil and gas extraction, construction, and
transportation.18 The present study was undertaken to
examine the effect of Nb and Al on the microstructure
and mechanical properties of sintered PM steels. The
mechanical properties were determined and the micro-
structures were investigated in the sintered condition to
assess the role of precipitation strengthening and grain
refinement.
2 MATERIALS AND EXPERIMENTAL
PROCEDURE
In this investigation, Fe, Nb and Al powders of 180
μm, <45 μm, and <75 μm supplied by Aldrich were used,
with purities of 99.9 %, 99.8 % and 93 %, respectively.
Electron micrographs of the powders presented in Fig-
ure 1 reveal an irregular shape of the powder particles
for all the studied samples. The required mass of
Fe-0.25C (Alloy 1), Fe-0.25C-0.05Nb-0.05Al (Alloy 2),
Fe-0.25C-0.075Nb-0.075Al (Alloy 3) and Fe-0.25C-
0.1Nb-0.1Al (Alloy 4) powders was accurately weighed
and mixed in an industrial conic mixer for 1 h. A total of
0.45 of graphite was added to reach a carbon content of
0.25 % in the sintered test pieces. Zn-stearate was added
as a lubricant. The mixed powder mass was then com-
pacted into dog-bone tensile specimens using a hydraulic
press with a capacity of 100 tons and a compaction
pressure of 700 MPa. Standard cross-section tensile-test
specimens were produced in accordance with the stan-
dard of ASTM E8/E8M 19 as shown in Figure 2. The
specimen has two large shoulders with a smaller
cross-section gauge in between. The shoulders are large
so they can be readily gripped, whereas the gauge sec-
tion has a smaller cross-section.
The test pieces were sintered in a tube furnace with
an argon atmosphere. The sintering cycle was: heating to
1350 °C at a rate of 5 °C/min, holding at this tempera-
ture for 1 h, cooling to room temperature at a rate of
5 °C/min. The sintered density was determined by
Archimedes’ principle using pure water according to
ASTM B 328-96.20 Four measurements were made for
each composition and the variation of those values was
less than 1 %. A tensile test at room temperature was
performed using a Schimadzu tensile-testing machine at
a crosshead speed of 1 mm min–1 and perfect alignment
of the specimens in the pull direction and no slippage.
S. GÜNDÜZ et al.: EFFECT OF THE ADDITION OF NIOBIUM AND ALUMINIUM ON THE MICROSTRUCTURES ...
642 Materiali in tehnologije / Materials and technology 50 (2016) 5, 641–648
Figure 1: SEM micrographs of Fe, Nb and Al powders: a) Fe 180 μm, b) Nb <45 μm and c) Al <75 μm
Slika 1: SEM-posnetki prahov Fe, Nb in Al: a) Fe 180 μm, b) Nb <45 μm in c) Al <75 μm
Figure 2: General view of tensile test specimen sintered at 1350 °C
for 1 h
Slika 2: Izgled preizku{anca za natezni preizkus, 1 h sintranega na
1350 °C
Three specimens of each alloy were tested and the mean
value was used.
The examination of the microstructure was carried
out using optical and scanning electron microscopy
(SEM), and energy-dispersive spectrometry (EDS) was
used to provide elemental analyses of the particles. The
average elemental composition of alloys was determined
with the X-ray fluorescence technique (XRF). The che-
mical composition of the produced PM steels is pre-
sented in Table 1. The specimens were mechanically
polished using a standard metallographic procedure and
etched with 2 % Nital solution. The microstructures were
examined in a Nikon ECLIPSE L150 microscope with a
magnification of 50× to 1000×. An electrochemical
extraction technique was used to characterise the precipi-
tates in the examined specimens. This procedure in-
volves the electrochemical dissolution of tensile speci-
men in the electrolyte (10 % HCl in methanol) and
filtering to separate the precipitate from the solution. The
dissolution time of 8 hours sufficed for the dissolution of
around two grams of the specimen. For the filtering,
cellulose acetate filters were used with a pore size of 0.4
μm and X-ray diffraction (XRD) analysis of the collected
residues with precipitates was carried out. The XRD data
were obtained using Cu-K radiation, a scan step of 4°, a
step time of 1 min and 2 range from 20° to 90°.
Table 1: Chemical composition of PM steel and microalloyed PM
steels
Tabela 1: Kemijska sestava PM jekla in mikrolegiranih PM jekel
Fe C Mn Al Nb
Alloy 1 99.232 0.249 0.2110 0.0000 0.0000
Alloy 2 98,5174 0.2355 0.2316 0.0514 0.0468
Alloy 3 99.0191 0.2624 0.2223 0.0752 0.0740
Alloy 4 99.0452 0.2653 0.2223 0.0974 0.0918
The grain size measurement was carried out using the
mean linear intercept (mli) method, with the line inclined
by 45°. At least 500 grains cut by the intersecting line
were counted for each sample. From the results the mean
linear intercept grain sizes21,22 was determined. The sta-
tistical errors for the assessment of the mean linear
intercept were examined by Woodhead and reviewed by
J. R. Blank and T. Gladman,23 and it was suggested that
the relative error of the individual intercept value (i/i) =
0.7 is almost constant for a variety of regular and space-
filling polyhedrals. The relative error, , of the mean
linear intercept based on the measurement of n grains
(SEi/i) is deduced as shown in Equation (1):
f
SE
i n n
i
i
= = =

 0 7.
(1)
where i is the standard deviation of the assessment of
the intercept lengths.
The volume fraction of ferrite or pearlite was also
calculated using the systematic point-count method.24,25
According to this method, when the grid points intersect
the ferrite boundary, they are counted as half. Errors in
the point counting were also calculated by T. Gladman
and J. Woodhead26 in Equation (2):
 =
−f f
N
( )1
(2)
where  is the standard deviation, f is the measured
volume fraction of ferrite or pearlite and N is the total
number of points counted. All the volume fractions were
expressed ±  (standard deviation)
3 RESULTS AND DISCUSSION
The light micrographs of the Fe-0.25C (Alloy 1),
Fe-0.25C-0.05Nb-0.05Al (Alloy 2), Fe-0.25C-0.075Nb-
0.075Al (Alloy 3) and Fe-0.25C-0.1Nb-0.1Al (Alloy 4)
PM steel in Figure 3 show that the microstructures of the
examined steels consist of ferrite and pearlite grains of
varying sizes. In Table 2 the relative density, phase
volume fractions and mean linear intercept grain sizes of
the specimens are listed.
Figure 3 and the data in Table 2 indicate the grain
sizes decreasing with an increasing content of Nb-Al,
from 0.1, 0.15 or 0.2 %. A major benefit of micro-alloy-
ing is the decrease of the grain growth rate during
S. GÜNDÜZ et al.: EFFECT OF THE ADDITION OF NIOBIUM AND ALUMINIUM ON THE MICROSTRUCTURES ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 641–648 643
Figure 3: Microstructures of sintered PM steel and microalloyed PM
steels: a) Alloy 1, b) Alloy 2, c) Alloy 3, and d) Alloy 4
Slika 3: Mikrostrukture sintranega PM jekla in mikrolegiranih PM
jekel: a) zlitina 1, b) zlitina 2, c) zlitina 3 in d) zlitina 4
Table 2: Relative density, mean linear intercept grain sizes and
volume fractions of ferrite and pearlite phases in the PM steel and
microalloyed PM steels
Tabela 2: Relativna gostota, srednja velikost zrn pri linearni
intercepciji in volumenski dele` ferita in perlita v PM jeklu in v PM
mikrolegiranem jeklu
Alloy Relativedensity (%)
Grain size
(μm)
Ferrite
(%)
Pearlite
(%)
Alloy 1 92 29.7±0.93 78.4±0.018 21.6±0.018
Alloy 2 94 27.2±0.85 75±0.019 25±0.019
Alloy 3 94 23.4±0.73 74±0.02 26±0.019
Alloy 4 94 22.9±0.71 73±0.02 27±0.02
austenitizing, and if fine precipitates exist during austen-
tizing the growth of grains is restricted and a finer grain
size is obtained after quenching.16,27 Niobium has a low
solubility product that allows the substantial dissolution
of niobium carbonitrides only at elevated temperatures.
At low temperatures in the austenite range, the carbo-
nitride shows such low solubility and the dispersion
strengthening is not generally observed. The undissolved
carbonitride at these temperatures acts mostly as an
effective grain refiner. The marked change in carboni-
tride dissolution between the high and low temperatures
(1300 °C and 900 °C) in the austenite temperature range
allows substantial strain-accelerated precipitation at
temperatures below about 1000 °C, and produces what is
arguably the most obvious distinctive effect of niobium,
i.e., the marked retardation of recrystallization at these
temperatures.4,28–30
Aluminium only forms a nitride that is stable at low
temperatures in the austenite range, but will dissolve
progressively as the temperature is increased. The extent
of the dissolution at high temperatures in the austenite
range (e.g. 1350 °C) depends upon the content of alumi-
nium and nitrogen, but the solubility product for alumi-
nium nitride is low and the only common micro-alloy
nitride having a lower solubility product is titanium
nitride. Aluminium nitride has the distinction of forming
a separate nitride, which has a hexagonal crystal struc-
ture and forms no solid solution with the face-centred-
cubic micro-alloy carbonitrides.12
Table 3 shows the yield strength (YS), ultimate
tensile strength (UTS) and elongation of the examined
PM steels, and Figure 4 shows typical examples of the
stress-strain curves obtained with the tensile test and a
general increase of YS and UTS of steel with the addition
of Nb and Al. Elongation tends to improve with increas-
ing Nb-Al content. These changes are a consequence of
the differences in the precipitation distribution.31 High
strength and good toughness in micro-alloyed steels are
achieved by a combination of micro-alloying and con-
trolled rolling. During sintering and slow cooling from
the sintering temperature NbC(N) or AlN precipitates
form in the austenite and the ferrite during the austenite-
ferrite transformation or after transformation, as
suggested by M. A. Erden at al.32 These lead to an in-
crease in the strength compared with the titanium-free
alloy.
Table 3: Mechanical properties of sintered PM steel and microalloyed
PM steels
Tabela 3: Mehanske lastnosti sintranega PM jekla in mikrolegiranih
PM jekel
Alloy Yield strength(MPa)
Ultimate tensile
strength (MPa)
Elongation
(%)
Alloy 1 144 252 13
Alloy 2 198 356 12
Alloy 3 209 375 12
Alloy 4 220 394 13
The alloying elements have widely differing effects
due to the different solubilities of their carbides and
nitrides in both austenite and ferrite as well as their
different precipitation kinetics. The strength is increased
by grain refinement and precipitation hardening, by a
sufficient content of carbon and nitrogen in the steel.33,34
In modern micro-alloyed steels, the requirements for
specific properties may call for the use of more than one
micro-alloying element. An example are Al-Nb steels,
where aluminium is used for the grain refinement and
niobium for controlling the hot-rolled microstructure and
dispersion strengthening. The behaviour of the micro-
alloying elements can be modified by the presence of
another of them and changes dependent on the particular
elements. In principle, it depends on their mutual insolu-
bility or mutual solubility. AlN has a close-packed-hexa-
gonal structure, with little or no solubility for niobium
and the NbN has a cubic structure with little or no
solubility for aluminium. Under these conditions, the
two separate nitrides can co-exist in the austenite accord-
ing to their own solubility products.35
In the present experimental work, the solubility pro-
duct of AlN, NbN and NbC at 1350 °C was calculated
using the equations given by K. Narita.36 At 1350 °C the
solubility products of AlN, NbN and NbC are 2.3×10–3,
3.7×10–3 and 3.6×10–2. It is clear that the solubility of
NbC is higher than NbN and AlN, and therefore niobium
and carbon atoms should be present in the solid solution
during sintering at 1350 °C. The dissolved Nb will
precipitate as NbC(N) in austenite or ferrite, depending
on the cooling rate. The strength increment in micro-
alloyed PM steels is mainly due to precipitation harden-
ing, resulting in the formation of AlN and NbC(N).
The steel’s mechanical properties and toughness were
analysed in terms of the influence of grain size. Tradi-
tionally, the approaches developed by E. O. Hall37, based
on experimental observations, and by N. S. Petch,38
based on both experimental and theoretical approaches.
S. GÜNDÜZ et al.: EFFECT OF THE ADDITION OF NIOBIUM AND ALUMINIUM ON THE MICROSTRUCTURES ...
644 Materiali in tehnologije / Materials and technology 50 (2016) 5, 641–648
Figure 4: Variation of stress–strain curves of the PM steel and
microalloyed PM steels at different percentages of Nb and Al content:
a) Alloy 1, b) Alloy 2, c) Alloy 3, and d) Alloy 4
Slika 4: Spreminjanje krivulj napetost-raztezek PM jekla in
mikrolegiranih PM jekel pri razli~nih vsebnostih Nb in Al: a) zlitina 1,
b) zlitina 2, c) zlitina 3 in d) zlitina 4
The relation between the yield strength and the grain size
is now commonly known as the Hall-Petch equation (3):
 y yk d= +
−
0
1 2/ (3)
where y is the lower yield stress, o is the friction
stress, ky is the strengthening coefficient and d is the
grain size. To study the influence of precipitates and
clusters on the strength of micro-alloyed PM steel it is
necessary to calculate the value of p, which represents
the strength obtained from precipitates and clusters in
the micro-alloyed PM steel containing a different weight
percentage of Nb-Al. This was done using the F. B.
Pickering and T. Gladman39 equation (4):
 y pd= + +
−54 17 4 1 2. / (4)
where d is the grain diameter in mm, and p is the
strength obtained from precipitates and clusters. In the
present study the p values were calculated using
equation (2) for micro-alloyed PM steels. A value for
the level of p was derived by subtracting the predicted
yield strength from the actual yield strength. The values
of p are given in Table 4 and vary from -10 to 51 MPa
for the Nb-Al-free and Nb-Al-added micro-alloyed PM
steels tensile tested at room temperature.
As seen in Table 4, different values for p were ob-
served in the PM steel and microalloyed PM steels. For
example, the additions (0.1, 0.15 and 0.2) of Nb-Al
increased the precipitation contribution (p). This is a
result of the formation of carbonitrides after sintering at
1350 °C, which led to both precipitation strengthening
and grain refinement. Alloy 1 without Nb-Al does not
show any measurable precipitation strengthening. The
role of niobium in sintered PM steels was investigated by
several investigators.40–43 To examine this effect a simple
iron-carbon system was investigated by 0.16 % of mass
fractions of Nb and carbon contents from 0.10 to 0.50 %
of mass fractions and transverse rupture strength and
apparent hardness were measured in the sintered condi-
tion. At low carbon there is an increase in the transverse
rupture strength, presumably due to the precipitation of
NbC(N). The precipitation of nitrides, as AlN may also
occur during sintering and/or cooling after sintering,
since nitrogen cannot be completely avoided in PM
steels using current sintering technologies, such as
sintering under an argon atmosphere. It has been shown
that the  	  +  transformation in isothermal condi-
tions accelerates the precipitation kinetics of AlN, due to
the lower solubility of nitrogen in ferrite.44
Table 4: Structure-property analyses of PM steel and microalloyed
PM steels tensile tested at room temperature
Tabela 4: Analiza lastnosti PM jekla in mikrolegiranih PM jekel pri
nateznem preizkusu na sobni temperaturi
Alloy o(MPa) d (μm)
kyd–1/2
(MPa)
ytotal
(MPa)
y test
(MPa)
p
(MPa)
Alloy 1 54 29.7±0.93 100 154 144 –10
Alloy 2 54 27.2±0.85 106 160 198 38
Alloy 3 54 23.4±0.73 113 167 209 42
Alloy 4 54 22.9±0.71 115 169 220 51
o: the friction stress, d: grain size, ky: strengthening coefficient,
y total: predicted yield strength, y test: actual yield strength,
p: difference between predicted yield strength and actual yield
strength (strength obtained from precipitates)
The density is also expected to significantly affect the
properties of PM steel and micro-alloyed PM steels
because the pores are potential crack-initiation sites, and
can also guide and propagate cracks through the mater-
ial. This reduces the strength as well as the heat-transfer
and cooling rates after sintering.45 A strong microstruc-
ture may be obtained by incorporating small amounts of
alloying elements to compensate for the effect of pores.46
S. GÜNDÜZ et al.: EFFECT OF THE ADDITION OF NIOBIUM AND ALUMINIUM ON THE MICROSTRUCTURES ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 641–648 645
Figure 6: SEM micrograph for Alloy 4 and corresponding EDS of the
indicated points
Slika 6: SEM-posnetek zlitine 4 in EDS-analiza ozna~enih to~k
Figure 5: SEM micrograph for Alloy 2 and EDS line scan of the
indicated particle
Slika 5: SEM-posnetek zlitine 2 in EDS-analiza preko ozna~enega
delca
In the present work PM steel and micro-alloyed PM
steels showed a similar relative density, i.e., 92 % and
94 %, for the as-sintered condition. This may explain
that the strength increase is the effect of the different
content of NbC(N), AlN in micro-alloyed PM steel.
Figure 5 illustrates EDS Fe, Nb and C line analysis
of the cross-section of the matrix and precipitate particle
in Alloy 2 (Fe-0.25C-0.05Nb-0.05Al). Two distinct con-
stituents are detectable: a Fe-rich matrix and a Nb-rich
phase depicted a sharp increase from the Fe-rich matrix
to the particle and inverse Fe behaviour. Figure 6 also
shows EDS analyses at points of 1 (precipitate particles)
and 2 (matrix) marked in the micrograph of Alloy 4
(Fe-0.25C-0.1Nb-0.1Al). As seen in Figure 6, at point 1
the Fe, C, N, Al, Nb and at point 3 the Fe, C, N contents
are detected. Concerning the literature and the current
results, some NbC(N) and AlN particles probably did not
dissolve during sintering. The results of the EDS anal-
ysis agree with the precipitates visible on the SEM
micrograph of the fracture of micro-alloyed PM steels
(Alloys 2 and 4).
Past studies to characterize precipitation in micro-
alloyed steels have been mainly carried out using the
transmission electron microscopy (TEM) of extraction
replicas and thin foils.47–50 However, quantitative chemi-
cal analysis is also helpful to determine the amounts of
micro-alloying elements in solid solution and pre-
cipitates. Such investigations can be performed using a
chemical or electrochemical procedure, selectively dis-
solving the steel matrix, and separating the undissolved
particles from the matrix by filtration.51 Figure 7 shows
the XRD precipitate peaks of the filter residue of Alloy 3
(Fe-0.25C-0.075Nb-0.075Al). The diffraction peaks in
this Figure 7 match well with NbC(N) and AlN. There-
fore, the precipitates rich in Nb and Al observed in
Figure 7 correspond to these two types of precipitates.
In an alloy containing aluminium and without tita-
nium or niobium, AlN precipitation occurs in the auste-
nitic or ferritic regions. Several investigations report a
fine precipitation («1 μm) with a large number density of
nitride particles for steels containing between 29.96
mg/L and 299 mg/Lnitrogen.15 This precipitation is
known to have significant effects upon recrystallization
and austenite grain growth.52 The strengthening effect of
Nb micro-alloying on steels also occurs with ferrite grain
refinement due to austenite grain-boundary pinning,
retardation of recrystallization and precipitation
strengthening with an increase of the steel’s strength.53
The elemental composition of the micro-alloyed
powder metallurgy steels was performed using the X-ray
fluorescence technique (XRF), which has many advan-
tages: it is fast, accurate, non destructive and has a limit
of detection in the range of few ppm for most elements.54
For these reasons, the XRF analysis method is widely
used in many fields such as metallurgy, geology and mi-
neralogy, the food industry and environmental manage-
ment. However, most routine steel analyses involve
standard wet-chemical methods or inductively coupled
plasma atomic emission. These methods are destructive
and require dissolution of the alloy and a long sample-
preparation time. The use of the X-ray fluorescence
technique is very attractive in many fields and especially
for metal and alloy analyses.55 The sample preparation
for XRF is relatively simple, so that it requires less time
and work. For example, when the solid sample is homo-
geneous, then it only needs polishing to be ready for
analysis.56 In this experimental work, micro-alloyed PM
steels were analysed using the XRF technique to deter-
mine their elemental compositions. The chemical com-
position in % of mass fractions of Alloy 4 (Fe-0.25C-
0.1Nb-0.1Al) are presented in Table 5. It can be noted
from Table 5 that most elemental compositions obtained
conform to the values claimed by micro-alloyed steels.
Table 5: Chemical composition of Alloy 4 obtained by XRF analysis
Tabela 5: Kemijska sestava zlitine 4, dobljene z XRF analizo
No. Com-ponent Result Unit El. line
Inten-
sity
w/%
normal
Analyz-
ing
depth
1 Fe 99.0452 w/% 99.0452
2 C 0.2653 w/% C-KA 0.0400 0.2653
3 Mn 0.2223 w/% Mn-KA 0.7588 0.2223 0.0281
4 Al 0.0974 w/% Al-KA 0.1224 0.0974 0.0009
5 Nb 0.0918 w/% Nb-KA 1.1323 0.0918 0.0658
Figure 8 shows the tensile-testing fracture surfaces
of Alloy 1 and Alloy 3. The changes were observed on
the fracture surface of the PM steel and micro-alloyed
PM steels sintered at 1350 °C with respect to the micro-
voids’ size, shape and depth. The mode of fracture for
Alloy 1 (Figure 8a) is purely ductile. This is evident
from the presence of numerous dimples along with fine
and rounded pores. The mechanisms of the fracture were
void formation and coalescence in the necks between
adjoining fracturing microvolumes. Some microvoids
pre-existed in the material, others nucleated and grew
from defects or inclusions. The microvoids nucleate at
strain discontinuities, such as those associated with
S. GÜNDÜZ et al.: EFFECT OF THE ADDITION OF NIOBIUM AND ALUMINIUM ON THE MICROSTRUCTURES ...
646 Materiali in tehnologije / Materials and technology 50 (2016) 5, 641–648
Figure 7: X-ray diffraction pattern of residues collected in the filter
for Alloy 3
Slika 7: Rentgenogram ostankov zbranih na filtru pri zlitini 3
a-priori defects (pores, microcracks), second-phase par-
ticles, inclusions, grain boundaries, dislocations pile-up.
As the strain increases, the microvoids grow, coalesce,
which involves plastic deformation, and eventually form
a continuous fracture surface.57
However, Alloy 3 showed dimples and cleavage
facets (Figure 8b) indicating that the fracture is of the
mixed type. Large voids were also observed in the alloy.
These voids are an indication of the removal of NbC or
AlN particulates in the minor fracture surfaces. The
pull-out of the NbC or AlN particulates on the crack
faces indicates a possible mechanism for crack-tip bridg-
ing in the micro-alloyed PM steel, as well M. Hajisafari
et al.58 showed that void nucleation and growth in mi-
cro-alloyed steel are a function of the shape factor of
second-phase particles and the mismatch between the
second-phase particle and the metal matrix. Voids firstly
nucleate around second-phase particles and consequently
grow.
The Al-N-based particle inside the large microvoids
of micro-alloyed PM steel is clearly shown in the SEM
fractograph and the corresponding EDS results in Figure
8c. M. A. Erden et al.32 investigated the tensile behaviour
of sinter-forged Ti-alloyed PM steel and they observed a
mixed (ductile-brittle) type of fracture in the micro-
alloyed PM steel with Ti due to the formation of carbides
and the carbide pull-off during heavy deformation.
4 CONCLUSIONS
Micro-alloyed PM steels with three different volume
fractions of Nb-Al were processed through cold pressing
and sintering in an argon atmosphere. The important
findings obtained can be summarised as follows:
Micro-alloyed PM steels were analysed using the
XRF technique to determine their elemental composi-
tions. The results show that most elemental compositions
obtained confirm to the values claimed by micro-alloyed
steels.
Micro-alloying with Nb and Al requires high-tem-
perature sintering to dissolve the Nb and Al in the
austenite. The relative density of the sintered Nb-Al
micro-alloyed PM steels can reach 94 % after sintering at
1350 °C.
The addition of Nb-Al can improve the strength of
micro-alloyed PM steel through precipitation hardening
and microstructure refining. The NbC(N) and AlN preci-
pitates inhibit the grain growth during the sintering step,
leading to enhanced strength.
The EDS and XRD analyses of the micro-alloyed PM
steels revealed Nb- and Al-rich particles in micro-
alloyed PM steel. The presence of Nb and Al elements in
the particles indicates that NbC(N) and AlN formed
during sintering and/or precipitate during the cooling
after sintering.
The analysis of the fracture surface indicated that PM
steel without micro-alloying addition exhibited ductile
fracture. In the case of the micro-alloyed PM steels the
mode of fracture is the mixed type with ductile and
brittle regions. This could be attributed to the pull-out of
NbC or AlN particles during deformation, indicating the
possible mechanism for crack tip is bridging in the
micro-alloyed PM steels.
Acknowledgements
The authors gratefully acknowledge financial support
from TÜBÝTAK, the Research Project of The Scientific
and Technological Research Council of Turkey with the
NO: 113M350.
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648 Materiali in tehnologije / Materials and technology 50 (2016) 5, 641–648
L. A. DOBRZAÑSKI et al.: CHARACTERISTICS OF DYE-SENSITIZED SOLAR CELLS WITH CARBON NANOMATERIALS
649–654
CHARACTERISTICS OF DYE-SENSITIZED SOLAR CELLS WITH
CARBON NANOMATERIALS
ZNA^ILNOSTI NA FIKSIRANO BARVO OB^UTLJIVIH SOLARNIH
CELIC Z OGLJIKOVIMI NANOMATERIALI
Leszek Adam Dobrzañski, Agnieszka Mucha, Marzena Prokopiuk vel Prokopowicz,
Marek Szindler, Aleksandra Dryga³a, Krzysztof Lukaszkowicz
Silesian University of Technology, Konarskiego St. 18A, 44-100, Gliwice, Poland
krzysztof.lukaszkowicz@polsl.pl
Prejem rokopisa – received: 2014-07-30; sprejem za objavo – accepted for publication: 2015-09-21
doi:10.17222/mit.2014.134
Dye-sensitized photovoltaic cells consisting of a layered structure have been developed for 20 years and they are a basis for the
new development trend of photovoltaics. One of the examined aspects of their application is building-integrated photovoltaics.
Dye-sensitized photovoltaic cells (DSSCs) were developed by Michael Grätzel and Brian O’Regan in 1991 and have been
intensively examined ever since. Because of their low production costs, easy transfer, the relatively high efficiency of the photon
conversion to the current and an easy production technology, dye-sensitized cells might represent an alternative to silicon cells.
Basically, a dye-sensitized photovoltaic cell consists of five elements: a mechanical base covered with a layer of transparent
conductive oxides (TCOs), a semiconductor film, e.g., TiO2, dye absorbed on the semiconductor’s surface, an electrolyte
including a redox carrier, and a counter electrode suitable to regenerate a redox carrier, e.g., platinum. As part of this work we
produced dye-sensitized solar cells. First, the glass with transparent conductive oxides was thoroughly cleaned. Then, the glass
with TCO was coated with a layer of TiO2 using the doctor-blade technique, and fired in a furnace at 450 °C. The plate prepared
in this way was then sensitized in a ruthenium-based dye. The counter electrode was obtained by applying it on the glass with
TCO carbon nanomaterials, including graphite, carbon black and carbon nanotubes. The photo-anode and the counter electrode
were combined and between them was injected the redox electrolyte. This paper provides an analysis of the microstructure and
electrical properties of nanostructural coatings with the carbon nano-element of the integrated dye-sensitized photovoltaic cells.
Keywords: dye-sensitized solar cell, carbon elements, counter electrode
Na fiksirano barvo ob~utljive fotovoltai~ne celice sestojijo iz plastovite strukture in so zadnjih dvajset let tematika razvoja na
tem podro~ju, predstavljajo namre~ nov razvojni trend v fotovoltaiki. Eden od raziskovanih vidikov njihove uporabe je
fotovoltaika, integrirana v zgradbe. Na fiksirano barvo ob~utljive fotovoltai~ne celice (DSSC), sta razvila Michael Grätzel in
Brian O’Regan leta 1991 in so od tedaj pogost predmet raziskav. Zaradi nizkih stro{kov njihove proizvodnje, enostavne
prenosljivosti, relativno visoko u~inkovite konverzije fotonov v tok in enostavne proizvodnje, so na fiksirano barvo ob~utljive
celice lahko nadomestek silicijevim celicam. V osnovi na fiksirano barvo ob~utljive fotovoltai~ne celice sestojijo iz 5
elementov: mehanska podlaga je prekrita s plastjo prosojnih, prevodnih oksidov TCO, polprevodnega sloja, npr. TiO2, fiksne
barve absorbirane na povr{ini polprevodnika, elektrolita z vklju~no redoks nosilcem nasprotna elektroda, ki je sposobna
regeneracije redoks nosilca, kot je npr. platina. Kot del tega dela, so bile izdelane na fiksirano barvo ob~utljive solarne celice.
Najprej je bilo steklo s presevnimi prevodnimi oksidi dobro o~i{~eno. Nato je bilo steklo s TCO, z uporabo tehnike kirur{kih
no`ev, prekrito s plastjo TiO2 in `gano v pe~i pri 450 °C. Tako pripravljena plo{~a je bila ob~utljiva za barve na osnovi rutenija.
Nasprotna elektroda je bila dobljena z nanosom TCO ogljikovih nanomaterialov, vklju~no z grafitom, ~rnim ogljikom in
ogljikovimi nanocevkami na steklu. Fotoanoda in nasprotna elektroda sta bili kombinirani in med njiju je bil vbrizgan redoks
elektrolit. ^lanek predstavlja analizo mikrostrukture in elektri~nih lastnosti prevlek z nano strukturo, z ogljikovimi
nanoelementi integrirane fotovoltai~ne celice, ob~utljive na fiksirano barvo.
Klju~ne besede: solarna celica, ob~utljiva na fiksirano barvo, ogljikovi elementi, nasprotna elektroda
1 INTRODUCTION
The increasing rate of energy consumption in the
world is directly related to the increase in the human
population. The progress of civilization is associated
with an increase in the demand for energy, and particu-
larly the most useful of its forms – electricity. Nowadays,
mankind consumes 13,500 GW of power. The solar
power reaching the Earth is 170,000,000 GW. If even a
part of this energy could be used it could reduce envi-
ronmental problems involving atmospheric pollution,
arising as a result of the excessive use of conventional
energy sources. Photovoltaics is an alternative and envi-
ronmentally friendly technology for electricity produc-
tion. Photovoltaic cells (also known as solar cells or PV
cells) are used to convert solar energy into electricity,
and this phenomenon is called the photovoltaic effect.
The main advantage that enhances the development of
organic photovoltaics is the potentially several times
lower price of energy production per unit cell area than
conventional solar cells based on silicon. Other
advantages include 1–4 much better aesthetics compared
to silicon solar cells, low toxicity, high transparency,
possibility to choose the colour, flexibility, low dead-
weight, low power loss due to the unfavourable angle of
incidence of the sunlight, which is used in BIPV
(Building Integrated Photovoltaics), working under
reduced radiation (cloudy, darkening), where these cells
Materiali in tehnologije / Materials and technology 50 (2016) 5, 649–654 649
UDK 67.017:620.3:621.383.51 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)649(2016)
have a much better performance than silicon solar cells,
a low production price due to the use of small amounts
of material and the simplicity of production technolo-
gies, and the performance is independent of temperature
changes in the range of 25–65 °C.
The construction of a dye-sensitized solar cell is
based on a layered structure, which consists of two trans-
parent glass plates with a Transparent Conductive Oxide
(TCO) on it, placed parallel to each other and spaced
about 40 μm apart (Figure 1). On one of the plates is
applied a nanocrystalline titanium oxide layer coated
organometallic photosensitive dye (photosensitizer) –
this system retrieve in the cell function photo-anode
(illuminated anode). On the surface of the second plate
glass with TCO is usually nanoplatinum, which is a cata-
lytic layer – this system is in the cell cathode. The space
between the plates is filled with an electrolyte containing
a redox system I–/I3–. Each component shows the depen-
dence between many other materials. If at least one
element in dye-sensitized solar cells is changed, e.g., the
dye, the particle size of the TiO2, the film thickness, or
the composition of the electrolyte, the DSSC cell re-
quires adjustment to ensure optimal system perfor-
mance.5,6
The operating principle of a dye-sensitized solar cell
is shown in Figure 2. The dye and the electrolyte are
essential components of the cell.
The task of the counter electrode is to gather elec-
trons flowing from the outer current and to catalyse the
reduction of the triiodide ions. Platinum is the most
common material used as a counter electrode. Despite
the fact that platinum shows a high catalytic activity, its
shortage in resources, high costs and corrosion possibi-
lity through a triiodide solution, inhibit its application on
a large scale in the future.7 For this reason, there is a
need for research on alternative materials that are charac-
terized by electrochemical activity and chemical stabi-
lity. So far platinum8, carbon,9,10 conductive plastics,11
CoS,12 WO2,13, Mo2C and WC,14 TiN,15 have been used
as the counter electrodes. There are many publications
about the methods of shaping the surface and structure of
materials to improve their properties.16–18
Carbon nanotubes conduct electricity. They are
almost transparent, flexible and strong, which makes
them the ideal material for transparent electrodes for
DSSC. The only drawback is that photo-generated
charge carriers in the nanotube may recombine with ions
in the dye, which reduces the power-conversion efficien-
cy of the solar cell.
In the present work inexpensive and available carbon
materials such as carbon black, graphite, and carbon
nanotubes have been used as alternative materials to
platinum because of their high corrosive resistance, high
reactivity for triiodide reduction and low costs.7,9,10,19–25
The disadvantages in catalytic activity in comparison to
platinum may be compensated for by increasing the
active surface of a catalytic layer using the porous struc-
ture.7,25 Forming high-quality carbon bands on a sub-
strate gives us prospects for using carbon as a counter
electrode.
L. A. DOBRZAÑSKI et al.: CHARACTERISTICS OF DYE-SENSITIZED SOLAR CELLS WITH CARBON NANOMATERIALS
650 Materiali in tehnologije / Materials and technology 50 (2016) 5, 649–654
Figure 1: Construction of dye-sensitized solar cell6
Slika 1: Zgradba solarne celice ob~utljive na fiksirano barvo6
Figure 2: The operating principle of dye-sensitized solar cells6
Slika 2: Princip delovanja solarnih celic, ob~utljivih na fiksirano barvo6
Figure 3: The production steps of dye-sensitized solar cells
Slika 3: Proizvodni koraki pri solarnih celicah, ob~utljivih na fiksi-
rano barvo
2 EXPERIMENTAL PROCEDURE
The production steps for dye-sensitized solar cells
were shown in Figure 3. As the counter electrodes we
used three carbon materials: carbon black, graphite and
carbon nanotube. The dye-sensitized solar cells have the
following arrangement of layers (Figure 4):
• FTO glass/TiO2/dye/electrolyte/carbon black/FTO
glass,
• FTO glass/TiO2/dye/electrolyte/graphite/FTO glass,
• FTO glass/TiO2/dye/electrolyte/nanotube/FTO glass,
where:
FTO glass is glass with a layer of fluorine-doped tin
oxide (FTO).
2.1 Fabrication of the photoanode
Glass plates with dimensions 30 mm × 30 mm
(10
/sq.) were used. In order to remove any surface con-
tamination and to degrease, the FTO glasses were dipped
and held in an ultrasonic, deionized water, acetone,
ethanol and isopropanol. In order to reduce the active
surface of the dye-sensitized solar cells the FTO glass
was covered with a layer of tape (Figure 5). In order to
obtain a TiO2 paste nitric acid with pH 3–4 was mixed
with ethanol.26–28 To the solution was added titanium
oxide nanopowder. This solution was stirred until a
uniform paste was obtained. A drop of TiO2 paste was
applied to the FTO glass and then it was uniformly
spread over the glass surface using the doctor-blade
technique (Figure 6). After removing the tape, the glass
plate with the titanium paste was annealed in a furnace at
450 °C in air and then air-cooled. In order to sensitize
the electrode it was immersed in the dye solution with
absolute ethanol for 24 h at room temperature, without
access to light. Once removed, the electrode was washed
with ethanol to remove any excess dye and allowed to
dry. A glass plate with a layer of titanium oxide, with
and without dye, is shown in Figure 7.
2.2 Fabrication of the counter electrode
The preparation of the carbon nanotubes’ surface
layers on the silicon substrates was performed using
EasyTube®2000. During the process the default settings
were used. EasyTube®2000 is an advanced, turnkey,
thermal catalytic chemical vapor deposition CVD pro-
cess tool for the synthesis of a wide variety of nano-
structured materials. The catalyst needed for the CNT
growth was a transition metal plus iron, and the catalyst
was introduced to the process together with the CNT
precursor.
It this case the counter electrodes were prepared by
spring carbon nanotubes on the FTO glass. The CNT
solution was prepared by direct mixing of the acid highly
conductive PEDOT:PSS and applied on the glass surface
with FTO. The second electrode that was used for this
experiment was the electrode with carbon black. The
L. A. DOBRZAÑSKI et al.: CHARACTERISTICS OF DYE-SENSITIZED SOLAR CELLS WITH CARBON NANOMATERIALS
Materiali in tehnologije / Materials and technology 50 (2016) 5, 649–654 651
Figure 5: FTO glass covered with a layer of tape
Slika 5: FTO-steklo, prekrito s plastjo traku
Figure 6: Photo-anode preparation
Slika 6: Priprava fotoanode
Figure 4: Schema of DSSC with: a) carbon black, b) graphite, c) car-
bon nanotube
Slika 4: Shema DSSC z: a) ~rnim ogljikom, b) grafitom, c) ogljiko-
vimi nanocevkami
carbon black is cheap in industrial mass production. Ti is
also used in printing toners, so we can easily spray it on
to FTO glass.
A thin layer of carbon materials, i.e., carbon black
(Figure 8a), graphite (Figure 8b), carbon nanotube
(Figure 8c), were deposited on the FTO glass that was
previously cleared of impurities.
2.3 Fabrication of the DSSC
An anode and a cathode were combined with the
sealing strip, which simultaneously serves as a separator.
Careful electrode bonding is a very important step in the
preparation of DSSC cells. It prevents the leakage and
evaporation of the electrolyte.
The last step was the placement the electrolyte,
which is a solution of iodine and iodide in an organic
solvent containing a redox couple I–/I3–, between the
photo-anode and a counter electrode.
2.4 Measurements
Because the used carbon elements and titanium oxide
consist of nano-metric structural units or single carbon
layers, modern research equipment was used. First of all,
an Atomic Force Microscope, a High Resolution Trans-
mission Electron Microscope and a Scanning Electron
Microscope.
Atomic force microscopy (AFM, XE-100, Park
Systems) was used to observe the surface morphology of
the TiO2 layer with and without the dye.
The counter electrodes’ morphology was observed
using the scanning electron microscope (SEM; SUPRA
35, ZEISS).
The High Resolution Transmission Electron Micro-
scope S/TEM (TITAN 80-300, FEI) was used to observe
the surface morphology of the carbon nanotubes.
The voltages of the DSSCs were recorded using a
multimeter (Meter Link Extech EX845) as a source
measure unit, which was connected between the FTO
glass and the counter electrode.
3 RESULTS AND DISCUSSION
Figure 9 shows the AFM images of the TiO2 layer
and the dye-adsorbed TiO2 layer. The size and the dis-
L. A. DOBRZAÑSKI et al.: CHARACTERISTICS OF DYE-SENSITIZED SOLAR CELLS WITH CARBON NANOMATERIALS
652 Materiali in tehnologije / Materials and technology 50 (2016) 5, 649–654
Figure 8: Counter electrode with a layer of: a) carbon black, b) graphite, c) PEDOT: PSS with carbon nanotubes
Slika 8: Nasprotna elektroda s plastjo: a) ~rnega ogljika, b) grafita, c) PEDOT: PSS z ogljikovimi nanocevkami
Figure 9: AFM images of the: a) TiO2 layer and b) dye-adsorbed TiO2
layer
Slika 9: AFM-posnetek: a) plast TiO2 in b) s fiksirano barvo adsorbi-
rana plast TiO2
Figure 7: FTO glass with TiO2 layer: a) before dying, b) after dying
Slika 7: FTO steklo s plastjo TiO2: a) pred su{enjem, b) po su{enju
tribution of the particles characterize the dye-absorbed
surface of the photo-anode. The SEM images show that
TiO2 surface with the absorbed dye is smoother than the
surface without the dye.
Figures 10 to 12 show the SEM surface images of
various counter electrodes on the FTO glass. Figure 10
shows the microstructure of the carbon black, where a
rectangular atomic arrangement can be observed. Be-
cause the fluorine-doped tin oxide film deposited on the
glass has a rough surface, in Figure 11 where the gra-
phite is shown, we can also observe the rough pyramid
microstructure. Figure 12 shows the microstructure of
the carbon nanotubes’ electrode, where severely agglo-
merated CNTs are observable. We can see that the sur-
face with the CNT and the highly conductive
PEDOT:PSS has a uniform structure and fewer pores
were generated in the structure compared with the micro-
structure of the carbon black or the graphite.
By adding carbon nanotubes to the electrode it is
possible to obtain a larger surface area and therefore a
larger contact surface than the graphite and carbon black
counter electrode.
Table 1 shows the electrical parameters by means of
voltage for the three dye-sensitized solar cells with three
different counter electrodes. As was expected after the
SEM images, the highest voltage comes from the DSSC
with carbon nanotubes as a counter electrode.
Table 1: Electrical properties of the investigated DSSCs
Tabela 1: Elektri~ne lastnosti preiskovane DSSC
Type of counter electrode Voltage (mV)
carbon black 40
graphite 35
carbon nanotubes 52
4 CONCLUSIONS
In this study, for the purpose of decreasing the cost of
dye-sensitized solar cells (DSSCs), we investigated the
effect of carbon materials, i.e., carbon nanotubes, gra-
phite and carbon black, added as a counter electrode.
Three different low-cost DSSCs (CNT-counter electrode
DSSC, carbon black-counter electrode DSSC, graphite-
electrolyte DSSC) were fabricated.
The costly platinum electrode is able to be replaced
by the modified counter electrode (using CNTs), with
only little change in the efficiency of the cell. The effect
of the carbon nanotubes on the performance of the DSSC
showed that the DSSC fabricated with the CNT and
PEDOT:PSS had the highest photovoltaic performances.
These results are attributed to increasing the surface area
of the counter electrode.
The conductivity of the carbon black is lower than
carbon materials such as the graphite and carbon nano-
tubes, because those carbon nanotubes have a high con-
ductivity and a large surface area. The literature proves
that the carbon nanotubes with PEDOT:PSS layers have
a good adhesion, so it can be used on the DSSC
L. A. DOBRZAÑSKI et al.: CHARACTERISTICS OF DYE-SENSITIZED SOLAR CELLS WITH CARBON NANOMATERIALS
Materiali in tehnologije / Materials and technology 50 (2016) 5, 649–654 653
Figure 12: SEM images of carbon nanotubes
Slika 12: SEM-posnetek ogljikovih nanocevk
Figure 10: SEM image of carbon black
Slika 10: SEM-posnetek ~rnega ogljika
Figure 11: SEM image of graphite
Slika 11: SEM-posnetek grafita
electrodes. In the future it is planned to perform tests of
the adhesion for the products’ layers of low-cost DSSCs.
The investigated effect of the CNT counter electrodes
on the efficiency of DSSC showed that these kinds of
solar cells had the best efficiency for the counter electro-
des.
Acknowledgment
The project was funded by the National Science
Centre on the basis of the contract No. DEC-2013/09/B/
ST8/02943.
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654 Materiali in tehnologije / Materials and technology 50 (2016) 5, 649–654
S. E. ERDOGAN, U. HUNER: THE EFFECT OF THE WELDING PARAMETERS AND THE COUPLING AGENT ...
655–662
THE EFFECT OF THE WELDING PARAMETERS AND THE
COUPLING AGENT ON THE WELDING OF COMPOSITES
VPLIV PARAMETROV VARJENJA IN SREDSTVA ZA SPAJANJE
NA VARJENJE KOMPOZITOV
Selcuk Ertugrul Erdogan, Umit Huner
Trakya University, Faculty of Engineering, Department of Mechanical Engineering, 22050 Edirne, Turkey
umithuner@trakya.edu.tr
Prejem rokopisa – received: 2015-03-11; sprejem za objavo – accepted for publication: 2015-09-14
doi:10.17222/mit.2015.059
This paper presents an experimental investigation of the welding of a glass-fiber-reinforced PP composite. The goals of this
paper are to investigate the issues of local changes of the welding strength that depend on the heating time and the coupling
agent (MAPP). Composite samples were prepared by using an extruding (for mixing) process and a hot-press method. The PP
matrix was reinforced by unidirectional short glass fibers. The welding process for the specimens was carried out using a
non-contact heated tool butt welding process. Tensile and fatigue tests were conducted to investigate the effects of the heating
time parameter and the coupling agent. The highest weld strength dependent on the heating time was achieved with 94 %
relative to the base strength of the material. The fatigue behavior of short-fiber-reinforced thermoplastic composites
(polypropylene/20 % of volume fractions of E-glass fiber) is presented in terms of stress versus the number of cycles to failure.
The specimens were fatigue tested at various percentages of their static tensile strengths at a load ratio R = 0.1 and frequency f=
5 Hz. An indefinite fatigue life was obtained at 35 % of the static damage initiation load for glass-fiber-reinforced specimens.
Then, these specimen’s maximum welding strengths and fatigue properties that were dependent on the heating time were
compared.
Keywords: plastic material, composite, heated tool, welding process, reinforcement, fatigue, polypropylene
^lanek predstavlja preiskavo varjenja PP kompozita, oja~anega s steklenimi vlakni. Cilj ~lanka je bil preiskati vpliv lokalnih
sprememb na trdnost zvara, ki je odvisna od ~asa ogrevanja in sredstva za spajanje (MAPP). Kompozitni vzorci so bili
pripravljeni z uporabo metode ekstruzije in vro~ega stiskanja. PP osnova je bila oja~ana z usmerjenimi kratkimi steklenimi
vlakni. Postopek ~elnega varjenja vzorcev je bil izveden z brezkontaktno ogrevanim orodjem. Izvedeni so bili natezni preizkusi
in preizkusi utrujenosti, da bi ugotovili vpliv ~asa ogrevanja in sredstva za spajanje. Najve~ja trdnost zvara v odvisnosti od ~asa
ogrevanja je bila 94 % trdnosti osnovnega materiala. Obna{anje pri utrujanju termoplasti~nega kompozita (polipropilen/20 %
prostorninskih dele`ev E-steklenih vlaken), oja~anega s kratkimi vlakni, je prikazano na krivulji utrujanja kot odvisnost
napetosti od {tevila ciklov. Utrujenost vzorcev je bila preizku{ana pri razli~nih odstotkih stati~ne natezne trdnosti, pri hitrosti
obremenjevanja R = 0,1 in frekvenci f = 5 Hz. Zdr`ljivost na utrujanje s steklenimi vlakni oja~anih vzorcev je bila dobljena pri
35 % nazivne stati~ne obremenitve. Primerjane so bile maksimalne trdnosti zvarov, z obna{anjem pri utrujanju v odvisnosti od
~asa ogrevanja.
Klju~ne besede: plasti~ni material, kompozit, ogrevano orodje, postopek varjenja, oja~anje, utrujenost, polipropilen
1 INTRODUCTION
In high-technology applications, a composite is suit-
ably qualified due to the fact that it has a higher strength
and a better stiffness-to-weight ratio. The vital properties
of thermoplastic composites include higher damage tol-
erance, corrosion resistance, higher impact resistance
and enhanced fatigue life. Due to their recyclable and
re-formable natures, thermoplastic composites are se-
lected for environmentally benign applications.1–4
The welding process is one of the preferred methods
to realize the assembly structure of thermoplastic com-
posites. This method eliminates disadvantages such as
stress concentration and the galvanic corrosion of me-
chanical fastenings. Also, the thermoplastic matrix has
its own specific welding parameters with much the same
reinforcements. In the welding zone the dispersion of the
reinforcement and orientation could be affected by the
welding pressure, heating, heating time and some similar
properties, which results in the corruption on unity for
the reinforced thermoplastic.5–7
Some researchers have worked on different welding
methods for thermoplastic composites. At present, there
is still not enough literature available related to the pro-
cess parameters and their influence on the joint’s
strength.
C. B. Bucknall et al.8 reported that the weld strengths
of glass-fiber-reinforced polypropylene were strongly af-
fected by the hot-plate temperature, heating time, and
melt flow during welding. K. V. Stokes5 has studied the
fatigue life of vibration-welded unreinforced polycarbo-
nate (PC), polyetherimide (PEI), modified polyphe-
nylene oxide resin (M-PPO) and poly (butylenes
terephthalate) (PBT), under tension–tension loading at
R = 0.1. The first three polymers are amorphous and
PBT is semi-crystalline. K. V. Stokes5 found that the ra-
tio of the endurance limit stress to the tensile strength
was 0.29 for PC, 0.34 for PEI, 0.22 for M-PPO and 0.31
Materiali in tehnologije / Materials and technology 50 (2016) 5, 655–662 655
UDK 621.791.92:669.018.25:678.742.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)655(2016)
for PBT. T. T. Lin et al.9 investigated the effect of weld-
ing parameters on the non-contact hot-plate butt welding
of polypropylene. For a given hot-plate temperature, an
optimum heating time and forging pressure were found.
M. Watson et al.10 investigated the parameters of heated
tool welding. The heated tool temperature has been
found to be a less critical parameter than either the heat-
ing pressure or time.
These studies mostly concentrate on the parameters
of the welding process, such as the heating time, the tool
temperature and the pressure. But they lack the compari-
son of composites that had a chemical treatment of a ma-
trix material like maleic anhydride (MA). There is a need
to better understand the influence of various processing
parameters and the coupling agent (MAPP) on the joint
properties of hot tool welded thermoplastic composites.
In accordance with previous studies published in the lit-
erature about the fatigue life of SGFR thermoplastics,
the fatigue scattering is small, especially in comparison
with the fatigue of metallic materials.
The final goal of this study is to gain a better under-
standing of the effects of using a coupling agent (MAPP)
and the heating time parameter on the welding properties
of hot tool welded composites. For constant temperatures
and reinforced MAPP with a different ratio, thermoplas-
tic composite component’s butt welded joint strength,
failure strain, modulus of elasticity and fatigue properties
that depend on the heating time and cyclic number were
compared by using the tensile and the fatigue-tests
method. Non-welded material data are also included as a
reference. The results are analyzed using curves for the
stress versus the number of cycles to failure (S–N) for
the fatigue test.
2 MATERIALS AND EXPERIMENTAL
PROCEDURES
2.1 Materials
The resins used in this study were commercially
available, virgin-grade polypropylene (PP) S.R.L., poly-
propylene-grafted maleic anhydride (PP-g-MA (Sigma
Aldrich), MA content = 1 % of mass fraction) chopped
into strands of glass fiber PA2-4.5 (Cam Elyaf Inc.). Ta-
ble 1 lists the properties of the resins as provided by the
resin producer. Glass-fiber-reinforced PP granules were
prepared with a lab-type single-screw extruder (L/D: 28).
And then the granules were shaped as 200×200 mm2
plates using a hot press. The samples were obtained from
the plates with a cutting press. ISO 527 tensile-test pro-
cedures were used in this investigation. Three identical
samples of each composition (Table 2) were measured
and the average values were reported.
Table 2: Thermoplastic composites used in the experiments
Tabela 2: Termoplasti~ni kompoziti, uporabljeni za preizkuse
Material type PP (w/%) MA-g-PP (w/%) GF (w/%)
PPv (virgin) 100 – –
PP20GF 80 0 20
MA2.5PPGF 77.5 2.5 20
MA5PPGF 75 5 20
2.2 Welding method
In the non-contact hot-tool welding process, the parts
being welded are placed near the hot tool separated from
it by a distance referred to as the non-contact gap. The
hot tool is removed during the change-over phase. Pres-
sure is applied to hold the parts in close contact during
weld cooling and solidification. The heat is transferred
by thermal radiation and convection. The process is oth-
erwise identical to hot-tool welding: the hot plate is re-
moved in the changeover phase, and pressure is applied
to achieve close contact as the weld cools and solidifies.8
The non-contact hot-tool or hot-plate welding has pro-
cessing parameters that influence the weld strength,
which include the size of the non-contact gap, the platen
temperature, heating time, change-over time, weld pres-
sure and duration. A butt-type joint has a lower weld line
strength at a low welding pressure.11 The welding and
tensile test steps are shown as an example in Figure 1.
During the welding process, the heat transfer raises
the temperature of the part and the resulting thermal ex-
pansion causes a small rightward (away from the hot-tool
surface) motion in the part and fixture. When the surface
temperature reaches the melting point of the plastic the
part surface begins to melt. The externally applied pres-
sure causes the molten material to flow laterally out-
wards, thereby inducing a leftward motion of the part.8
In non-contact heated tool welding, the contami-
nation of the weld surfaces is minimized, the heating is
S. E. ERDOGAN, U. HUNER: THE EFFECT OF THE WELDING PARAMETERS AND THE COUPLING AGENT ...
656 Materiali in tehnologije / Materials and technology 50 (2016) 5, 655–662
Table 1: Thermoplastic materials and glass-fiber properties used in the experiments
Tabela 1: Lastnosti termoplasti~nega materiala in steklenih vlaken, uporabljenih pri preizkusih
Material Producer Density MFI (g/10min) Melting temp. Vicat soft. temp.(°C)
Tensile strength
(MPa)
Polipropilen
(S.R.L) ROM Petrol 0.90 20 165 132 32
PP-g-MA Sigma Aldrich 0.95 115 152 147 –
Material Producer Dimensions Density(gr/cm3)
Tensile strength
(MPa)
Melting temp.
(°C) Annotations
Glass Fiber Cam Elyaf A.ª. D: 10.5 μmL: 4.5 mm 2.54 3450 1722
treated 0.6 %
silane
uniform, and a small weld bead is produced, providing
good, consistent weld strengths.11 An important aspect
was the heat-soak time, which was influential in obtain-
ing high joint strengths. Ideally, the stops should be as
close to the joint interface as possible, consistent with
bringing the joint interface into close contact with the
heating element.
The experimental activity was carried out with the
aim of evaluating the static and fatigue behavior under
tension loadings of single lap welded joints in composite
materials and to investigate the mechanics of damage
evolution.
In this study the mechanical performance of the weld
was obtained using rectangles of injection-molded sam-
ples that were welded together. The dimensions of the
samples have enough tolerance for the welding process.
Two parts were joined under heat that was generated by
the stainless-steel hot-tool plate with dimensions of
40 mm × 2 mm × 100 mm. The weld line temperature
was manually controlled with an InfraRed thermometer
(CEM DT-8835, K-Type). The sets of these test samples
and test parameters are given in Table 3. The weld area
is equal to 40 mm2. The welding pressure was held at
constant values of 0.5 MPa and 4 MPa. For each type of
thermoplastic composite a welding temperature of 260 °C
was applied.12 The non-contact gap was 1 mm. A con-
stant heating displacement and weld displacement were
maintained during the experiments. The heating tempera-
ture was chosen as an optimum value that is common to
welding on all types of thermoplastic composite welding.
The range for the heating time was 40 s to 70 s. The
weld was always situated in the middle of the specimen,
perpendicular to the load line.
2.3 Tensile test
In this study the tensile tests were performed using an
Instron Universal Testing Machine Model 8501, equip-
ped with a 500-kg load cell, a strain-gauge extensometer
(Instron, model 2620, UK) after conditioning at 23±2 °C
according to the ISO 527 standard. The cross-head speed
used for the type IA tensile specimens was 5 mm/min.
All the samples were 150 mm in length with a bonded
butt (flat) type. The tests were performed in triplicate
and the results reported are the arithmetic average of the
parallel samples.
The dog-bone-shaped sample is routed down to a
standard ISO 527 tensile test specimen with a butt joint
at its center. The tensile sample, which has a transverse
butt weld at mid-length, is then subjected to a constant-
displacement-rate tensile test in which the strain across
the weld is monitored with an extensometer. In this way
the average failure strain across the weld over a 25 mm
gauge length can be monitored.
2.4 Fatigue test
Load-controlled fatigue tests were performed in the
tension-tension mode at ambient temperature (approxi-
mately 23 °C). The specimens were tested under a
sinusoidal waveform at various loads between 35 % and
80 % of the static damage initiation load according to
ASTM D 3479 (tension-tension fatigue behavior). The
tests were conducted under a load ratio R = 0.1 and a
frequency f = 5 Hz. No significant heating was noticed
during the fatigue testing. When specimen failure could
S. E. ERDOGAN, U. HUNER: THE EFFECT OF THE WELDING PARAMETERS AND THE COUPLING AGENT ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 655–662 657
Figure 1: Welding and tensile test procedure of a sample: a) divide
injection-molded sample into two parts, b) joint samples by welding,
c) tensile testing of welded part, d) failure on welded parts after tensile
test
Slika 1: Postopek varjenja in preizku{anja vzorca: a) dva dela razde-
ljenega tla~no litega vzorca, b) zvarjen vzorec, c) natezni preizkus
zvara, d) poru{en zvar po nateznem preizkusu
Table 3: Some parameters of the experimental procedure for the study and sample codes (welding temp. 260 °C)
Tabela 3: Nekateri parametri preizkusov in oznake vzorcev (temperatura varjenja 260 °C)
Materials Heating time(s)
Welding pressure (MPa)
Part I/Part II Materials
Heating time
(s)
Welding pressure (MPa)
Part I/Part II
PPv
40 0.5/4 MPa
MA2.5PPGF
40 0.5/4 MPa
50 0.5/4 MPa 50 0.5/4 MPa
60 0.5/4 MPa 60 0.5/4 MPa
70 0.5/4 MPa 70 0.5/4 MPa
PP20GF
40 0.5/4 MPa
MA5PPGF
40 0.5/4 MPa
50 0.5/4 MPa 50 0.5/4 MPa
60 0.5/4 MPa 60 0.5/4 MPa
70 0.5/4 MPa 70 0.5/4 MPa
not be obtained within 1 million cycles, the test was ter-
minated and an indefinite fatigue life was reported. Bet-
ween 10 and 15 specimens were tested for each material.
The fracture surfaces of the broken specimens were
observed visually and using scanning electron micro-
scopy (SEM). The samples were first sputter coated with
a fine layer of gold under vacuum for 60 s. An accelerat-
ing voltage of 20 kV was used to collect the SEM
images.
3 RESULTS AND DISCUSSION
3.1 Tensile test results
The results of the static tests are presented in terms of
both nominal tensile stress on the adherents and shear
stress on the adhesive. The aim of this is to provide
information on the load-carrying capability of the joints
and the adhesive properties. It is a well-known fact that
the welding quality is influenced by many processing
factors, some of these being the welding time and the
welding pressure. It is therefore important to explore the
best combination of these factors to obtain the best
welding result. Considering the used filler content, it was
verified that for all the tested composites the tensile
strength increases by the welding time and considering
the shape of filler content, the type of fiber for the filler
content achieved a relatively high strength compared to
the other reinforced PP composites. This happens be-
cause the fiber adhesion to the matrix is enough, when
compared with the others, to increase the matrix ten-
sion’s transfer efficiency through the interface and that is
essential to get an improvement of the mechanical pro-
perties in the composite.
For each time period (40 s, 50 s, 60 s, 70 s) three
parallel samples were used and the results reported are
the arithmetic average of the parallel samples. The joint
S. E. ERDOGAN, U. HUNER: THE EFFECT OF THE WELDING PARAMETERS AND THE COUPLING AGENT ...
658 Materiali in tehnologije / Materials and technology 50 (2016) 5, 655–662
Table 4: Comparing the results of joint strength and failure strain for the welded samples
Tabela 4: Primerjava rezultatov trdnosti spojev in raztezek pri poru{itvi zvarjenih vzorcev
Materials Heatingtime(s)
Joint strength
(MPa) I Sd
*** w/b*
I
Joint strength
(MPa) II Sd
w/b
II
PPv
(Bulk strength 32 MPa)
40 17.84 7.55 0.55 14.53 5.86 0.45
50 20.16 9.77 0.63 17.66 6.53 0.55
60 23.42 8.02 5.73 20.45 7.94 0.63
70 26.11 8.51 5.81 23.71 12.84 0.74
PP20GF
(Bulk strength 41 MPa)
40 24.15 10.87 4.58 20.31 9.12 0.49
50 26.34 11.14 7.55 22.59 6.35 0.55
60 33.18 15.97 9.11 29.16 8.74 0.71
70 37.77 16.23 8.92 33.94 10.81 0.82
MA2.5PPGF
(Bulk strength 44 MPa)
40 30.24 10.88 1.68 27.76 8.54 0.63
50 35.61 12.01 6.80 32.59 8.62 0.74
60 38.22 14.11 5.77 38.93 10.88 0.88
70 40.78 18.14 6.92 40.56 11.99 0.92
MA5PPGF
(Bulk strength 48 MPa)
40 36.45 10.78 1.75 32.98 9.21 0.68
50 39.27 12.28 4.81 35.26 10.54 0.73
60 41.48 17.33 3.86 36.48 13.33 0.76
70 46.25 19.71 6.96 40.58 10.27 0.92
Materials Heatingtime(s)
Failure strain
(%) I Sd
*** w/b**
I
Failure strain
(%) II Sd
w/b
II
PPv
(Bulk material strain
0 = 4.56%)
40 3.51 1.3 0.77 3.14 1.72 0.69
50 3.76 0.7 0.82 3.05 2.02 0.67
60 4.11 2.2 0.90 2.87 0.97 0.63
70 4.18 2.4 0.92 2.56 1.12 0.56
PP20GF
(Bulk material strain
0 = 0.58 %)
40 0.51 0.16 0.88 0.40 0.14 0.69
50 0.45 0.04 0.78 0.34 0.12 0.59
60 0.41 0.13 0.71 0.29 0.09 0.50
70 0.36 0.08 0.62 0.26 0.11 0.45
MA2.5PPGF
(Bulk material strain
0 = 0.35 %)
40 0.22 0.04 0.63 0.16 0.06 0.46
50 0.17 0.04 0.49 0.11 0.08 0.31
60 0.11 0.03 0.31 0.06 0.02 0.17
70 0.08 0.01 0.23 0.04 0.02 0.11
MA5PPGF
(Bulk material strain
0 = 0.24 %)
40 0.17 0.04 0.71 0.14 0.07 0.58
50 0.15 0.06 0.63 0.11 0.03 0.46
60 0.07 0.06 0.29 0.04 0.01 0.17
70 0.05 0.02 0.21 0.02 0.01 0.08
* w/b (Relative strength),** w/b (Relative strain),*** Sd (Standard deviation)
strength of the welded samples and the bulk strength of
unwelded samples are compared in Table 4. All of the
welded assemblies of virgin and reinforced polypropy-
lene failed at the weld lines. All the glass-fiber-rein-
forced parts were found to fracture at the weld interfaces.
The results suggested that the hot-plate-welded
MA5PPGF composite parts exhibited the highest joint
strength on 0.5 MPa weld pressure, followed by the
MA2.5PPGF and PP20GF composites. The welded
virgin polypropylene showed the lowest joint strengths.
It is clear that a weld does not represent just a discon-
tinuity in the material, but is the source of an extended
mechanical disturbance, which strongly influences the
local material’s behavior, even at a distance outside the
weld.13 It can be noticed that for the higher MA-g-PP/GF
content the local strain (Table 4) in the accompanying
zones is lower, which corresponds to the lower com-
pliance of the reinforced material. Also, in this case the
lower pressure results in a local strain super elevation at
the weld. The range of fiber orientation inside the weld is
only very low, so that the material showing a higher
degree of deformation and damage at the same weld
pressure governs the mechanical behavior for the whole
weld range.
As the amount of MAPP increases the joint strength
increases to a 0.5-MPa weld pressure as reported in
Table 4, but the joint strength of the samples decreases
with less than 4 MPa weld pressure. Similar results have
been reported by J. S. Liu, H. F. Cheng13 Under the
influence of the welding pressure, the plasticized mater-
ial in the weld zone flows. With increasing penetration,
fibers push out of the surface. This fiber bridging causes
a high weld strength. But increasing the welding
pressure causes a high degree of orientation transverse to
flow (injection) to the original fiber alignment. This in
turn leads to a decrease in the joint strength.
The modulus (E) of the material is also reported in
Figure 2. With respect to PPv, they have higher moduli
and strengths, but they break at a lower strain. The
differences in the E of the PP-based composites cannot
be ascribed to their different composition because in that
case, E should monotonically decrease with the welding
pressure.
Increasing the hot-plate heating time increases the
temperatures of the materials. The temperature provides
the necessary movement of the melt flow, and time is
needed for the diffusion to occur. A higher material
temperature therefore aids increasing the weld strengths.
Homogenous filler orientation such as fiber on the
welding zone and matrix material dominates the stress
behavior. Fiber reinforcements lead to an increase of the
melting point of the matrix. The random orientation of
glass fiber on the welding zone affects the welding
strength depending on the heating time. Increasing the
heating time thereby the melting length leads to increas-
ing fiber orientation horizontal to the direction of the
load. With the increasing melting length fibers move to
the welding zone while the melt flows out. The results
that depend on heating time have provided a comparison
for a variety of filler-reinforced PP composites. M. Gehde
et al.14 have investigated the effect of melting length and
joining length for different heating times (0 s, 60 s and
120 s) on the material reinforced PP with a random glass
mat (PP-GM). PP-GM attains the highest strength at a
low joining length. Varying the melting length does not
affect the maximum strength of 28 MPa. In this study,
46.25 MPa, the highest strength value of the welding,
was obtained for a low welding pressure, which means a
low joining length. This strength was reached at a time
of 60 s. The higher strength values from the literature
can be said to be associated with the use of MA. MA-g-
PP usage, increasing the fiber adhesion, also on the
welding interface, has fulfilled its function of improving
the strength in the basic structure. The MA-g-PP (5 %),
approximately 20 % difference between using and not
using the welding strength of the composite were deter-
mined at low/high weld pressures.
The SEM observation at the fracture surface suggests
that the welded MA5PPGF materials have fewer fibers
S. E. ERDOGAN, U. HUNER: THE EFFECT OF THE WELDING PARAMETERS AND THE COUPLING AGENT ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 655–662 659
Figure 2: a) Modulus of welded samples under 0.5 MPa weld
pressure, b) modulus of welded samples under 4 MPa weld pressure
Slika 2: a) Modul zvarjenih spojev pri 0,5 MPa tlaku zvara, b) modul
zvarjenih vzorcev pri 4 MPa tlaku zvara
that are oriented horizontal to the direction of load 0.5
MPa, when compared to that of the welded MA5PPGF
composites (Figure 3) on the pressure load of 4 MPa.
The joint strengths of the welded MA5PPGF at a 4-MPa
weld pressure may thus be inferior to that of the
MA5PPGF composites welded a 0.5-MPa weld pressure.
Many transverse fibers were widely detached and some-
times fibers in longitudinal direction of the specimen
were broken. Additionally, the squeezed-out plastic melt
generates a flash, resulting in a sharp transition in the
cross-section of the product.8 For a welding pressure of 4
MPa, the SEM micrographs show that the fibers that
bridge the weld zone are shorter than those at a pressure
of 0.5 MPa. Therefore, the region with fibers protruding
from the weld is smaller, and the fiber orientation in the
direction of the flash increases.
3.2 Fatigue test results
There are many factors that govern the fatigue beha-
vior of discontinuous fiber-reinforced polymer-matrix
composites. Some of these include the processing con-
ditions, the fiber length and the orientation with respect
to the loading axis, the properties of the matrix,
interfacial properties, and testing conditions. The fibers
tend to orient along the flow direction, which leads to
superior mechanical properties along the flow direction.
As the degree of fiber disorientation with respect to the
S. E. ERDOGAN, U. HUNER: THE EFFECT OF THE WELDING PARAMETERS AND THE COUPLING AGENT ...
660 Materiali in tehnologije / Materials and technology 50 (2016) 5, 655–662
Figure 4: S-N curves at R = 0.1 for glass-fiber-reinforced PP with weld pressure of 0.5 MPa (LP), 4.0 MPa (HP) and PPv (unreinforced), (arrows
indicate unbroken specimens)
Slika 4: S-N krivulje pri R = 0,1 za PP oja~an s steklenimi vlakni, pri tlaku varjenja 0,5 MPa (LP), 4 MPA (HP) in PP (neoja~an), (pu{~ice ka`ejo
na neporu{ene vzorce)
Figure 3: a) Increasing fiber orientations in the direction of the flash, b) MA5PPGF composite interface with fiber orientation at 4-MPa welding
pressure, c) MA5PPGF composite interface, fibers oriented horizontal to the direction of load 0.5 MPa
Slika 3: a) Nara{~anje usmerjenosti vlaken v smeri svetlobe, b) MA5PPGF stika kompozita z usmerjenostjo vlaken pri tlaku varjenja 4 MPa, c)
MA5PPGF stik kompozita, vlakna so usmerjena horizontalno v smeri obremenitve z 0,5 MPa
loading axis increases, the strength of the composite is
increasingly dominated by the matrix and interfacial
properties.15–20
All the fatigue data were preliminarily analyzed
using the classic stress-life approach and drawing the
fatigue curves based on the nominal stress and the num-
ber of cycles to failure, identified as the complete separa-
tion of the joints. Figure 4 presents the fatigue perfor-
mance for all four materials as a plot of the maximum
cyclic stress versus the number of cycles to failure, on a
semi-log scale (S–N plot). Figure 4 also reports the S-N
curves of glass-fiber-filled polypropylene samples that
were welded under a 0.5-MPa weld pressure and 4 MPa
weld pressure. The vertical axis represents the maximum
cyclic stress and the horizontal axis is the number of
fatigue cycles to failure in this figure.
The stress-life approach allows the influence of the
design parameters under investigation to be clearly
identified.21–23 In fact, an increase in the fatigue strength
of the joints and therefore in their load capability can be
observed when the welding pressure decreases for the
joint process. Even the influence of the weld pressure
itself can be easily identified in Figure 4. In terms of
absolute stress, as could be expected from the static tests,
MA5PPGF specimens exhibited better performance.
Analyzing the fatigue diagrams, the low-pressure weld-
ing process led to a higher fatigue strength. At the
welding pressure of 0.5 MPa, the highest strength values
were achieved with the MA5PPGF material. Based on
this result, in the welding zone, the fiber orientation is
parallel to the applied force direction. The samples were
welded under a pressure of 4 MPa, and the test results
showed that the fatigue strength decreased. While
MA5PPGF material strength again showed the highest
fatigue strength of 26.88 MPa at about 9500 cycle, this
load level represents 80 % of the static damage initiation
load, which is similar to the findings of previous studies
on the welding of thermoplastic composites.15,16 K. J. Tsang
et al.16 also reported that the fatigue life of specimens
welded at low pressure was 50 % greater than that of the
specimens welded at high pressure, this study’s results
showed an 18 % fatigue life increase at low pressure.
The low-weld-pressure condition shows more fibers ori-
ented perpendicular to the weld plane, which may have
increased the weld strength and resulted in a longer
fatigue life. When the effects on the strength of the
welding reinforcement by MAPP are analyzed, using
MAPP material at a low rate (2.5 %) provided protection
of the weld strength at a low weld pressure. The
MAPP2,5PP material showed similar results (Figure 5)
under a fatigue load, neither at a low welding pressure
nor at a high welding pressure This tendency is also seen
in the tensile-strength value (Table 4).
Figures 5a to 5c show the fatigue fracture surfaces of
reinforced PP specimens welded at low and high pres-
sures. The low-weld-pressure condition shows more
fibers oriented perpendicular to the weld plane (Figure
5b), which may have increased the weld strength and re-
sulted in a longer fatigue life. This is consistent with the
fracture surface observed in previous studies.8 For rein-
forced material, the fatigue crack propagation process is
more local and on a smaller scale compared with that of
the unreinforced material due to the presence of glass
fibers.
4 CONCLUSIONS
This study has examined the effect of different pro-
cessing parameters on the joint strength of hot-plate
welded thermoplastic composites, including the heating
time, weld pressure and using MA-treated polypropy-
lene. Four materials were used in the study: virgin
polypropylene, 20 % glass fiber, 2.5 % MAPP (and 20 %
S. E. ERDOGAN, U. HUNER: THE EFFECT OF THE WELDING PARAMETERS AND THE COUPLING AGENT ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 655–662 661
Figure 5: a) The matrix adhering to the fiber (at high pressure), b) matrix sticking to the fiber can be clearly seen and the orientation of the fibers
perpendicular to the weld line. The sample exhibited a similar fracture surface as shown by the tensile specimen with the matrix sticking to the
surface, c) fibers oriented to the center of the weld interface (rectangle area).
Slika 5: a) Matrica, ki se je prijela na vlakno (pri visokem tlaku), b) osnova, ki se je prijela na vlakno se dobro vidi in orientacija vlakna je
pravokotna na linijo zvara. Vzorec ka`e enako povr{ino preloma kot pri nateznem preizku{ancu, kjer se je osnova prilepila na povr{ino, c) vlakna
v sredini zvara (pravokotna podro~ja).
glass fiber), 5 % MAPP (and 20 % glass fiber) reinforced
polypropylene composites.
With the increased welding time, the tensile strength
increased and this was found to be reversed by corre-
lating the welding pressure.
At a low welding pressure the MAPP material pro-
vides increased tensile strength of the glass-fiber-rein-
forced PP at about a rate of 50 %. Under a high welding
pressure the MAPP increases the tensile strength, but it
is seen that when using 2.5 % of that 5 % of any con-
tribution when used.
The fatigue behavior gets worse and the fatigue limit
of the tested material decreases when the weld pressure
increases. As well as this, the welding time shows simi-
lar trends to the weld pressure. While contributing posi-
tively to the use of MAPP fatigue strength, the strength
drop caused by the pressure increase was partially
blocked.
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662 Materiali in tehnologije / Materials and technology 50 (2016) 5, 655–662
S. POURANVARI et al.: CHEMICAL CROSS-LINKING OF CHITOSAN/POLYVINYL ALCOHOL ...
663–666
CHEMICAL CROSS-LINKING OF CHITOSAN/POLYVINYL
ALCOHOL ELECTROSPUN NANOFIBERS
KEMIJSKO ZAMRE@ENJE ELEKTRO SPREDENIH NANOVLAKEN
IZ HITOSAN/POLIVINIL ALKOHOLA
Sara Pouranvari1, Firouz Ebrahimi2, Gholamreza Javadi1, Bozorgmehr Maddah3
1Islamic Azad University, Department of Biology, Science and Research Branch, Tehran, Iran
2IHU, Basic Sciences Faculty, Biology Research Center, Tehran, Iran
3IHU, Basic Sciences Faculty, Department of Chemistry, Tehran, Iran
febrhimi@ihu.ac.ir
Prejem rokopisa – received: 2015-04-17; sprejem za objavo – accepted for publication: 2015-07-08
doi:10.17222/mit.2015.083
Electrospun nanofibrous scaffolds have great potential for many biomedical applications. In the present study, we fabricated and
characterized chitosan/polyvinyl alcohol (Chi/PVA) nanofibrous scaffolds through electrospinning. Cross-linking was per-
formed using chemically with 5 % glutaraldehyde vapor. The morphology and chemical banding of the electrospun nanofibers
before and after cross-linking were evaluated using scanning electron microscopy (SEM) and Attenuated Total Reflectance-
Fourier Transform InfraRed (ATR-FTIR) spectroscopy. SEM micrographs and FTIR spectra showed that the cross-linking
process was accomplished successfully. With the biocompatibility and non-toxicity of chitosan and PVA, it is expected that this
electrospun nanofibrous scaffold could be an excellent candidate for biomedical applications.
Keywords: electrospinning, chitosan, polyvinyl alcohol, cross-linking
Mre`e iz elektro spredenih nanovlaken imajo velik potencial za uporabo v biomedicini. V {tudiji smo izdelali in karakterizirali
nanovlaknasto mre`o, izdelano z elektro predenjem nanovlaken iz hitosan/polivinil alkohola (Chi/PVA). Zamre`enje je bilo
izdelano s pomo~jo kemijske metode s 5 % glutaraldehidne pare. Morfologija in kemijsko povezovanje elektro predenih
nanovlaken, pred in po zamre`enju, sta bila ocenjena z uporabo vrsti~nega elektronskega mikroskopa (SEM) in z metodo z
oslabljenim odbojem infrarde~e spektroskopije s Fourierjevo transformacijo (ATR-FTIR). SEM-posnetki in FTIR-spekter sta
pokazala, da je bil postopek zamre`enja uspe{no dose`en. Glede na biokompatibilnost in netoksi~nost hitosana in PVA se
pri~akuje, da bodo mre`e iz elektro spredenih nanovlaken odli~en element za uporabo v biomedicini.
Klju~ne besede: elektro-predenje, hitosan, polivinil alkohol, zamre`enje
1 INTRODUCTION
Electrospinning is a simple, versatile and cost effec-
tive method for forming non-woven fibrous scaffolds.
Technically, the electrospinning process uses a high
voltage source to draw a polymer fluid into fine fibers
which are deposited on a collector.1 In recent years, the
use of electrospun nanofibers for biomedical applications
such as tissue engineering2, wound dressing3, protein
immobilization4, materials for artificial blood vessels5,
barriers for the prevention of induced adhesion after
operation6, and vehicles for drug or gene delivery7 has
attracted a great deal of attention from scientists. Elec-
trospinning of synthetic and natural polymers has been
reported for collagen8, gelatin9, silk fibroin10, polygly-
colide (PGA)11, polylactide (PLA)12 and poly(-capro-
lactone) (PCL)13, polyurethane14, poly(vinylalcohol)15,
PEO16, polydioxanone17, and polyphosphazene deriva-
tives.18 Furthermore, the blending of two or more
polymers and copolymerization are effective methods for
the preparation of composites with new and desirable
properties. Obviously, by adjusting the ratio of the
components, structure and morphology of the nanofibers
and the biological properties of the electrospun scaffolds
can be tailored to the desired traits and functions.1 For
example PLGA7, P(LA-CL) copolymers,19 and mixtures
of collagen with elastin,20 gelatin with PCL,9 chitosan
with poly(ethylene oxide) (PEO)21 and chitosan with
PVA22 have all been utilized to fabricate electrospun
nanofibrous scaffolds for biomedical applications.
In biomedical applications, after electrospinning,
different cross-linking methods can be uses to provide
stabilization against aqueous environments for those
scaffolds produced from aqueous soluble polymers (For
example: PVA).
In the present study, electrospinning of a chitosan and
PVA blend was performed. Chitosan was selected due to
its cytocompatibility, biocompatibility, biodegradability
and antibacterial activity.23 PVA was used due to its bio-
compatibility, biodegradability, non-toxicity, chemical
resistance, and good fiber-forming properties.24
2 MATERIALS AND METHODS
2.1 Materials
PVA (average molecular weight of 70000–100000
g/mol) and chitosan (medium molecular weight) were
purchased from Sigma-Aldrich (St. Louis, MO). Acetic
Materiali in tehnologije / Materials and technology 50 (2016) 5, 663–666 663
UDK 620.192.4:660.017:678.744.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)663(2016)
acid and glutaraldehyde were obtained from Merck
(Germany).
2.2 Preparation of the solutions
Chitosan and PVA were dissolved in 50 % aqueous
solution of acetic acid at a concentration of 2 % mass
fraction and 15 % mass fraction, respectively. The
chitosan solution and PVA solution were mixed together
with a weight ratio of 40/60 (Chi/PVA) under magnetic
stirring at 60 °C.
2.3 Preparation of nanofibrous membranes
The optimal conditions for the electrospinning were
as follows: 25 kV applied voltage, 15 cm tip-to-collector
distance, and 1 ml/h flow rate. Moreover, a 5 ml syringe
with a 21 gauge stainless-steel needle was used for the
delivery of the polymer solution via a syringe pump.
2.4 Crosslinking of nanofibrous membranes
Samples were exposed to the 5 % glutaraldehyde
(GA) vapor at room temperature for 48 h for cross-link-
ing to stabilize them against aqueous media solubility
and enhance their biomechanical properties biomedical
applications. After crosslinking, the samples were care-
fully washed several times with 2 % glycine for the inac-
tivation and removal of the GA.25
2.5 Characterization of nanofibrous membranes
The morphology and microstructure of the electro-
spun nanofibers before and after cross-linking were
determined by Scanning Electron Microscopy (SEM).
The average diameter of fibers were calculated using the
ImageJ (US National Institute of Health, Bethesda, MD)
image analysis program by analyzing at least 50 fibers in
ten SEM micrographs. The chemical structures of the
chitosan and PVA powders and Chi/PVA nanofiber mem-
branes before and after cross-linking were investigated
by Attenuated Total Reflectance-Fourier Transform
InfraRed (ATR-FTIR) spectroscopy (Bruker Tensor 27,
USA). FTIR spectra were obtained in the 4000 cm–1 to
400 cm–1 wavenumber range, with the data analyzed
using OPUS software.
RESULTS AND DISCUSSION
3.1 Morphology of the nanofibrous scaffold
SEM images of Chi/PVA nanofibers before and after
GA cross-linking are shown in Figure 1. As seen in
Figure 1a, relatively fine, continuous, uniform fiber-
structures (no bead), and randomly oriented fibers were
obtained. The average fiber diameter was found to be
180±2.28 nm. In accordance with the relatively fine
fibers fabricated, it is expected that a suitable porosity
exists for biomedical applications.
3.2 Crosslinking of nanofibrous scaffold
The SEM micrographs of the cross-linked nanofibers
after immersion in water (at least 48 h) are shown in
Figure 1b. As PVA is a water-soluble polymer, cross-
linking should be performed for the use of Chi/PVA
nanofibers in biomedical applications. Several studies
have been reported on GA cross-linking for medical
application. For example, Jafari et al. used a saturated
vapor of a 25 % GA aqueous solution for cross-linking
of chitosan-gelatin electrospun nanofibers.26 In another
study cross-linking of electrospun water-soluble carbo-
S. POURANVARI et al.: CHEMICAL CROSS-LINKING OF CHITOSAN/POLYVINYL ALCOHOL ...
664 Materiali in tehnologije / Materials and technology 50 (2016) 5, 663–666
Figure 1: SEM micrographs of electrospun Chi/PVA nanofibers, after
48 h immersion in water at 37 °C: a) before GA cross-linking and b)
after GA cross-linking
Slika 1: SEM posnetek elektro spredenih Chi/PVA nanovlaken, po
48 h namakanja v vodi s 37 oC: a) pred GA zamre`enjem in b) po GA
zamre`enju
xyethyl chitosan/poly (vinyl alcohol) nanofibrous mem-
brane towards wound dressings for skin regeneration was
performed using a GA vapor.27 Unlike the previous work,
in the present study, the cross-linking was performed
using a lower concentration of GA vapor (5 %). Despite
using this low concentration, it was proved that the
fabricated membranes’ structure is stable in an aqueous
solution. Moreover, according to the SEM micrographs
shown in Figure 1, the porous structure of the fabricated
membranes remained intact implying that they are insol-
uble in water. The cross-linking mechanism of chitosan
and PVA with GA is shown in Figure 2.28,29
3.3 ATR-FTIR analysis
FTIR spectra were taken of the electrospun nano-
fibers before and after cross-linking, to assess their
chemical groups. The FTIR spectrum of the Chi/PVA
blended nanofibers before cross-linking, is shown in
Figure 3. The two peaks at 1423 cm–1 and 1565 cm–1
arise fromcarboxylic acid and symmetric deformation of
–NH3+ groups due to ionization of primary amino groups
in the acidic medium, respectively. The peak at 1703 cm–1
is attributed to the carboxylic acid dimer.22 In this study,
this peak is due to the acetic acid utilized for dissolving
the chitosan. The peak located at 1244 cm–1 is related to
the C–O of the CH2OH chitosan group forming a hydro-
gen bond with the OH of PVA, confirming the fabrica-
tion of Chi/PVA blend nanofibers.30 The FTIR spectra of
the Chi/PVA blended nanofibers before cross-linking is
given in Figure 3. Chemical crosslinking of the chito-
san/PVA is verified by the peak located at 1586 cm–1
attributed to the C-N band. All chitosan-derived blends
cross-linked with GA, have shown the presence of the
imine (C=N) band. The imine band was formed by the
nucleophilic reaction of the amine from chitosan with the
aldehyde group of GA.31 Due to imine band instability
with temperature and pH, this group can transform to a
C-N group.32
4 CONCLUSION
In this work, nanofibrous Chi/PVA was fabricated via
electrospinning and stabilized by chemical cross-linking
using 5 % GA. The porous structure of the electrospun
scaffolds, antimicrobial properties of the chitosan and
chemical resistant traits of PVA, make our fabricated
electrospun scaffold an excellent candidate for biomedi-
cal applications. However, in vitro and in vivo experi-
ments for evaluation of the biocompatibility of these
Chi/PVA nanofiberous membranes is necessary.
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S. POURANVARI et al.: CHEMICAL CROSS-LINKING OF CHITOSAN/POLYVINYL ALCOHOL ...
V. YÝLMAZ: INVESTIGATION OF HOLE PROFILES IN DEEP MICRO-HOLE DRILLING OF AISI 420 ...
667–675
INVESTIGATION OF HOLE PROFILES IN DEEP MICRO-HOLE
DRILLING OF AISI 420 STAINLESS STEEL USING
POWDER-MIXED DIELECTRIC FLUIDS
PREISKAVA PROFILOV LUKNJE PRI GLOBOKEM VRTANJU
MIKRO LUKNJE V AISI 420 NERJAVNEM JEKLU S POMO^JO
DIELEKTRI^NE TEKO^INE S PRIME[ANIM PRAHOM
Volkan Yýlmaz
Gazi University, Faculty of Technology, Manufacturing Engineering, Ankara, Turkey
volkan@gazi.edu.tr
Prejem rokopisa – received: 2015-05-18; sprejem za objavo – accepted for publication: 2015-07-13
doi:10.17222/mit.2015.100
A series of experiments were carried out to drill deep micro-holes, using Electrical Discharge Machining (EDM) with different
concentrations of carbon powder/dielectric fluid mixture into, and varying machining parameters such as elecrode rotation speed
and dielectric fluid spray pressure. Four dielectric fluid concentrations, three electrode rotation speeds, and three dielectric fluid
spray pressure settings were tested; the resulting holes were investigated with respect to their average radial overcut and surface
roughness (Ra) values. Study results indicate that as carbon powder concentrations, dielectric fluid spray pressure and electrode
rotation speeds are increased, ARO values increased while Ra values are observed to decrease. It was determined that
improvements in Ra and ARO values may be attained by the right configuration of machining parameters and the selection of the
appropriate mixtures of carbon powder with dielectric fluid.
Keywords: AISI 420 stainless steel, deep micro-hole drilling, electro discharge machining, hole profile, surface roughness
Izvedena je bila vrsta eksperimentov vrtanja globokih mikroizvrtin s pomo~jo EDM in razli~ne koncentracije prahu ogljika,
prime{anega dielektri~ni teko~ini pri spreminjanju parametrov obdelave, vklju~no s hitrostjo vrtenja elektrode in tlaka curka
dielektri~ne teko~ine. Preizku{ene so bile {tiri koncentracije dielektri~ne teko~ine, tri hitrosti vrtenja elektrode in trije razli~ni
tlaki curka dielektri~ne raztopine; nastale luknje so bile preiskovane glede na povpre~ni radij preseka in vrednosti hrapavosti
povr{ine (Ra). Rezultati ka`ejo, da nara{~anje koncentracije prahu ogljika, tlak curka teko~ine in hitrost rotacije elektrode
pove~uje vrednosti ARO, medtem ko se vrednost Ra zmanj{uje. Bilo je ugotovljeno, da je mogo~e izbol{anje vrednosti Ra in
ARO mogo~e dose~i s konfiguracijo parametrov obdelave in z izbiro primerne koli~ine prahu ogljika, ki je prime{an dielektri~ni
teko~ini.
Klju~ne besede: AISI 420 nerjavno jeklo, vrtanje globokih mikroizvrtin, obdelava z elektroerozijo, profil luknje, hrapavost
povr{ine
1 INTRODUCTION
In recent years, the use of Electrical Discharge Ma-
chining (EDM) for drilling small diameter holes in very
hard materials has gathered momentum, as the materials
used in newly developed systems exhibit advanced
mechanical properties and conventional drilling methods
fall short in such materials. In EDM, there is no direct
contact between the drilling tool and the workpiece, a
significant benefit which eliminates the problems
associated with physical contact.1–8 The EDM method is
based on principles of thermal and electrical conductivity
whereby small regions on the workpiece surface are
removed by melting and vaporization. EDM systems
allow the drilling of specific diameter holes provided that
the material is electrically conductive.9 The EDM me-
thod is used for applications such as machine assembly
points, fuel injection spray nozzles and aircraft engine
cooling holes. EDM hole drilling has found applications
at both the macroscale and microscale, with good surface
quality and acceptable tapering.10
Due to electrode wear and lateral erosion, EDM holes
exhibit some, albeit a small, amount of tapering. This is a
complication for the EDM industry to overcome.
Research is being conducted into reducing the tapering,
and disparities between hole entry and exit diameters are
reported to have been reduced to acceptable levels.11
Additionally, it has been reported that the problem may
be surmounted with the use of coating on the tool elec-
trode.12,13 The parameters used for EDM hole drilling
have a significant effect on hole tapering and surface
roughness.14–16 Rotation of the electrode has enhanced
hole drilling performance in EDM operations.17,18 The
type of dielectric fluid and the manner in which it is
applied have improved the process.19,20 In addition to the
tapering of the holes in EDM hole drilling, the surface
roughness of hole walls is also of importance. It has been
reported that the electrode plunge action is facilitated by
rotational movement, and that surface roughness is
simultabneously improved.21 Various types of electrodes
are used in EDM and by changing their polarization, the
surface roughness of holes have been reported to be
Materiali in tehnologije / Materials and technology 50 (2016) 5, 667–675 667
UDK 620.1:669.14.018.8:621.95 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)667(2016)
reduced.22 It is also reported that ancillary equipment can
achieve significant enhancements in surface roughness.23
In EDM hole drilling in tandem with grinding, improved
surface roughness results have been achieved.24 Addi-
tionally, system hybridization has been implemented
resulting in the process duration being decreased while
quality has been enhanced.25
One factor affecting machining performance in EDM
is the conductive powder mixed with the dielectric fluid.
Thanks to the conducting particulate, EDM processing,
including hole drilling, is improved. The presence of
conductive particulates not only improves hole drilling
performance, but also has positive effects on hole taper-
ing and hole surface roughness.26–28 However, despite the
existence of several studies concerning hole drilling
using powder-mixed dielectric fluids, the benefits and
drawbacks of the method have yet to be established.
This study focuses on carbon (C) powder mixed with
the dielectric fluid, as well as two specific EDM para-
meters, and attempts to determine their effect on hole
profile and surface roughness. In a survey of the litera-
ture, no studies have been identified targeting hole pro-
file and surface roughness properties of the EDM
micro-hole drilling method using either classical or
powder-mixed EDM in AISI 420 material; this study
attempts to fill the gap.
2 EXPERIMENTAL SETUP
For machining, a Furkan brand EEI M50A model
EDM machine was used. A custom attachment was
mounted on the EDM tool head to allow for the rotation
of the electrode at various min–1 (revolutions per minute)
speeds and to provide for the dielectric fluid to be
delivered through the tool, at the required pressure, to the
area to be machined on the workpiece. The attachment is
shown in Figure 1; the EDM machine and the testing
set-up are shown in Figure 2.
Single-hole brass pipe with diameter of 0.5 mm was
used as tool electrode; AISI 420 stainless steel, common-
ly utilized in the manufacturing industry as hot work tool
steel and also used for turbine blade construction, was
selected as the workpiece; distilled water in pure form as
well as distilled water mixed with carbon powder at three
concentrations (2 g/L, 4 g/L and 8 g/L) were used as
dielectric fluids. The chemical composition of AISI 420
is presented in Table 1 and the workpiece, electrode and
carbon powder physical characteristics in Table 2. Table
3 shows the machining parameters used in the testing.
Table 1: Chemical composition of the AISI 420 material in mass
fractions, (w/%)
Tabela 1: Kemijska sestava materiala AISI 420, v masnih dele`ih
(w/%)
C Mn Si P S Cr
0.1 1.00 1.00 0.04 0.030 12.0
The average radial overcut value is determined by
taking the sum of the individual measurements of the
hole diameter along the depth of the hole (Figure 3),
averaging the sum, subtracting the result from the tool
electrode diameter de, and dividing by 2:
ARO( m)
e
μ =
+ + + + + +⎛
⎝
⎜ ⎞
⎠
⎟ −
⎡
⎣⎢
⎤
⎦⎥
d d d d d d d
d
1 2 3 4 5 6 7
7
2
(1)
The surface roughness values were measured using a
Mitutoyo Surftest SJ-210. The values were determined
through an arithmetical averaging of several readings for
each hole.
V. YÝLMAZ: INVESTIGATION OF HOLE PROFILES IN DEEP MICRO-HOLE DRILLING OF AISI 420 ...
668 Materiali in tehnologije / Materials and technology 50 (2016) 5, 667–675
Figure 2: Experiment set-up
Slika 2: Eksperimentalni sestav
Figure 1: Custom attachment to enable electrode rotation
Slika 1: Prilagoditev, ki omogo~a vrtenje elektrode
3 RESULTS AND DISCUSSION
Using EDM, micro-holes with a diameter of 0.5 mm
and depth of 20 mm were drilled into AISI 420 steel
workpiece using four dielectric fluid concentrations
(pure distilled water, and distilled water with carbon
powder concentrations of 2 g/L, 4 g/L, and 8 g/L), three
electrode rotation speeds (100 min–1, 200 min–1 and
400 min–1), three spray pressure settings for the dielectric
fluid (20 bar, 40 bar, and 80 bar), with constant discharge
current (6 A), pulse duration (200 μs) and pulse interval
(100 μs). The effects of the machining parameters on the
ARO values for hole profile and surface roughness have
been investigated. Sample drillings are presented in
Figure 4.
The study was designed to test the dielectric fluid
mixes against dielectric fluid pressure and electrode rota-
tion speed. Test numbers in the following description are
described in Table 3: Four groups of dielectric fluid
were used for testing (Di-0: pure distilled water with no
powder mix, Di-1: distilled water with 2 g/L carbon
powder mix, Di-2: with 4 g/L mix, Di-3: with 8 g/L
mix). Three separate powder mixes were subjected to
three groups of tests (A, B, and C), with three tests in
each test group (1, 2, and 3). For the three test groups,
the pressure setting of the dielectric fluid was varied (A:
20 bar, B: 40 bar and C: 80 bar). For each individual test
within a test group, the electrode rotation speed was
varied (1: 100 min–1, 2: 200 min–1 and 3: 400 min–1). For
each individual test, the average radial overcut (ARO)
V. YÝLMAZ: INVESTIGATION OF HOLE PROFILES IN DEEP MICRO-HOLE DRILLING OF AISI 420 ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 667–675 669
Figure 4: Sample drillings
Slika 4: Vzor~ne izvrtine
Figure 3: Measurements along hole profile
Slika 3: Meritve vzdol` profila izvrtine
Table 2: Workpiece, electrode and carbon powder physical characteristics
Tabela 2: Fizikalne zna~ilnosti obdelovanca, elektrode in prahu ogljika
Material Density (gr/cm3) Specific heatcapacity Electrical resistivity
Thermal conduc-
tivity (W/m K) Melting point (°C)
AISI 420 7.75 460 (J/kg K) 55 (μ
-cm) 24.9 1450
Brass 8.25 0.380 (J/g °C) 6.39 (μ
-cm) 159 900
C powder 2.25 – 12.2 (μ
-m) 100 3527
Table 3: Machining parameters
Tabela 3: Obdelovalni parametri
Discharge current (I) (A) 6
Pulse duration (On-time) (μs) 200
Pulse interval (Off-time) (μs) 100
Electrode rotational speed (min–1) 100, 200, 400
Dielectric spray pressure (bar) 20, 40, 80
Polarity Electrode (+), Workpiece (–)
Dielectric Fluid Pure distilled water, pure distilled water with carbon powder
Workpiece AISI 420
Electrode Brass
Processing depth (mm) 20
Electrode diameter (mm) 0.5
and surface roughness (Ra) values were recorded. Results
of the tests are presented in Table 4.
3.1 Examination of the results obtained for surface
roughness (Ra)
Variations in Ra values in response to machining
parameters and the concentration of carbon powder used
in the dielectric fluid are presented graphically in Fig-
ures 5 and 6. A fundamental goal of EDM processing in
manufacturing is to achieve low surface roughness (Ra)
values. While having a direct relationship with the
machining parameters, Ra values are also influenced by
factors such as the workpiece itself, the type of dielectric
fluid used, the dielectric application method and the type
of electrode used.1,7,29,30,31
Figures 5 and 6 show the Ra values obtained for.
Figure 5 indicates that as the electrode rotation speed is
increased, the Ra values decrease across all tests. This is
a welcome outcome in terms of the overall industry goal
of obtaining lower Ra values in general. A preliminary
explanation offered for the decrease observed in Ra
values is the easier electrode plunge action due to its
rotation. Note that during testing, as the rotation speed
was increased, the drilling duration was observed to
decrease.
In tests conducted using the pure distilled water
dielectric fluid (column Di-0) at 20 bar pressure (test
group A), the Ra value decreased by 8% (from 3.25 μm
to 2.98 μm) as the electrode rotation speed was increased
from 100 min–1 to 200 min–1, the Ra value decreasing by
another 6% (from 2.98 μm to 2.79 μm) as electrode
rotation speed increased from 200 min–1 to 400 min–1
(corresponding to surface roughness results for indivi-
dual tests A-1, A-2, and A-3, under column Di-0). For
tests conducted with dielectric fluid of distilled water
mixed with 2 g/L carbon powder (column Di-1) at 20 bar
pressure (test group A), the decreases were 9 % and 8 %,
for rotation speed increases from 100 min–1 to 200 min–1,
and from 200 min–1 to 400 min–1, respectively (corres-
ponding to surface roughness results for individual tests
A-1, A-2, and A-3, under column Di-1). For the same
V. YÝLMAZ: INVESTIGATION OF HOLE PROFILES IN DEEP MICRO-HOLE DRILLING OF AISI 420 ...
670 Materiali in tehnologije / Materials and technology 50 (2016) 5, 667–675
Figure 6: Dielectric pressure vs. Ra (constant rotation of 100 min–1)
Slika 6: Odvisnost tlaka dielektri~ne teko~ine od hrapavosti povr{ine
Ra (pri konstanti hitrosti vrtenja 100 min–1)
Figure 5: Electrode rotation speed vs. Ra (constant spray pressure at
20 bar)
Slika 5: Odvisnost hitrosti vrtenja elektrode od hrapavosti Ra (kon-
stanten tlak curka 20 bar)
Table 4: Test organization and results
Tabela 4: Zasnova preizkusov in rezultati
Di-0
Pure distilled water
(no powder mix)
Di-1
Distilled water with
carbon
(2 g/L)
Di-2
Distilled water with
carbon
(4 g/L)
Di-3
Distilled water with
carbon
(8 g/L)
Test
group
Dielectric
fluid spray
pressure
(bar)
Electrode
rotation
(min–1)
Average
radial
overcut
(μm)
Surface
roughness
Ra (μm)
Average
radial
overcut
(μm)
Surface
roughness
Ra (μm)
Average
radial
overcut
(μm)
Surface
roughness
Ra (μm)
Average
radial
overcut
(μm)
Surface
roughness
Ra (μm)
A-1
20
100 42 3.25 62 2.46 77 2.14 81 1.92
A-2 200 48 2.98 65 2.24 79 2.06 88 1.87
A-3 400 51 2.79 77 2.06 84 1.92 93 1.75
B-1
40
100 53 3.04 71 2.35 82 2.05 92 1.85
B-2 200 59 2.82 78 2.18 86 1.95 98 1.82
B-3 400 64 2.70 83 2.02 94 1.85 104 1.63
C-1
80
100 65 2.66 86 2.26 91 1.92 101 1.81
C-2 200 72 2.62 88 2.11 94 1.88 105 1.74
C-3 400 76 2.51 94 1.96 102 1.73 112 1.68
test group (A) conducted with dielectric fluid mixed with
4 g/L carbon powder (column Di-2), the decreases were
4 % and 7 %, respectively (results for individual tests
A-1, A-2, and A-3, under column Di-2). For the corres-
ponding tests conducted with dielectric fluid mixed with
8 g/L carbon powder (column Di-3), the decreases were
3 % and 6 % (results for individual tests A-1, A-2, and
A-3, under column Di-3).
When the same set of tests were repeated using 40
bar pressure, the observed decreases in Ra values were as
follows: 7 % and 4 % (B-1, B-2 and B-3 under Di-0),
7 % and 7 % (B-1, B-2 and B-3 under Di-1), 5 % and
5 % (B-1, B-2 and B-3 under Di-2), and 2 % and 10 %
(B-1, B-2 and B-3 under Di-3). For tests at 80 bar
pressure, the observed decreases in Ra values were as
follows: 2 % and 4 % (C-1, C-2 and C-3 under Di-0),
7 % and 7 % (C-1, C-2 and C-3 under Di-1), 2 % and
8 % (C-1, C-2 and C-3 under Di-2), and 4 % and 3 %
(C-1, C-2 and C-3 under Di-3).
Normally, EDM processing results in craters being
formed on the surface of the workpiece due to spark
discharges, the presence of craters increasing the surface
roughness. The decrease in surface roughness values
observed here, in relation to increased rotational speeds
of the electrode, is a significant outcome, as it proves
that electrode rotation improves the process. The expla-
nation for this outcome is that increased electrode rota-
tion speeds reduce the force with which spark discharge
occurs, as rotation of the electrode hinders the direct
flow of spark to a single spot on the workpiece. Instead
of a single spot, the spark is relayed to an area on the
surface of the workpiece. As the spark is spread on the
workpiece, crater depth is reduced, resulting in shallower
craters.2,7,10,18,30 As Ra values are directly related to the
presence of craters on the workpiece, craters with less
depth mean improved (lower) Ra values.
In addition to the electrode rotation speed, the study
also looked at the effects of dielectric fluid spray pres-
sure on surface roughness (Figure 6). It was observed
that Ra values decreased as the dielectric fluid spray pres-
sure was increased. For tests conducted with the elec-
trode rotation speed kept constant at 100 min–1 and using
pure distilled water dielectric fluid (individual tests A-1,
B-1 and C-1 under column Di-0), increasing the spray
pressure from 20 bar to 40 bar resulted in the Ra value
decreasing by 6 %; increasing the spray pressure from 40
bar to 80 bar resulted in the Ra value decreasing by
another 12 %. When the same tests were repeated for
dielectric fluid concentration with 2 g/L carbon, the
decreases were 4 % and 4 % (individual tests A-1, B-1
and C-1 under column Di-1); for dielectric fluid con-
centration with 4 g/L carbon, the results were 4 % and
6 % (same individual tests under column Di-2); and for
dielectric fluid concentration with 8 g/L carbon, the
results were 4 % and 2 % (same individual tests under
column Di-3), respectively.
The decrease in Ra values as a result of the increase in
dielectric fluid spray pressure is also a significant
outcome for the study. It is interpreted that as the spray
pressure is increased, the spark gap is flushed clean to a
greater degree, allowing debris to be removed faster. By
faster flushing, a continuous supply of clean dielectric
fluid is maintained. This prevents the debris buildup
from contributing to unwanted spark discharges, result-
ing in less damage to the workpiece surface. Controlling
the spark discharge is very important to obtain the
desired Ra values.1–3,19,20,31 A spark gap region cluttered
with debris also impedes the plunging action of the elec-
trode, causing larger craters to be formed when the
electrode’s movement is obstructed; increasing the
dielectric fluid spray pressure curtails this problem.
Another factor that affected Ra values was the carbon
powder mixture (Figures 5 and 6). In test A-1 with
dielectric fluid spray pressure of 20 bar and electrode
rotation speed of 100 min–1, conducted using distilled
water with carbon powder at 2 g/L concentration (Di-1),
Ra values were observed to decrease by 24 % compared
to the same test conducted with pure distilled water (A-1
under Di-0).For the same test conducted using distilled
water with carbon powder at 4 g/L concentration (Di-2),
Ra values were observed to decrease by 34 % compared
to the test with pure distilled water (A-1 under Di-0).
When using distilled water with carbon powder mixed in
at 8 g/L concentration (Di-3), Ra values were observed to
decrease by 41 % compared to the same test conducted
with pure distilled water (A-1 under Di-0). In test A-2
with electrode rotation speed of 200 min–1, the observed
decreases in Ra values were 25 %, 30 % and 37 %,
respectively, for Di-1, Di-2, and Di-3, when compared to
the Ra value for Di-0. In test A-3 with electrode rotation
speed of 400 min–1, the observed decreases in Ra values
were 26 %, 31 % and 37 %, respectively, for Di-1, Di-2,
and Di-3, when compared to the Ra value for Di-0.
In the case of a dielectric fluid spray pressure of 40
bar with electrode rotation at 100 min–1 (test B-1), the
observed decreases in Ra values were 23 %, 33 % and
39 %, respectively, for Di-1, Di-2, and Di-3, when com-
pared to the Ra value for Di-0. In the case of dielectric
fluid spray pressure of 80 bar with electrode rotation at
100 min–1 (test C-1), the observed decreases in Ra values
were 15 %, 28 % and 32 %, respectively, for Di-1, Di-2,
and Di-3, when compared to the Ra value for Di-0.
In all of the tests conducted, it was observed that the
carbon powder had a positive effect in decreasing Ra
values. The cause is interpreted to be the wider discharge
gap that is possible due to the use of the carbon powder
mix. A wider discharge gap allows sparks to act on a
larger workpiece surface area, leading to craters with less
depth, and in turn, lower Ra values. During the hole drill-
ing operation, loose debris from the workpiece increases
the spark discharge between the electrode’s lateral sur-
faces and the workpiece. This in turn leads to high points
V. YÝLMAZ: INVESTIGATION OF HOLE PROFILES IN DEEP MICRO-HOLE DRILLING OF AISI 420 ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 667–675 671
on the workpiece surface being melted and vaporized (Fig-
ure 7), enabling lower Ra values to be attained.21,26–28,32
An additional reason for the observed decrease in Ra
values when carbon powder mixes are used is the for-
mation of low-power sparks among the carbon particu-
lates and the workpiece surface. These low-power sparks
cause shallow craters, which improve Ra values.26,33
The answer to whether further increasing carbon
powder concentrations would help boost the observed
decreases in Ra values was investigated through addition-
al testing not included in this study. In the additional
tests, carbon powder concentrations of 16 g/L and 20 g/L
were used, but did not yield further improvements. Thus
this study has been limited to reporting carbon powder
mixed to an 8 g/L concentration. The interpretation for
the limit reached in the efficiencies is two-fold: conges-
tion within the electrode, and short circuits created
within the spark gap. This study has demonstrated that
specific carbon powder concentrations are highly effec-
tive on Ra and that an optimal carbon powder mix is an
effective path to achieving low Ra values.
3.2 Analysis of the findings relating to average radial
overcut (ARO)
The average radial overcut (ARO), an important
parameter in describing the results obtained from EDM
hole drilling operations, is the average difference bet-
ween the diameter of the hole and the diameter of the
electrode. The ARO value provides information on the
hole profile; a lower value indicates that hole diameter is
close to the electrode diameter. Figures 8 and 9 present
the ARO values attained in the study in relation to
machining parameters as well as the carbon powder
concentrations used.
The results presented in Figures 8 and 9 indicate that
increases in electrode rotation speed, dielectric fluid
spray pressure and carbon powder concentrations, in-
crease ARO values. In tests conducted using pure
distilled water dielectric fluid (without carbon powder)
(column Di-0) at 20 bar pressure (test group A), the ARO
value increased by 14 % (from 42 μm to 48 μm) as
electrode rotation speed was increased from 100 min–1 to
200 min–1, and the ARO value increased by another 6 %
(from 48μm to 51 μm) as electrode rotation speed was
increased from 200 min–1 to 400 min–1 (corresponding to
ARO results for individual tests A-1, A-2, and A-3, under
column Di-0). For tests conducted with dielectric fluid of
distilled water mixed with 2 g/L carbon powder (column
Di-1) at 20 bar pressure (test group A), the increases
were 5 % and 18 %, corresponding to the rotation speed
increase from 100 min–1 to 200 min–1, and from 200 min–1
to 400 min–1, respectively (corresponding to ARO results
for individual tests A-1, A-2, and A-3, under column
Di-1). For the same test group (A) conducted with
dielectric fluid mixed with 4 g/L carbon powder (column
Di-2), the increases were 3 % and 6 %, respectively
(results for individual tests A-1, A-2, and A-3, under
column Di-2) and for tests conducted with 8 g/L carbon
powder (column Di-3), the increases were 9 % and 6 %
(results for individual tests A-1, A-2, and A-3, under
column Di-3), respectively. Similar increases in ARO
values at all electrode rotation speeds were observed for
test groups B and C, corresponding to dielectric fluid
spray pressures of 40 bar and 80 bar, respectively.
In addition to the electrode rotation speed, the study
also looked at the effects of dielectric fluid spray pres-
sure on ARO values. It was observed that ARO values
increased as dielectric fluid spray pressure was in-
creased. For tests conducted with the electrode rotation
speed kept constant at 100 min–1 and using dielectric
fluid of pure distilled water (individual tests A-1, B-1
and C-1 under column Di-0), increasing the spray
pressure from 20 bar to 40 bar resulted in the ARO value
increasing by 26 %; increasing the spray pressure from
V. YÝLMAZ: INVESTIGATION OF HOLE PROFILES IN DEEP MICRO-HOLE DRILLING OF AISI 420 ...
672 Materiali in tehnologije / Materials and technology 50 (2016) 5, 667–675
Figure 7: Debris in the spark gap region18
Slika 7: Drobci v podro~ju re`e18
Figure 9: Dielectric pressure vs. ARO (constant rotation of 100 min–1)
Slika 9: Odvisnost med dielektri~nim tlakom in ARO (konstantna
hitrost rotacije 100 min–1)
Figure 8: Electrode rotation speed vs. ARO (constant spray pressure at
20 bar)
Slika 8: Odvisnost hitrosti vrtenja elektrode od ARO (pri konstantnem
tlaku curka 20 bar)
40 bar to 80 bar resulted in the ARO value increasing by
another 23 %. When the same tests were repeated for
dielectric fluid concentration with 2 g/L carbon, the
increases were 15 % and 21 % (individual tests A-1, B-1
and C-1 under column Di-1); for the dielectric fluid
concentration of 4 g/L carbon, the results were 6 % and
11 % (same individual tests under column Di-2); and for
dielectric fluid concentration of 8 g/L carbon, the results
were 14 % and 10 % (same individual tests under
column Di-3).
It is argued that better flushing of the machining area
with increased circulation speed of the dielectric fluid
and a more effective spark discharge, explain the in-
crease in ARO values as electrode rotation speed and
dielectric fluid spray pressures are increased. With faster
flushing of the spark gap, the molten debris is better
removed from the region, which leads to uninterrupted
spark discharges.1,2,7,16,32 With uninterrupted spark dis-
charges, continuous arcing between the electrode lateral
surfaces and cavity walls increases ARO values (depicted
in Figure 10).
In EDM operations, debris inside the cavity may
impede spark discharges, leading to a nonuniform hole
profile in terms of diameter; large scale craters may also
be formed.1-3,33,34 In this study, fluctuations in the hole
profile were prevented by the use of electrode rotation
and dielectric fluid spray through the electrode itself.
Through the selection of machining parameters, the
creation of large depressions in the cavity were pre-
vented; however, a certain increase in ARO values could
not be avoided. This phenomena is related to electrode
insulation, a current research topic with recent citations
in literature, which requires further development, as
research has demonstrated that insulating the lateral
surfaces of the electrode helps to prevent increases in the
ARO values.
Another factor that affected ARO values was the use
of the carbon powder. In the A-1 test with dielectric fluid
spray pressure of 20 bar and electrode rotation speed of
100 min–1, conducted using distilled water with carbon
powder at 2 g/L concentration (A-1 under Di-1), ARO
values were observed to increase by 48 % compared to
the same test conducted with pure distilled water (A-1
under Di-0). For the same test conducted using distilled
water with carbon powder at 4 g/L concentration (A-1
under Di-2), ARO values were observed to increase by
83 % compared to the same test conducted with pure
distilled water (A-1 under Di-0).When using distilled
water with carbon powder at 8 g/L concentration (A-1
under Di-3), ARO values were observed to increase by
93 % compared to the same test conducted with pure
distilled water (A-1 under Di-0). In test A-2 with elec-
trode rotation speed of 200 min–1, the observed increases
in ARO values were 35 %, 65 % and 83 %, respectively,
for Di-1, Di-2, and Di-3, compared to the ARO value for
Di-0. In test A-3 with electrode rotation speed of 400
min–1, the observed increases in ARO values were 51 %,
65 % and 82 %, respectively, for Di-1, Di-2, and Di-3,
compared to the Ra value for Di-0.
In the case of dielectric fluid spray pressure of 40 bar,
the observed increases in ARO values were 34 %, 55 %
and 74 % for electrode rotation speed of 100 min–1 (test
B-1), 32 %, 46 % and 66 % for electrode rotation speed
of 200 min–1 (test B-2), and 30 %, 47 % and 63 % for
electrode rotation speed of 400 min–1 (test B-3), respec-
tively, for Di-1, Di-2, and Di-3, when compared to the
ARO value for Di-0. In the case of dielectric fluid spray
pressure of 80 bars, the observed increases in ARO
values were 32 %, 40 % and 55 % for electrode rotation
speed of 100 min–1 (test C-1), 22 %, 31 % and 46 % for
electrode rotation speed of 200 min–1 (test C-2), and
24 %, 34 % and 47 % for electrode rotation speed of 400
min–1 (test C-3), respectively, for Di-1, Di-2, and Di-3,
when compared to the ARO value for Di-0.
At higher dielectric fluid spray pressure, it is ob-
served that the increase in ARO values slows down as the
carbon powder concentration is increased. In the study,
the lowest increase observed for ARO values as a result
of an increase in carbon powder concentration was for
the case of the spray pressure of 80 bar. Indeed, the
increases of 22 %, 31 % and 46 % for test C-2 are the
lowest figures for the study in terms of ARO increases
with carbon powder concentrations. The interpretation
for this surprising finding is that, as spray pressure is
increased, the spark gap is flushed clean to a greater
extent, allowing faster removeal of debris.20,26,32–34 With
faster flushing, a continuous supply of clean dielectric
fluid is provided, preventing the debris from generating
unwanted spark discharges and an excessive rise in the
ARO values. It also results in a decreased machining
duration. An increase in machining duration would have
increased the hole diameter as well as the accompanying
ARO values.
ARO values that increase in tandem with increases in
carbon powder concentrations bring to the fore a funda-
mental problem in EDM hole drilling. Expulsion of
carbon powders through the cavity as a result of
dielectric flushing leads to spark generation between
V. YÝLMAZ: INVESTIGATION OF HOLE PROFILES IN DEEP MICRO-HOLE DRILLING OF AISI 420 ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 667–675 673
Figure 10: Increase in ARO values
Slika 10: Pove~anje vrednosti ARO
electrode lateral surfaces and cavity walls. This has
caused ARO values to increase in an accelerated fashion.
The most significant result achieved in this study has
been the low increases in ARO values observed at high
dielectric fluid spray pressures for increasing powder
concentrations, as decreased machining duration has
positively influenced the ARO values. However, when
the test results are considered in their entirety, it is clear
that increased carbon powder concentrations in turn
increase ARO values. While this outcome may in fact be
exploited in certain applications, it is evident that carbon
powder concentrations have a significant effect on hole
profiles.
4 CONCLUSION
Using EDM, micro-holes with diameter of 0.5 mm
and depth of 20 mm were drilled into an AISI 420 steel
workpiece using four dielectric fluid concentrations,
three electrode rotation speeds, and three dielectric fluid
spray pressure settings. Effects of the machining para-
meters on the ARO values for hole profile and surface
roughness (Ra) have been investigated. The results of the
tests are presented below.
Significant reductions in Ra values were observed by
mixing carbon powder with the dielectric fluid. As the
carbon powder concentration was increased, Ra values
decreased. In the study, Ra values were observed to
decrease as electrode rotation speed was increased. As
dielectric fluid spray pressure was increased, Ra values
were again observed to decrease.
The study also investigated the hole profile (i.e. ARO
values); as carbon powder concentrations, electrode rota-
tion speeds and the dielectric fluid spray pressure were
increased, ARO values were observed to increase. The
findings were interpreted to be the result of a continuous
spark discharge being enabled. It was determined that Ra
and ARO values may be controlled by configuring the
machining parameters and using an optimal mix of
carbon powder with the dielectric fluid.
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Materiali in tehnologije / Materials and technology 50 (2016) 5, 667–675 675

W. OZGOWICZ et al.: THE PHENOMENON OF REDUCED PLASTICITY IN LOW-ALLOYED COPPER
677–682
THE PHENOMENON OF REDUCED PLASTICITY IN
LOW-ALLOYED COPPER
POJAV ZMANJ[ANJA PLASTI^NOSTI MALO LEGIRANEGA
BAKRA
Wojciech Ozgowicz1, El¿bieta Kalinowska-Ozgowicz2, Barbara Grzegorczyk1,
Klaudiusz Lenik2
1Silesian University of Technology, Mechanical Engineering Faculty, Institute of Engineering Materials and Biomaterials,
Konarskiego Str. 18A, 44-100 Gliwice, Poland
2Lublin University of Technology, Fundamentals of Technology, Nadbystrzycka Str. 38, 20-618 Lublin, Poland
kalinowska-ozgowicz@tlen.pl
Prejem rokopisa – received: 2015-05-19; sprejem za objavo – accepted for publication: 2015-10-12
doi:10.17222/mit.2015.101
This paper presents the results of investigations that allow us to determine the influence of the temperature of plastic
deformation in the range from 20 °C to 800 °C during static tensile tests on the mechanical properties and structure of low-alloy
copper alloys of the type CuCo2 and CuCo2B, completed by measurements of the microhardness and observations of the
structure in a light microscope, and also of fractures in a scanning electron microscope. Based on the results of these
investigations the temperature range for the occurrence of the reduced plasticity of the alloys CuCo2 and CuCo2B could be
determined.
Keywords: low-alloy copper, plastic deformation, structure, mechanical properties, brittleness
^lanek predstavlja rezultate preiskav, ki omogo~ajo opredelitev vpliva temperature na plasti~no deformacijo v obmo~ju od
20 °C do 800 °C s stati~nimi nateznimi preizkusi na mehanske lastnosti in strukturo malo legiranih bakrovih zlitin, vrste CuCo2
in CuCo2B, izvedenih z merjenjem mikrotrdote ter opazovanjem mikrostrukture v svetlobnem mikroskopu in prelomov v
vrsti~nem elektronskem mikroskopu. Na osnovi rezultatov teh preiskav je bilo mogo~e opredeliti temperaturno podro~je pojava
zmanj{anja plasti~nosti zlitin vrste CuCo2 in CuCo2B.
Klju~ne besede: malo legirani baker, plasti~na deformacija, struktura, mehanske lastnosti, krhkost
1 INTRODUCTION
Low-alloy copper is applied in various ways. How-
ever, most of it is applied in electrical engineering and
electronics. It is also used in the production of welding
electrodes, elements of bearings, non-sparking tools and
chemical apparatus.1–3 High-temperature brittleness
results in a reduced plasticity at the given temperature of
deformation, called the temperature of minimum plasti-
city (TMP).4–6 The reason for this phenomenon concern-
ing the brittleness of copper alloys has not been fully
explained yet; it depends on many factors, mainly on the
chemical composition, the structure of the alloy and the
parameters of the deformation.7–12
The purpose of the present investigations was to
determine the influence of the temperature of deforma-
tion on the mechanical properties, the structure, and par-
ticularly the range of temperature for the reduced plas-
ticity of low-alloy copper, containing cobalt and boron of
the type CuCo2 and CuCo2B.
2 MATERIALS AND METHODS
The investigations concerned low-alloy copper type
CuCo2 and CuCo2B smelted in the laboratory in a
crucible induction furnace with a frequency from 500 Hz
to 4000 Hz and the mass of the charge up to 100 kg. In
the course of smelting to liquid the CuCo2B alloy, boron
was added in an amount of 0.005 %. The ready melts
were passed to a graphite gate with a diameter of 30 mm.
After cooling the obtained ingots, re-forged to rods,
15 mm in diameter, on a pneumatic forging hammer, the
weight of its ram amounting to 200 t. For the chemical
compositions of the investigated alloys CuCo2 and
CuCo2B (Table 1).
Table 1: Chemical composition of the investigation alloys
Tabela 1: Kemijska sestava preiskovanih zlitin
Alloy
type
Mass contents in mass fractions, (w/%)
Cu + Ag Co Si Fe Ni P B
CuCo2 96.71 2.76 0.29 0.16 0.01 0.05 –
CuCo2B 96.88 2.86 0.16 0.01 0.01 0.07 0.005
After forging the rods were supersaturated at 900 °C
and cooled in water. The temperature during this proce-
dure was determined based on an analysis of a binary
system of the phase equilibrium of copper with co-
balt.11,12 The temperature of supersaturation was assumed
to be 100 °C higher than the boundary temperature of the
solubility of Co on Cu concerning the tested alloys. The
Materiali in tehnologije / Materials and technology 50 (2016) 5, 677–682 677
UDK 669.3:67.017:620.178.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)677(2016)
operation of supersaturation was carried out in an elec-
tric chamber furnace equipped with a controller ensuring
measurements of the temperature with an accuracy of
±2 °C. After their supersaturation the rods were cut into
segments, from which samples were used for testing the
mechanical properties, applying a threaded grip.
The chemical compositions of the alloys CuCo2 and
CuCo2B were tested on monolithic samples in the shape
of disks with a thickness of about 5 mm and a diameter
of 30 mm, cut out from the ingots.
The mechanical properties of the alloys CuCo2 and
CuCo2B were tested on an Instron 1115 universal testing
machine provided with a high-temperature resistance
furnace, including a microprocessing system controlling
the temperature. The procedure of heating was per-
formed in a protective atmosphere containing 95 % nitro-
gen and 5 % hydrogen. Static tensile tests were accom-
plished in the temperature range 20 °C to 800 °C at a
tensile rate of 20 mm/min, corresponding to the strain
rate  = 1.28 10–3s–1. Based on the data on the curves of
tension for the investigated alloys, the tensile strength
(Rm) was determined, and the elongation (A) and the
reduction of the area of the sample (Z) were calculated
based on the geometrical features of the sample previous
to and after the rupture. The result of the tests is the
arithmetic mean of the three measurements.
Metallographic investigations were carried out on
longitudinal microsections of the alloy CuCo2 and
CuCo2B after supersaturation and hot tensile tests. The
samples were immersed in self-hardening resin, and then
mechanically polished. In order to reveal their structure
the samples were etched in a reagent containing 5 g
FeCl3, 10 cm3 HCl and 100 cm3 C2H5OH. Metallographic
observations were performed using an Olympus GX71
(Japan) light microscope with a magnifying power of up
to 1000 times. The size of the grains was measured by
applying the method of sections.
Fractographic tests of the fractions after decohesion
in the tensile test were produced by means of a DSM940
scanning electron microscope by the firm Opton, accom-
plished at an accelerating voltage of 20 kV and magnify-
ing power of up to 3000 times. The precipitation ob-
served on the fractures was investigated by means of an
EDAX X-ray microanalyzer. Prior to the fractographic
test, the sample was ultra-sound cleaned in ethyl alcohol
for 3 min.
The microhardness was measured on a Vickers scle-
rometer, applying a load of 50 g. These measurements
were carried out on metallographic microsections of the
alloys CuCo2 and CuCo2B after tension at a temperature
of 20 °C to 800 °C.
3 RESULTS AND DISCUSSION
The results of the analysis of the chemical compo-
sitions of the investigated alloys have been gathered in
Table 1. The analysis revealed that in these alloys there
is a presence of cobalt and boron, as well as admixtures
of silicon, iron, nickel and phosphorus. These elements
affect mainly the electrical conductivity of copper,
reducing it. Moreover, cobalt and iron increase the hard-
ness of these alloys, and phosphorus is a de-oxidant,
increasing their viscosity.
The results of the static tensile tests allowed us to
determine the effect of temperature on the strength and
plastic properties of the alloys CuCo2 and CuCo2B, and
thus also to assess the range of temperature at which the
plasticity of the investigated alloys decreases due to the
dependence of elongation, contraction and strength on
the temperature of deformation (Figures 1 to 3). Ana-
lyzing the dependence of the reduction of the area of the
sample on the temperature of deformation of these
alloys, it has been found to be more or less the same. In
both these alloys the range of temperatures at which such
a contraction attains its minimum value is quite evident.
The diagram of the dependence of elongation on the
temperature of deformation of the alloy CuCo2 is cha-
racterized by a varying shape (Figure 1). At 20 °C the
elongation of the alloy amounts to 34 %. An increase in
the temperature of deformation is accompanied by a de-
crease in the value of A, reaching its minimum of 4.7 %
at the temperature of 600 °C. A further rise in the
temperature of deformation to 800 °C involves an in-
crease of the elongation to 22 %. At 20 °C the elongation
of the alloy CuCo2B amounts to 45.7 %. If the tempe-
rature of deformation rises from 20 °C to 550 °C, the
elongation decreases, reaching its minimum of 10 % at
550 °C. A further rise of the temperature of deformation
results in an elongation amounting to 53 % at 800 °C.
The alloy CuCo2B is characterized by a much larger
elongation in the range of the temperature of deforma-
tion from 700 °C to 800 °C than the alloy CuCo2.
Analyzing the dependence of the course of contraction
on the temperature of deformation of the alloys CuCo2
and CuCo2B, it has been found to be similar (Figure 2).
In the case of both these alloys the range of temperature
W. OZGOWICZ et al.: THE PHENOMENON OF REDUCED PLASTICITY IN LOW-ALLOYED COPPER
678 Materiali in tehnologije / Materials and technology 50 (2016) 5, 677–682
Figure 1: The influence of the temperature of plastic deformation in
the tensile test on the elongation (A) of the alloys CuCo2 and CuCo2B
Slika 1: Vpliv temperature plasti~ne deformacije pri nateznem preiz-
kusu na raztezek (A) zlitin CuCo2 in CuCo2B
characterized by a minimum contraction is quite distinct.
The contraction of the alloy CuCo2, deformed in the
range of temperature from 20 °C to 600 °C, decreases,
reaching its minimum at 600 °C (Z = 5.5 %). At a tem-
perature of 600 °C to 800 °C the contraction increases up
to a value of 30.6 %. At the temperature of deformation
20 °C the contraction of the alloy CuCo2 amounts to
71 % (Figure 2).
On the curve of the dependence of the contraction on
the temperature of deformation of the alloy CuCo2B
there occurs a local minimum (Figure 2). In the range of
the temperature of deformation 20 °C to 550 °C the
value of the contraction decreases from 85.7 % at 20 °C
and attains its minimum Z = 10.9 % at 550 °C. A further
rise of the temperature of deformation (to 800 °C) leads
to an increase of the contraction to a value of 84.2 %.
Comparing the diagrams of the dependence of elon-
gation and contraction on the temperature of deformation
concerning the alloys CuCo2 and CuCo2B, we find that
in both cases there exists a range of temperature in which
these alloys indicate a minimum of the plastic properties,
characteristic for the phenomenon of brittleness (Figures
1 and 2). The alloy with the addition of boron is charac-
terized by brittleness in the range of lower temperatures
than the alloy without boron. The elongation and con-
traction of the alloy CuCo2B exceed those of the alloy
CuCo2 in the entire range of the investigated tempe-
rature. The alloy CuCo2 displays a minimum plasticity
in the range of temperature from 500 °C to 700 °C, and
the alloy CuCo2B at a temperature of 450 °C to 600 °C.
Subjected to a static tensile test in the range of tem-
perature from 20 °C to 800 °C, the investigated alloys
display similar values of tensile strength (Figure 3). The
curve in the diagram of the dependence of the tensile
strength on the temperature of deformation concerning
the alloys CuCo2 and CuCo2B is a decreasing function.
The tensile strength of the alloy CuCo2, deformed at a
temperature of 20 °C amounts to 237 MPa and drops to
38 MPa at 800 °C, where as in the case of the alloy
CuCo2B it amounts, respectively, to 232 MPa and
34 MPa.
The results of the metallographic investigations
allowed us to determine the influence of the temperature
of deformation on the structure of the CuCo2 and
CuCo2B in the range from 20 °C to 800 °C (Figures 4 to
9). After a hot tensile test the alloys CuCo2 and CuCo2B
have a variated structure in the zone of rupture and the
central zone of the sample, with sliding bands. In the
central part of the samples the grains have been found
with a hardness of 71–93 HV and twins with straight and
curve-linear boundaries. In the zone of rupture the
structure of the alloy CuCo2, stretched at a temperature
of 200 °C, is characterized by the occurrence of micro-
cracks at the boundaries of elongated grains of the phase
 (Figure 4), and the central part of the sample by axial
grains  with twins (Figure 5). The alloy CuCo2B has a
similar structure in the zone of rupture. The structures of
Materiali in tehnologije / Materials and technology 50 (2016) 5, 677–682 679
W. OZGOWICZ et al.: THE PHENOMENON OF REDUCED PLASTICITY IN LOW-ALLOYED COPPER
Figure 4: Elongated grains of the phase  with a micro-crack in the
structure of the alloy CuCo2 after stretching at temperature of 200 °C
(zone of rupture)
Slika 4: Razpotegnjena zrna  faze z mikrorazpokami v strukturi zli-
tine CuCo2, po nateznem preizkusu na temperaturi 200 °C (podro~je
preloma)
Figure 3: The influence of the temperature of plastic deformation in
the tensile test on the strength (Rm) of the alloys CuCo2 and CuCo2B
Slika 3: Vpliv temperature plasti~ne deformacije pri nateznem preiz-
kusu na trdnost (Rm) zlitin CuCo2 in CuCo2B
Figure 2: The influence of the temperature of plastic deformation in
the tensile test on the reduction of area (Z) of the alloys CuCo2 and
CuCo2B
Slika 2: Vpliv temperature plasti~ne deformacije pri nateznem
preizkusu na kontrakcijo (Z) zlitin CuCo2 in CuCo2B
alloys stretched at elevated temperatures display a larger
amount of microcracks occurring at the boundaries of the
grains and at the contact of three grains and twin boun-
daries in the zone of rupture. In the central part of the
sample a heterogeous structure was detected consisting
of a diversified size of the grains (20 μm to 60 μm)
depending on the temperature of the deformation and
due to the process of recrystallization.
In alloys stretched at the temperature of minimum
plasticity (550 °C) the structure in the zone of rupture is
characterized by numerous micro-cracks. In the structure
of the alloy CuCo2B, stretched at the temperature 550 °C,
the zone of rupture contains axial recrystallized grains of
the phase , 40 μm in diameter, and also many micro-
cracks (Figure 6). Also in the central part of the sample
there are micro-cracks at the boundary of the phase 
(Figure 7). The structure of this part of the sample con-
tains grains of the phase  with many twins with
straight-lined boundaries as well as stepped boundaries,
testifying to the advanced recrystallization of the alloy.
In the central part the sample of the alloy CuCo2B there
are grains of the phase  with micro-cracks and precipi-
tations. After their deformation at 600 °C the investi-
gated alloys display the structure of grains of the phase
, varying in their size, with twins and micro-cracks
(Figures 8 and 9). The structure of the alloy CuCo2B,
elongated at a temperature of 800 °C, displays numerous
micro-cracks both in the zone of rupture and in the
central part of the sample. In the structure of the central
part of the sample the micro-cracks occurred in the front
recrystallization due to the presence of large grains of the
W. OZGOWICZ et al.: THE PHENOMENON OF REDUCED PLASTICITY IN LOW-ALLOYED COPPER
680 Materiali in tehnologije / Materials and technology 50 (2016) 5, 677–682
Figure 8: Differentiated grains of the phase  with the twins and
micro-cracks in the structure of the alloy CuCo2B after stretching at
the temperature 600 °C (central zone)
Slika 8: Diferencirana zrna  faze z dvoj~ki in mikrorazpokami v
strukturi zlitine CuCo2B po nateznem preizkusu na temperaturi
600 °C (sredina vzorca)
Figure 7: Micro-cracks at the boundaries of the phase  in the
structure of the alloy CuCo2B after stretching at a temperature of
550 °C (central zone)
Slika 7: Mikrorazpoke na mejah  faze v strukturi zlitine CuCo2B po
nateznem preizkusu na temperaturi 550 °C (sredina vzorca)
Figure 6: Recrystallized grains of the phase  and numerous cracks in
the structure of the alloy CuCo2B after stretching at a temperature of
550 °C (zone of rupture)
Slika 6: Rekristalizirana zrna  faze in {tevilne razpoke v strukturi
zlitine CuCo2B po nateznem preizkusu na temperaturi 550 °C (pod-
ro~je preloma)
Figure 5: Elongated grains of the phase  with twins and bands of
deformation in the structure of the alloy CuCo2 after stretching at a
temperature of 200 °C (central zone)
Slika 5: Razpotegnjena zrna  faze z dvoj~ki in deformacijskimi pa-
sovi v strukturi zlitine CuCo2 po nateznem preizkusu na temperaturi
200 °C (sredina vzorca)
phase (about 100 μm) with a hardness of about 60 HV
and a revealed substructure (Figure 9). The size of the
grains in the phase  of the structure of the alloy
CuCo2B results from the way of recrystallization in the
course of and after the plastic deformation during the
tensile test.
The results of fractographic investigations allowed us
to determine the influence of the temperature of defor-
mation on the character of the fractures of the alloys
CuCo2 and CuCo2B after decohesion in the tensile test
in the range of temperature from 20 °C to 800 °C.
The fracture of the alloy CuCo2 and CuCo2B after
the decohesion in the tensile test indicates a diversified
character depending on the temperature of tension. At
the temperature of deformation amounting to 200 °C, the
investigated alloys are characterized by a transcrystalline
ductile fracture with numerous craters differing in the
W. OZGOWICZ et al.: THE PHENOMENON OF REDUCED PLASTICITY IN LOW-ALLOYED COPPER
Materiali in tehnologije / Materials and technology 50 (2016) 5, 677–682 681
Figure 13: Result of the quantitative microanalysis of the chemical
composition of a precipitate in the alloy CuCo2 after a tensile test at
550 °C
Slika 13: Rezultati kvantitativne mikroanalize kemijske sestave
izlo~ka v zlitini CuCo2 po nateznem preizkusu na 550 °C
Figure 10: Transcrystalline ductile fracture in the alloy CuCo2 after a
tensile test at 200 °C
Slika 10: Transkristalni `ilav prelom zlitine CuCo2 po nateznem
preizkusu na 200 °C
Figure 9: Coarse grains of the phase  with sub-grains in the structure
of the alloy CuCo2B after stretching at a temperature of 800 °C
(central zone)
Slika 9: Velika zrna  faze s podzrni v strukturi zlitine CuCo2B po
nateznem preizkusu na temperaturi 800 °C (sredina vzorca)
Figure 12: Intercrystalline brittle fracture in the alloy CuCo2 after a
tensile test at 600 °C
Slika 12: Interkristalni krhki prelom zlitine CuCo2 po nateznem
preizkusu na 600 °C
Figure 11: Intercrystalline brittle fracture in the alloy CuCo2B after a
tensile test at 550 °C
Slika 11: Interkristalni krhki prelom zlitine CuCo2B, po nateznem
preizkusu na 550 °C
diameters and precipitations in the bottom (Figure 10).
The lateral planes of the craters are considerably corru-
gated.
At the temperature of deformation amounting to
550 °C and 600 °C, these alloys display brittle inter-
crystalline fractures with many micro-cracks and precipi-
tations (Figures 11 and 12).
The planes of the cracks indicate the effects of plastic
deformation. At the bottom of the crater on the fracture
of the alloy CuCo2 precipitations were found, the che-
mical composition of which was determined by means of
an X-ray analysis (EDAX) and proved to contain
96.55 % copper and 3.55 % cobalt (Figure 13). In the
alloy CuCo2 deformed at 600 °C; an inter-crystalline
brittle fracture was detected with micro-cracks at the
boundaries (Figure 12), whereas in samples deformed at
800 °C a fracture mixed with cracks on the grain boun-
daries was observed.
4 CONCLUSIONS
The performed investigations and analyses of the
obtained results allow us to draw the following conclu-
sions:
1. The low-alloy copper type CuCo2 reaches its
minimum plasticity in the tensile test at a temperature
of deformation from 500 °C to 700 °C, whereas in
the case of the alloy CuCo2B the minimum value is
attained at a temperature of 450 °C to 600 °C.
2. An increase in the temperature of plastic deformation
from 20 °C to 800 °C involves a decrease in the
tensile strength of the alloy CuCo2 from about 240
MPa to about 40 MPa, and that of the alloy CuCo2B
from about 230 MPa to about 25 MPa.
3. The temperature of the minimum plasticity (TMP) of
the alloy CuCo2B from 20 °C to 800 °C is about 50
°C lower than the TMP of the alloy CuCo2. With a
microaddition of boron the alloy is also more plastic
(A and Z by about 5 %) in the range of TMP if
compared with the alloy CuCo2.
4. The structures of the investigated alloys of copper in
the range TMP are characterized by homogeneous
grains in the solution , about 40 μm in size, with
numerous micro-cracks at the grain boundaries.
5. The investigated plastically hot-deformed low-alloys
beyond the region TMP have a typical structure of
the solution  with a differing degree of deformation
or dynamic or static recrystallization and a ductile
fracture.
6. Low-alloy copper, in the range of TMP characterized
by minimum plastic properties (A and Z about 5-10
%), displays after stretching a brittle intercrystalline
fracture.
7. A micro-addition of boron involves increased plastic
properties of the alloy CuCo2B in the entire range of
the temperature of plastic deformation and changes
the character of the fracture in the temperature
interval from 550 °C to 800 °C.
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Karel Dvoøák, Iveta Hájková
Brno University of Technology, Faculty of Civil Engineering, Veveøí 331/95, 602 00 Brno, Czech Republic
dvorak.k@fce.vutbr.cz
Prejem rokopisa – received: 2015-06-23; sprejem za objavo – accepted for publication: 2015-09-21
doi:10.17222/mit.2015.127
The aim of this work was to observe the impact of the milling technique employed by the DESI 11 disintegrator on the
properties of fly ash. This type of mill is a high-speed pin mill with two counter rotors. The device selected for study allows the
use of rotors with different working tools. In this case, two types of rotors were selected, identified as BR-AR and OR rotors.
The OR rotor has cylindrical teeth, while the teeth on the AR-BR rotor can be described as rhomboidal. The fly ash was ground
by 1, 2, 3, 5, and 10 passes through the mill. The Blaine specific surface area, particle size and particle size distribution were
measured for each sample. The pozzolanic activity was also determined via a modified Chapelle test, and its influence on the
shape of the grains was assessed by SEM. The results were compared together and with the original fly ash. There was a steep
increase in the specific surface area and pozzolanic activity after the first pass through the mill. The second pass of ash through
the mill did not increase the specific surface area due to strong aggregation, which gradually changes into agglomeration.
However, the pozzolanic activity still increased during the aggregation phase. This phenomenon was clearly observable for both
types of selected rotors. Based on the results, we can say that in the case of fly ash, a high-speed disintegrator can be a
promising means of improving its properties via grinding and mechanical activation.
Keywords: high speed grinding, fly ash, pozzolanic activity
Namen dela je opazovati posledice uporabljene tehnike mletja, v DESI 11 mlinu, na lastnosti lete~ega pepela. To je mlin z
veliko hitrostjo, z dvema nasprotnima rotorjema. Naprava, izbrana za {tudij, omogo~a uporabo rotorjev z razli~nimi orodji. V
tem primeru sta bila izbrana dva razli~na rotorja, ozna~ena kot AR-BR in OR rotor. OR rotor ima cilindri~en zob, medtem ko je
zob na AR-BR rotorju ozna~en kot romboidalen. Lete~i pepel je bil zmlet z 1, 2, 3, 5 in 10 prehodi skozi mlin. Za vsak vzorec
so bile izmerjene Blainova specifi~na velikost povr{ine ter velikost in razporeditev delcev. Pucolanska aktivnost je bila dolo~ena
z modificiranim Chapelle testom. Vpliv na obliko zrn je bil dolo~en s SEM. Rezultati so bili primerjani z originalnim lete~im
pepelom. Po prvem prehodu skozi mlin se je skokovito pove~ala specifi~na povr{ina in pucolanska aktivnost. Drugi prehod
pepela skozi mlin ni pove~al specifi~ne povr{ine zaradi mo~ne agregacije, ki se je postopoma spremenila v aglomeracijo.
Vseeno pa je pucolanska aktivnost {e nara{~ala med fazo agregacije. Ta pojav je bil jasno opa`en pri obeh vrstah rotorjev. Na
osnovi teh rezultatov lahko re~emo, da je v primeru lete~ega pepela dezintegrator z veliko hitrostjo obetajo~e sredstvo, da z
mletjem in mehansko aktivacijo izbolj{amo njegove lastnosti.
Klju~ne besede: mletje z veliko hitrostjo, lete~i pepel, pucolanska aktivnost
1 INTRODUCTION
One of the trends in the area of milling that have been
intensively examined recently is high-energy milling
(HEM). With HEM there are certain phenomena that
have not been observed in conventional grinding. These
phenomena can be summarized by the term mechano-
chemical activation.1 The idea of mechanochemical
activation is that an increase occurs in the value of the
internal energy, consequently increasing the enthalpy and
the formation of so-called active surfaces on the newly
formed grains. To increase the level of enthalpy, it is not
only necessary to supply a large amount of mechanical
energy, but it is also important how that energy is trans-
mitted to the material. Only if the material is capable of
absorbing the energy can the internal structure of the
substance be destabilized and new active surfaces be
developed.2 The effects of mechanical activation have
been described for model materials such as dolomite3 or
clay minerals,4–6 or, for example, for silica.7 In the field
of grinding and mechanical activation, research is fo-
cused on monitoring changes in the crystal structure and
amorphization, 6,8 changes in the granularity and the
aggregation and agglomeration of particles, 9,10 and
changes in surface properties, especially the specific
surface area and the zeta potential.4,8 One type of HEM is
high-speed grinding (HSG). HSG involves supplying
large amounts of energy using very short and intense
power pulses. The amount of energy that is effectively
transferred to the material is higher in the case of HSG
than with conventional grinding in mills with the same
power input. One of the types of mills suitable for HSG
is a high-speed pin mill with two counter-rotating rotors,
known as a disintegrator.11 This type of mill is particu-
larly suitable for the grinding and activation of fine
powder materials.12 The material is refined by very
intensive changes in the mechanical stress, which take
place at a very high frequency. Another advantage is the
variety of working tools that can be employed to affect
the grinding process.13 The disadvantage of this type of
milling in the case of silicates is the significant electro-
Materiali in tehnologije / Materials and technology 50 (2016) 5, 683–687 683
UDK 621.926:621.763:67.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)683(2016)
static charging of the particles and their easy and quick
aggregation.1 One of the important silicate materials that
are widely used in the construction industry is ash, va-
rious kinds of which exist. A number of authors describe
the utilization and effects of pin mills on the physical
and mechanical properties of both fluidized bed ashes14,15
and ashes based on SiO2.16 In the case of fly ash, during
traditional milling and even during HEM there is not
only an increase in the specific surface area, but also an
improvement in the reactivity. Different authors use a
wide spectrum of methods for the assessment of reacti-
vity, though it is typically done by FTIR and calorimetric
methods during the hydration of cement pastes contain-
ing fly ash.17,18 Increases in the reactivity are commonly
achieved via the refinement of large porous particles of
ash, i.e., plerospheres, which is relatively easy to do. The
refinement of cenospheres is more difficult, and requires
more grinding work, but leads to a further increase in the
pozzolanic activity of the fly ash.19 The particles remain
in the chamber of the high-speed disintegrator for only a
few seconds, which is because of the principle by which
the pin mill operates. Repeating the passing of the mate-
rial through the mill is the only effective way of
extending the milling time. However, in the case of fly
ash it entails the risk of grain agglomeration. The aim of
this work was to describe and assess the effect of a HSG
disintegrator’s milling technique on the fly ash’s
properties, especially its pozzolanic properties, grain
size, grain morphology and specific surface area.
2 MATERIALS AND METHODS
TEKO fly ash was used for monitoring the impact of
the milling technique in question on the material pro-
perties. The fly ash was analyzed before the milling
began, with its chemical composition being determined
by traditional chemical analyses. The density was deter-
mined using a Micromeritics AccuPyc II 1340 automatic
pycnometer. For the measurement of the Blaine specific
surface area, a PC-Blaine-Star automatic device was
used with a measurement cell capacity of 7.95 cm3. The
determination was performed three times to eliminate
errors, and the resultant value was the average of these
determinations. Milling of the samples was carried out in
a DESI 11 disintegrator, which is a high-speed pin mill
with two counter-rotating rotors. The total installed
power of this mill is 4.1 kW. The rotors’ rotation fre-
quency is up to 200 Hz, and the maximum speed of
impact is 240 ms–1. The material is dispensed by a conti-
nuous feeder and enters the grinding chamber through
the middle of the left rotor. To assess the effect of the
selected milling technique, two rotor types were used for
comparative purposes, i.e., types AR-BR and OR. The
rotors were designed and manufactured by FF Service
Ltd. Both types have two rows of teeth on the left-hand
rotor and three rows of teeth on the right-hand rotor. The
OR rotor has cylindrical teeth, while the shape of the
teeth on the AR-BR rotor can be described as rhom-
boidal. With this kind of rotor, the tilt direction of the
leading edge against the flow of material on the left-hand
rotor is always opposite to that of the right-hand rotor. A
continuous feeder was used for dosing the fly ash into
the mill. The ash was always fed in at a dose of 0.5 kg
and milled at the maximum Hz. The time required for all
the material to pass through the mill was 180 s. The sam-
ples for both rotor variants were prepared for 1, 2, 3, 5
and 10 passes through the mill. Grinding was performed
under standard laboratory conditions, at 22 °C and with a
relative humidity of 56 %. Between the steps of grinding,
the mill was cooled by an air stream so that the
temperature of the working chamber did not exceed
70 °C. The Blaine specific surface area was determined
for all the samples. All the fly ash samples, including the
control sample, were subjected to particle size distribu-
tion measurements by laser granulometry. This analysis
was performed on a Malvern Mastersizer 2000 with a
Hydro 2000 G wet dispergation unit; 2-isopropanol was
used as a dispersant. The effect of the milling technique
on the morphology of the grains was observed and
assessed by electron microscopy (SEM). A Tescan
MIRA 3 XMU SEM with a secondary electron detector
was used. The Chapelle test method20 was used to deter-
mine the pozzolanic activity This test was done on the
control sample of fly ash and the samples ground via
one, three, and ten passes. The method was adapted for
the needs of the experiment. The modified Chapelle test
consists of allowing pozzolan and freshly annealed CaO
to react together in an aquatic environment at 93 °C for
24 h. The reaction takes place in a tightly closed stain-
less-steel vessel and the suspension is stirred with an
electromagnetic stirrer. The result is expressed as the
amount of Ca(OH)2 bound in mg per 1 g of pozzolan.
The results were compared to each other and with those
from the control sample of fly ash.
3 RESULTS
The results for the chemical composition of the con-
trol TEKO fly ash are shown in Table 1. Only selected
oxides and loss on ignition are listed in the table.
Table 1: The chemical composition of the fly ash
Tabela 1: Kemijska sestava lete~ega pepela
Compo-
nents SiO2 Al2O3 Fe2O3 CaO SO3
Loss on
ign. Others
Content %
per mass 51.8 18.1 9.1 6.7 0.8 6 7.5
The specific surface area and density are then pre-
sented in Table 2.
Table 2: Density and specific surface area of the TEKO fly ash
Tabela 2: Gostota in specifi~na povr{ina TEKO lete~ega pepela
Density (kg/m3) 2423
Specific surface area (m2/kg) 494
K. DVOØÁK, I. HÁJKOVÁ: THE EFFECT OF HIGH-SPEED GRINDING TECHNOLOGY ON THE PROPERTIES OF FLY ASH
684 Materiali in tehnologije / Materials and technology 50 (2016) 5, 683–687
The chemical and physical properties are typical for
conventional siliceous ash.
The specific surface area values were determined
from the samples ground via the selected operating mode
for both types of rotors immediately after passing
through the mill. The specific surface area results for
both types of rotors are shown in Figure 1.
After the first pass of the samples through the mill, a
significant increase in the specific surface area is
apparent in both cases. However, after the second pass
through a significant decrease in the specific surface area
can be seen. The decrease in specific surface area in the
initial stages was greater in the case of the AR-BR
rotors. Each additional pass through the mill led to a
further reduction in the surface area. However, after ten
cycles, the sample milled by the OR rotor had a lower
specific surface area than the sample milled by the
AR-BR rotor.
The granulometry of the control fly ash and of all the
ground samples was analyzed by laser granulometry.
Selected laser granulometry results for both types of
rotor are shown in Figures 2a and 2b.
The dependence of the grain size d (0.1), d (0.5) and
d (0.9) on the number of passes of the material through
the mill is summarized in Table 3 in order to facilitate a
comparison of the impact of the milling technique on the
course of the grinding.
Table 3: The dependence of the grain size d (0.1), d (0.5) and d (0.9)
on the number of passes of the material through the mill
Tabela 3: Odvisnost velikosti zrn d (0,1), d (0,5) in d (0,9) od {tevila
prehodov materiala skozi mlin
Number of passes through the mill
d/μm 0 1 2 3 5 10
AR-BR
0.1 4.90 3.52 3.66 3.53 3.99 3.99
0.5 25.60 13.19 13.19 12.02 12.15 12.99
0.9 108.083 52.87 47.80 43.86 37.57 40.08
OR
0.1 4.90 3.68 3.47 3.39 4.27 4.85
0.5) 25.60 12.56 11.72 10.61 11.40 12.39
0.9 108.083 46.59 38.90 32.45 33.01 30.91
In both cases, after the first pass there was a signi-
ficant reduction in the proportion of coarse particles,
with each additional grinding causing a further decline in
the amount present. However, at the same time, there
was also a decrease in the proportion of ultra-fine parti-
K. DVOØÁK, I. HÁJKOVÁ: THE EFFECT OF HIGH-SPEED GRINDING TECHNOLOGY ON THE PROPERTIES OF FLY ASH
Materiali in tehnologije / Materials and technology 50 (2016) 5, 683–687 685
Figure 2: Particle size distribution: a) AR-BR and b) OR
Slika 2: Razporeditev velikosti delcev: a) AR-BR in b) OR
Figure 1: Specific surface area for all the samples
Slika 1: Specifi~na povr{ina pri vseh vzorcih
Figure 3: a) The morphology of the control fly ash, b) 1 pass, AR-BR
rotors, c) 1 pass, OR rotors, d) 10 pass, AR-BR rotors and e) 10 pass,
OR rotors
Slika 3: a) Morfologija kontrolnega lete~ega pepela, b) 1 prehod,
AR-BR rotorji, c) 1 prehod, OR rotorji, d) 10 prehod, AR-BR rotorji
in e) 10 prehod, OR rotorji
cles. The distribution curves after 10 passes were signi-
ficantly narrower. This phenomenon is more significant
in the case of the OR rotor. The results of the effect of
the rotors and the number of passes on the morphology
of the grains after one and ten passes of fly ash through
the mill are illustrated in Figures 3a to 3e.
The SEM analysis shows that a clearly visible refin-
ing process is occurring, along with the subsequent
agglomeration of particles. This effect is more distinctive
in the case of the OR rotors.
Based on the evaluation of the specific surface area,
samples were selected after one, three and ten passes
through the mill for the purpose of determining the
pozzolanic activity. The results of this determination are
shown in Figure 4.
In both cases the first significant increase in evident
pozzolanic activity did not occur until the third pass. A
significant decrease occurred after the tenth cycle, this
being mainly the case for the OR rotors.
4 DISCUSSION
A chemical composition with a high content of SiO2
and a low content of SO3 is typical for conventional sili-
ceous ash. Morphologically, it is a mixture of spherical
cenospheres and porous plerospheres, as is apparent
from the SEM image. Its specific surface area of 494
m2/kg is relatively high, which corresponds to the higher
pozzolanic activity of fly ash at 860 mg of Ca(OH)2/1g,
as determined by the modified Chapelle test. The pozzo-
lanic activity of fly ash tested by this method usually
ranges from 700 to 850 mg Ca(OH)2/1 g.21 This fly ash
can therefore be rated as reactive. In the case of both
types of rotor, a step increase in the specific surface area
occurred after the first pass of the ash through the mill.
The specific surface area increased by about 60 m2/kg to
a final 557 m2/kg in both cases. The growth in specific
surface area is associated with a decrease in the grain
sizes d (0.1), d (0.5) and d (0.9), which is clearly visible
from the particle size distribution curves in Figures 2a
and 2b. As regards the rotors, the AR-BR curve is clearly
wider than that for the OR type. At this stage, especially
large, soft and porous plerospheres are milled very inten-
sively, as evidenced by the SEM images above. However,
the amount of pozzolanic activity is already different at
this stage of the milling process. Fly ash ground on the
AR-BR rotors showed pozzolanic activity that was 24 mg
of Ca(OH)2/1g higher for the same specific surface area
than ash ground on the OR rotors. The results correspond
well with the particle size distribution curves when the
AR-BR sample contains more ultra-fine particles. A
sharp decline in the specific surface area can be observed
after the second and third passes. This phenomenon can
be explained as being the result of the beginning of the
aggregation process due to electrostatic forces. The
lower decrease in the specific surface area value in the
case of the OR rotors than for the AR-BR type can be
explained by the cylindrical shape of the teeth on the OR
rotors. The friction surface over which the particles roll
is smaller with this tooth shape. The charging of the
particles, which leads to aggregation, is thus smaller. The
aggregation process can be clearly observed on the
granulometric curves, which become narrower. However,
in both cases, the pozzolanic activity increased signifi-
cantly. This phenomenon can be explained as being just
due to the formation of aggregates, which are only
loosely bound together by electrostatic forces. Because
the Chapelle test takes place in an aquatic environment,
loosely bonded aggregates can easily disintegrate and the
fine particles can react with the calcium ions very swiftly
and easily. In the case of the AR-BR rotors, the increase
in pozzolanic activity is significantly higher than in the
case of the OR rotors. A possible explanation for this
abnormality is simply that the AR-BR rotor teeth have a
larger contact surface. They thus charge particles and
accelerate the formation of aggregates more than the
cylindrical teeth of the OR rotors, but may also create
more defects on the surfaces of the fly ash particles,
causing the particles to become more reactive. Between
the third and the tenth passes, a gradual decline in spe-
cific surface area can be observed. The particle distribu-
tion curves become narrower and there is a significant
increase in grain size d (0.1). At this point the aggrega-
tion phase has already changed to the stage of agglome-
ration, where the particles are bound by chemical bonds.1
This confirms the results of the pozzolanic activity
determination, and the SEM images. The specific surface
area and pozzolanic activity results achieved by the OR
rotors after ten passes are significantly worse than in the
case of the AR-BR rotors. Because most of the grinding
work in the phase of agglomeration is just consumed in
the breakage of new agglomerates, the results indicate
the AR-BR rotors have a higher efficiency compared to
the OR rotors.
5 CONCLUSION
High-speed grinding in a high-speed disintegrator is
very effective during the first pass of fly ash through the
K. DVOØÁK, I. HÁJKOVÁ: THE EFFECT OF HIGH-SPEED GRINDING TECHNOLOGY ON THE PROPERTIES OF FLY ASH
686 Materiali in tehnologije / Materials and technology 50 (2016) 5, 683–687
Figure 4: The impact of the technology of grinding on the pozzolanic
activity
Slika 4: Vpliv tehnologije mletja na pucolansko aktivnost
mill, when there is a step increase in the specific surface
area, and related pozzolanic activity. During the first pass
there is no aggregation or agglomeration of the grains.
Fly ash can thus be easily homogenized with the other
components of the cement composite, and its properties
can be improved at the same time. The disadvantage of
this type of mill is the process of rapid aggregation and
subsequent agglomeration, which makes it impossible to
achieve a higher specific surface area. However, this
phenomenon can be eliminated by adding grinding aids
or using particle separators, for example. We can con-
clude based on all the achieved results that the use of the
AR-BR rotor with rhomboid-shaped teeth for the grind-
ing of fly ash is more advantageous in the case of a HSG
disintegrator than using the OR rotor with its traditional
cylindrical tooth shape. Higher pozzolanic activity was
achieved for the same specific surface area. Based on the
results, we can say that in the case of fly ash, a high-
speed disintegrator can be a promising means of
improving its properties via grinding and mechanical
activation.
Acknowledgment
This work was financially supported by project No.
LO1408 "AdMaS UP – Advanced Materials, Structures
and Technologies", supported by the Ministry of Educa-
tion, Youth and Sports under "National Sustainability
Programme I" and by project No. 15-08755S: "Study of
the effects of samples preparation on the final properties
of inorganic binders".
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K. DVOØÁK, I. HÁJKOVÁ: THE EFFECT OF HIGH-SPEED GRINDING TECHNOLOGY ON THE PROPERTIES OF FLY ASH
Materiali in tehnologije / Materials and technology 50 (2016) 5, 683–687 687

S. CHERNEVA et al.: INVESTIGATION OF THE MECHANICAL PROPERTIES OF ELECTROCHEMICALLY DEPOSITED ...
689–693
INVESTIGATION OF THE MECHANICAL PROPERTIES OF
ELECTROCHEMICALLY DEPOSITED Au-In ALLOY FILMS
USING NANO-INDENTATION
PREISKAVA MEHANSKIH LASTNOSTI ELEKTROKEMIJSKO
NANE[ENEGA FILMA ZLITINE Au-In Z NANOVTISKOVANJEM
Sabina Cherneva1, Roumen Iankov1, Martin Georgiev2, Tsvetina Dobrovolska2,
Dimitar Stoychev2
1Bulgarian Academy of Sciences, Institute of Mechanics, Acad. G. Bonchev str., Bl. 4, 1113 Sofia, Bulgaria
2Bulgarian Academy of Sciences, Institute of Physical Chemistry, Acad. G. Bonchev str., Bl. 11, 1113 Sofia, Bulgaria
sabina_cherneva@yahoo.com
Prejem rokopisa – received: 2015-06-26; sprejem za objavo – accepted for publication: 2015-09-09
doi:10.17222/mit.2015.129
Thin Au-In alloy films containing different amounts of In were electrochemically deposited on a CuZn substrate with a 500-μm
thickness. The thicknesses of the obtained films varied from 0.4 μm to 2.7 μm. The chemical and phase compositions, as well as
the structures of the films, were investigated by XRF, XRD and SEM analysis. The mechanical properties of the films and
substrates were investigated using nano-indentation experiments. As a result, load–displacement curves were obtained and two
mechanical characteristics of the substrate and investigated films – indentation hardness and indentation modulus – were
calculated using the Oliver & Pharr approximation method. The dependence of the indentation modulus and the indentation
hardness on the depth of the indentation and the content of In, the structure and the phase compositions of the films were
investigated and discussed as well.
Keywords: gold-indium alloy, electrochemical deposition, mechanical properties, nano-indentation
Tanke plasti zlitine Au-In, z razli~no vsebnostjo In, so bile elektrokemijsko nane{ene na podlago iz CuZn, debeline 500 μm.
Debeline dobljene plasti so bile od 0,4 μm to 2,7 μm. Kemijska sestava in sestava faz, kot tudi mikrostruktura plasti, so bile
preiskovane z XRF, XRD in s SEM analizami. Mehanske lastnosti preiskovanih plasti so bile preiskane s pomo~jo preizkusa z
nanovtiskovanjem. Kot rezultat so bile dobljene krivulje obremenitev-raztezek. Dve mehanski lastnosti podlage in preiskovanih
plasti – trdota vtiskovanja in modul vtiskovanja – sta bili izra~unani s pomo~jo Oliver & Pharr metode pribli`ka. Preiskovana in
prediskutirana je bila odvisnost modula vtiskovanja in trdota vtiskovanja na globino vtiska od vsebnosti In.
Klju~ne besede: zlitina zlato-indij, elektrokemijsko nana{anje, mehanske lastnosti, nanovtiskovanje
1 INTRODUCTION
The phase diagram of the gold–indium1 system
shows the presence of several intermetallic compounds
existing at a temperature lower than the melting point of
Indium (~ 156 °C), including stable AuIn and AuIn2 pha-
ses with a cubic lattice, similar to the -phase of gold.
The indium phase starting from a 53 % mass fraction is
tetragonal. The average microhardness obtained for the
metallurgical alloy system Au–In is as follows: Au
(99.999 %) = 0.660,  AuIn (8 % of amount fractions of
In) = 1.700, 1-phase = 3.68, AuIn = 2.73, AuIn2 = 0.77
GPa and the alloy with 80 % of amount fractions of
In = 2.07 GPa.2 The microhardness in the hardened
condition of the compound Au3In2 ( phase) is 1.83 GPa,
and increases in the uniformity of the phase deviations in
the stoichiometric composition.3 In contrast to metallur-
gically obtained, the electrochemically deposited Au-In
alloys are not well studied. At the same time, electroche-
mically deposited thin layers of Au–In alloys will find
wide application (instead of pure gold coatings) in
electrical engineering, micro-electronics, the manufact-
uring of various sensors, the jewelery industry, etc. The
interest in studying the impact of the content of In in the
Au–In alloy on a number of colors, decorative, optical,
corrosion, mechanical and other properties, regardless of
the method of their production, has also increased.4–6 The
aim of the present work is to investigate the indentation
hardness and the indentation modulus of electrochemi-
cally deposited thin layers of Au–In alloys as a function
of the indentation depth and considering the effect on
them of the nature of the substrate, the content of In, the
structure and phase composition of the alloy films as
well as the surface roughness of the films.
2 EXPERIMENTAL PART
The Au-In alloy films with thicknesses between
0.4 μm and 2.7 μm were deposited onto brass sheet sub-
strates (2 cm × 1 cm × 0.03 cm) in a standard electro-
chemical glass cell equipped with two Pt anodes as the
counter electrodes. The standard preliminary treatment
of the brass cathode-substrates includes a procedure for
electrochemical decreasing, followed by pickling in a 20 %
water solution of sulphuric acid at room temperature.
The investigated Au–In alloy films were electrodeposited
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in galvanostatic mode (in the range of cathodic current
densities from 0.2 to 1.8 A dm–2) of an acetate-citrate
electrolyte (containing 1 g/L Au as a metal (KAu (CN)2);
3 g/L In as a metal (InCl3); 90 g/L CH3COONa; 14 g/L
citric acid). The electrolysis process was performed
without any stirring of the electrolyte and at room tem-
perature. The thickness, content of In and percentage
composition (Au:In) of the thin alloy films were deter-
mined by X-ray fluorescence analysis (Fischerscope
XDAL). The structure and morphology of the layer sur-
face were investigated by scanning electron microscopy
(SEM) using a JSM 6390 microscope. The phase com-
position was characterized by X-ray diffraction (XRD)
using a PANalytical Empyrean Equipped with a multi-
channel detector (Pixel 3D) using Cu–K (45 kV-40 mA)
radiation in the 20–115° 2 range with a scan step of
0.01° per 20 s. The mechanical properties of the Au–In
alloy films containing different amounts of In onto the
CuZn substrate were investigated by means of nano-in-
dentation experiments, using a Nano Indenter G200
(Keysight Technologies, USA), equipped with a Berko-
vich three-sided diamond pyramid with a centerline-
to-face angle of 65.3° and a 20-nm radius at the tip of the
indenter. We realized a series of 25 indentations on each
sample probe. We used an indentation method that was
proposed in 7. The indentation hardness and indentation
modulus are determined using the stiffness calculated
from the slope of the load–displacement curve during
each unloading cycle. As a result load–displacement
curves were obtained and two mechanical characteristics
of the substrate and the investigated films – indentation
hardness (HIT) and indentation modulus (EIT) – were
calculated using the Oliver & Pharr approximation
method.8
3 RESULTS AND DISCUSSION
Table 1 shows the results of the XRF analyses on the
chemical composition, thickness () and the micro
roughness (Rz and Ra) of the tested alloy samples and the
brass substrate on which they were deposited. From the
results it can be seen that the interval of changes in
content 49–63 % for the mass fractions of indium in the
resulting thin alloy layers and changes of their micro
roughness. Information about the surface morphology
and the structure of the electrodeposited pure Au and In
coatings of the working electrolyte for the preparation of
the alloy coatings of which in the first case the presence
of In ions is excluded, and in the second case, the
presence of Au ions is excluded, give the microphoto-
graphs presented in Figure 1a and 1f. While the Au
coating is dense and uniform, formed by spherical
crystallites having a size ~ 0.5–1.2 μm (Figure 1a), the
coatings of In have an uneven thickness – over the fine
crystal thin indium layer which covers the entire surface
690 Materiali in tehnologije / Materials and technology 50 (2016) 5, 689–693
S. CHERNEVA et al.: INVESTIGATION OF THE MECHANICAL PROPERTIES OF ELECTROCHEMICALLY DEPOSITED ...
Figure 1: SEM microphotographs of deposited a) Au 100 %: f) In 100
% and Au–In alloy layers in which the content of indium (in mass
fractions, (w/%) is: b) 49.4 % In, c) 54.2 % In, d) 56.0 % In and e) 63
% In (samples No. 2, 7, 3, 4, 5, 6 described in Table 1)
Slika 1: SEM-posnetek nanosa: a) Au 100 %, f) In 100 % in nanosi
AuIn z razli~no vsebnostjo In (v masnih odstotkih, (w/%): b) 49,4 %
In, c) 54,2 % In, d) 56,0 % In, e) 63 % In (vzorci {t. 2, 7, 3, 4, 5, 6
opisani v Tabeli 1)
Table 1: Chemical composition, thickness, Rz and Ra of the investigated Au–In alloy layers and the substrate on which they are deposited
Tabela 1: Kemijska sestava, debelina, Rz in Ra preiskovanih AuIn plasti in podlage, na katero so bile nane{ene
No Sample Content in mass fractions,(w/%) , ìm Ra, ìm Rz, ìm J, A dm
2 deposition time, min
1. Brass substrate (pickled) Cu – 65.80; Zn – 34.20 300 1.61 9.13
2. Àu/Brass Au – 100 0.64 1.13 4.77 1.0; 20
3. AuIn/Brass Au – 50.6; In – 49.4 0.56 1.14 4.90 1.8; 7
4. AuIn/Brass Au – 45.8; In – 54.2 0.75 1.40 5.37 1.2; 15
5. AuIn/Brass Au – 44.1; In – 56.0 1.42 1.15 5.00 0.6; 20
6. AuIn/Brass Au – 37.0; In – 63.0 2.76 1.50 8.93 0.2; 30
7. In/Brass In – 100 0.49 1.18 3.83 1.0; 20
of the brass substrate, they grew, not fully coalesced,
spheroidal agglomerates with size ~ 1–10 μm (Figure
1f). The influence of changes in the content of indium on
the surface morphology and structure of the Au–In alloy
layers is presented in Figures 1b to 1e. From the
microphotographs it is clear that at the lowest content of
indium (49.4 %) (Figure 1b) the film has a morphology
and structure that is different from that of the pure gold
film. The alloy coating is formed by homogeneously dis-
tributed agglomerates of a size of the base several times
larger than that of the spherical grains constituting the
gold coating. Moreover, there was no phase hetero-
geneity in the regime of back scattering electrons. The
reasons for this conclusion give the images on the left-
hand side of the SPI electron microscopic image (Figure
1b), obtained in the regime of back scattering electrons
(BEI), while the right-hand part of the photograph shows
an image that was obtained in the regime of a secon-
dary-electron image (SEI). With the same purpose (the
recording of possible phase heterogeneity) are the SPI
electron microscopic images for a higher content of In
(Figures 1c to 1e). Increasing the content of indium in
the alloy layer to ~ 54 % (Figure 1c), leads to a levelling
of the morphology, respectively, to a finer structure
compared with those at a content of ~ 49 % (Figure 1b),
in which an even greater degree was observed in the next
amount (~ 56 %) of indium (Figure 1d). Obviously, the
observed differences in morphology are not related to the
phase, but are related to the topographic heterogeneity.
When the content of co-deposited In, however, reached
63 %, the morphology and the structure drastically change;
they are characterized by large aggregates (3–15 μm),
composed of crystallites with dimensions of 0.5–1 μm.
Moreover, the BEI image (left-hand side) of the electron
microscopic micrograph, at this content (Figure 1e)
S. CHERNEVA et al.: INVESTIGATION OF THE MECHANICAL PROPERTIES OF ELECTROCHEMICALLY DEPOSITED ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 689–693 691
Figure 5: Dependence of the indentation hardness on the content of In
Slika 5: Odvisnost trdote vtiskovanja od vsebnosti In
Figure 3: Dependence of the indentation hardness on the depth of the
indentation
Slika 3: Odvisnost trdote vtiskovanja od globine vtiskovanja
Figure 4: Dependence of indentation modulus on the depth of the
indentation
Slika 4: Odvisnost modula vtiskovanja od globine vtiskovanja
Figure 2: XRD patterns of the: a) Au 100 %, f) In 100 % and Au–In
alloy layers, containing: b) 49.4 % In, c) 54.2 % In, d) 56.0 % In,
e) 63 % In; (o) – reflections of CuZn substrates, () – reflections of
AuIn2
Slika 2: Rentgenogram: a) Au 100 %, f) In 100 % in AuIn nanosa z:
b) 49,4 % In, c) 54,2 % In, d) 56,0 % In, e) 63 % In; (o) odboji CuZn
podlage, () odboji AuIn2
indicates the occurrence of phase heterogeneity. The
X-ray phase analysis of the same alloy samples (Figure
2) showed that for electrodeposited samples of electro-
lyte containing only gold ions (in the absence of indium
ions) the diffractogram showed the presence of reflec-
tions of the cubic lattice (a = b = c = 4.08) of the phase
of gold (pdf 98-004-4362), and reflections of the brass
substrate (pdf 98-062-9457) (Figure 2a). In the case of
the electrodeposition of an indium film (in the absence of
gold ions in the electrolyte), the diffractogram of the
obtained sample indicates the presence of reflections of
an In tetragonal phase (pdf 98 005-3091) with the lattice
parameters a = 0.3253 nm, b = 0.49455 nm and reflec-
tions of the substrate made of brass (pdf 98-062-9457)
(Figure 2f). Since the content of indium in the electro-
deposited alloy film is in the range 49–63 % of mass
fractions of In, then, according to the phase diagram for
the system Au–In, they fall into the area of the phase
AuIn2. This is confirmed by the recorded reflections in
the diffractograms presented in Figures 2b to 2e. The
AuIn2 phase has a cubic lattice with highly expanded
parameters regarding the -phase of the gold (a = b = c =
6.502). Only in the alloy composition containing over 63
% of mass fractions of In (Figure 2e) is there a presence
of both the phase AuIn2 and the tetragonal phase of In
(pdf 98005-3091), which is the most likely cause for
registered, strongly emphasized, morphological hetero-
geneity of the alloy coating (Figure 1). The dependence
of the indentation hardness and the indentation modulus
of the investigated alloy films on the depth of the inden-
tation are shown in Figures 3 and 4. With an increasing
depth of indentation, the indentation hardness and the
modulus change a great deal. There are two possible
reasons for this: the influence of the substrate and the
effect of the difference in the structure with depth. The
dependence of the indentation hardness and the inden-
tation modulus of the investigated alloy films (at load =
1.15 mN, in order to be far enough from the influence of
the substrate) on the content of In is shown in Figures 5
and 6. With an increase of the content of In from 0 to
49.4 %, the indentation hardness increases too, and after
this (from 54.2 % to 63 % content of In) it starts to
decrease. It is obvious from Figure 6 that with an
increase in the In content up to 56.0 % and 63 % the
indentation modulus of the investigated Au–In films
decreases. This could be explained by the influence of
the simultaneously existing two phases on the surface
electrode: In and AuIn2. The effect of the non-regularity
is very strong due to the different type of crystal lattice –
tetragonal in case of the In phase and cubic in the case of
the AuIn2 phase. Most probably, the inhomogeneity of
these two phases, as well as their roughness limit the
accuracy, due to the randomly of both phases during the
performed measurements.
4 CONCLUSIONS
In the present work the mechanical properties of
electrochemically deposited thin layers of Au–In alloys
as a function of the indentation depth and considering
the effect on them of the nature of the substrate, the
content of In, the structure and the phase composition of
the alloy films as well as the surface roughness of the
films were investigated. The results showed that with an
increasing content of In from 0 % to 49.4 %, the indenta-
tion hardness increased too and after this (from 54.2 %
to 63 % content of In) it starts to decrease. Moreover,
with an increase in the In content up to 56.0 % and 63 %
the indentation modulus of the investigated Au-In films
decreases. This could be explained by the influence of
the simultaneously existing two phases on the surface
electrode: In and AuIn2. The effect of the non-regularity
is very strong due to the different types of crystal lattice:
tetragonal in the case of the In phase and cubic in the
case of the AuIn2 phase.
Acknowledgments
Authors gratefully acknowledge the financial support
of Bulgarian National Science Fund under Grant No.
T02-22/12.12.2014.
5 REFERENCES
1 T. Massalski, J. Murray, B. Lawrence, B. Hugh, Binary Alloy Phase
Diagram, American Society for Metals, Metals Park, Ohio, 1 (1986)
90, 260–270
2 G. W. Powell, J. D. Braun, Diffusion in the gold-indium system,
Transactions of the Metallurgical Society of AIME, 230 (1964) 4,
694–699
3 V. K. Nikitina, A. A. Babitsina, U. K. Lobanova, Phase diagrams of
the system Au-In, Inorganic materials (in russian), Izvestia AN
SSSR, 7 (1971) 3, 421–427
4 L. C. Archibald, G. Sanderson, Electrodeposition of a White
Gold-Indium Alloy from an Acid Cyanide Electrolyte, Transactions
of the Institute of Metal Finishing, 55 (1978) 4, 149–154
5 C. Cretu, E. Van der Lingen, Coloured Gold Alloys, Gold Bulletin,
32 (1999) 4, 115–126, doi: 10.1007/BF03214796
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692 Materiali in tehnologije / Materials and technology 50 (2016) 5, 689–693
Figure 6: Dependence of the indentation modulus on the content of In
Slika 6: Odvisnost modula vtiskovanja od vsebnosti In
6 U. E. Klotz, Metallurgy and processing of coloured gold inter-
metallics – Part I: Properties and surface processing, Gold Bulletin,
43 (2010) 1, 4–10, doi: 10.1007/BF03214961
7 M. Datcheva, S. Cherneva, D. Stoychev, R. Iankov, M. Stoycheva,
Determination of Anodized Aluminum Material Characteristics by
Means of Nanoindentation Measurements, Materials Sciences and
Applications, 2 (2011) 10, 1452–1464, doi:10.4236/msa.2011.
210196
8 W. Oliver, G. Pharr, Measurement of hardness and elastic modulus
by instrumented indentation: Advances in understanding and
refinements to methodology, Journal of Materials Research, 19
(2004) 1, 3–20
S. CHERNEVA et al.: INVESTIGATION OF THE MECHANICAL PROPERTIES OF ELECTROCHEMICALLY DEPOSITED ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 689–693 693

A. ROUSTA, H. R. DIZAJI: GROWTH OF K2CO3-DOPED KDP CRYSTAL FROM AN AQUEOUS SOLUTION ...
695–698
GROWTH OF K2CO3-DOPED KDP CRYSTAL FROM AN
AQUEOUS SOLUTION AND AN INVESTIGATION OF ITS
PHYSICAL PROPERTIES
RAST KDP KRISTALOV Z DODATKOM K2CO3 IZ VODNE
RAZTOPINE IN PREISKAVA NJIHOVIH FIZIKALNIH LASTNOSTI
Abrisham Rousta1, Hamid Rezagholipour Dizaji2
1Islamic Azad University, Central Tehran branch, Faculty of Basic Sciences, Physics Department, Crystal Growth Laboratory,
2 Rasooli Alley, North Jamalzadeh st., Tehran, Iran
2Semnan University, Faculty of Physics, Semnan, Iran
hrgholipour@semnan.ac.ir
Prejem rokopisa – received: 2015-06-26; sprejem za objavo – accepted for publication: 2015-10-09
doi:10.17222/mit.2015.128
In this present work, KDP and 2M%-K2CO3-doped KDP crystals were grown by a slow-evaporation solution technique. The
grown crystals were characterized by Fourier Transform Infrared (FT-IR) spectroscopy, X-ray diffractometry (XRD), UV-Vis
spectroscopy, and laser damage threshold (LDT) analysis. The presence of the functional groups of the grown crystals was
identified from the FT-IR spectra. The XRD tests showed that the grown crystals had a tetragonal structure. A comparison of the
optical transmission of the grown crystals revealed that the K2CO3-doped KDP crystal had a higher transmission than the pure
KDP crystal for the entire UV and visible region.
Keywords: growth from solution, slow-evaporation solution technique, KDP crystal, K2CO3 additive
V predstavljenem delu so KDP in z 2M % K2CO3 dopirani KDP kristali rasli s tehniko po~asnega izhlapevanja teko~ine. Dob-
ljeni kristali so bili karakterizirani iz infrarde~o spektroskopijo s Fourierjevo transformacijo (FT-IR), z rentgensko difrakcijo
(XRD), z UV-Vis spektroskopijo in s pragom po{kodbe z laserjem (LDT). Prisotnost funkcionalnih skupin kristalov v rasti je
bila dolo~ena s FT-IR spektrom. XRD je pokazal, da imajo rasto~i kristali tetragonalno zgradbo. Primerjava prepustnosti
svetlobe v rasto~ih kristalih je odkrila, da imajo KDP kristali, dopirani s K2CO3, bolj{o prepustnost kot pa ~isti KDP v celotnem
UV in v vidnem podro~ju.
Klju~ne besede: rast iz raztopine, tehnika po~asnega izparevanja raztopine, KDP kristal, dodatek K2CO3
1 INTRODUCTION
Potassium dihydrogen phosphate KH2PO4 (KDP) is a
material that is soluble in water with a positive solubility
coefficient. The crystal structure is tetragonal with the
lattice parameters a = b = 0.7448 nm and c = 0.6977 nm.
A KDP single crystal is piezoelectric at room tempera-
ture, and below 123 K (Curie point) it transforms to the
ferroelectric phase and has an orthorhombic structure.
This crystal is an excellent electro-optic and nonlinear
optical (NLO) material, so it is used in optical modu-
lators such as a second-harmonic generator. In addition,
it is characterized by its good UV-visible transmission,
high damage threshold, etc. Many attempts have been
made to modify its properties, either by changing the
growth condition or by adding different impurities.1–7
P. V. Dhanaraj et al.8 showed that the addition of
K2CO3 could make the KDP solution more stable than
with other additives and enhanced the metastable zone
width of the KDP solution for all temperatures.8 They
also found that the laser-induced damage threshold of a
K2CO3-added KDP crystal was higher than that of the
pure KDP crystal.
In the present study, pure and 2M%-K2CO3-doped
KDP crystals were grown from an aqueous solution
using the slow-evaporation method at room temperature.
The grown crystals were then subjected to various
characterization techniques.
2 EXPERIMENTAL PROCEDURE
KDP crystals, pure and with added 2M% K2CO3,
were grown from an aqueous solution using the slow-
evaporation method at room temperature. A saturated
solution of KDP was prepared by dissolving an appro-
priate amount of commercially available KDP powder in
double distilled water without any further purification.
The solution was then stirred well for two hours using a
magnetic stirrer, filtered using Whatmann filter paper
and transferred into the growth container for the slow
evaporation. In a similar way, a saturated solution of
KDP with added 2 M% K2CO3 was prepared. Then each
container was covered with a perforated cover and kept
in a dust-free place. Within two weeks, transparent KDP
crystals of both pure (22 mm × 21 mm × 8 mm) and with
added 2 M% K2CO3 (24 mm × 20 mm × 14 mm) were
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grown. Figures 1a and 1b show the pure and doped
crystals, respectively.
3 RESULTS AND DISCUSSION
3.1 X-ray diffraction analysis (XRD)
Both the pure and 2M%-K2CO3-added KDP crystals
were subjected to powder X-ray diffraction (XRD)
analysis using an X-ray diffractometer (Advance Model
D8) with high-intensity Cu-K radiation ( = 0.15406 nm).
The grown crystals were ground using an agate mortar
and pestle in order to determine the crystal phases by
XRD. Figure 2 shows the XRD patterns of the pure and
doped KDP single crystals. From the data, both the
crystals were found to crystallize in the tetragonal
system. Comparing the two patterns, we found no extra
peaks due to the doping; hence adding K2CO3 to KDP
did not affect its crystal structure. The sharp peaks
observed in both patterns indicate the good crystallinity
of the grown crystals.
3.2 Optical transmittance
The optical transmittance spectrum in the wavelength
region 200–800 nm was recorded at room temperature
using a Perkin Elmer model lambda25 UV-Vis-NIR
spectro-photometer on a 1.8-mm-thick plate of the
grown K2CO3-added KDP crystal in the (001) direction.
This property is the most desirable one for an NLO
material. Figure 3 presents a comparison among the
transmittance spectra of pure and 5M%-K2CO3-added
KDP single crystals7 and a 2 M%-K2CO3-added KDP
single crystal (present work).
It is clear from the figure that the crystals are highly
transparent across the entire UV-visible region. It is also
obvious that the transmittance percentage of the doped
KDP crystal is higher than that of the pure one. This
improvement in the transparency of the KDP crystal after
the addition of K2CO3 to the solution may be attributed
to its ability to suppress the inclusions due to the heavy
metals usually present in the starting material.
A. ROUSTA, H. R. DIZAJI: GROWTH OF K2CO3-DOPED KDP CRYSTAL FROM AN AQUEOUS SOLUTION ...
696 Materiali in tehnologije / Materials and technology 50 (2016) 5, 695–698
Figure 3: UV-visible transmittance spectra of: a) pure KDP and b)
5M%-K2CO3-added KDP single crystals6, and c) 2 M%-K2CO3-added
KDP single crystal
Slika 3: Spekter prepustnosti UV in vidne svetlobe: a) ~isti KDP in b)
monokristali KDP z dodatkom 5M % K2CO36 in c) monokristal KDP
z dodatkom 2 M % K2CO3
Figure 1: Photographs of: a) pure and b) 2 M%-K2CO3-doped KDP
crystals
Slika 1: Posnetka: a) ~istega in b) z 2 M% K2CO3 dopiranega KDP
kristala
Figure 2: X-ray diffraction patterns of: a) pure KDP and b) 2 M%-
K2CO3-added KDP single crystals
Slika 2: Rentgenograma: a) ~isti KDP in b) KDP monokristali z
dodatkom 2 M % K2CO3
3.3 Fourier-transform infrared (FT-IR) analysis
The FT-IR spectra of the pure and 2 M%-K2CO3-
doped KDP crystals were recorded using a Perkin Elmer
model 410 Jasco company spectrometer in the wave-
number range from 400 cm–1 to 4000 cm–1 using the KBr
pellet technique. Figure 4 represents the FT-IR spectra
of the grown crystals. Also, the position of the peaks and
their functional group assignments are given in Table 1.
Table 1: Observed FT-IR wave numbers (cm–1) and their functional
group assignments for the grown pure and 2M%-K2CO3-added KDP
crystal
Tabela 1: Opa`ena FT-IR valovna {tevila (cm–1) in dodeljene funk-
cionalne skupine pri ~istem KDP kristalu in KDP kristalu z dodatkom
2M % K2CO3
Wave number (cm–1)
Functional group
assignmentsPure KDP
crystal
2M %K2CO3
added KDP
crystal
3739.30 3744.12 Free O-H stretchinghydrogen bonded
3448.10 3427.85 O-H stretching hydrogenbonded
2489.65 – O=P-OH asymmetricstretching
1649.80 1649.80 O-H bending out of plate
1302.68 1301.72 O-P=O stretching
1100.19 1096.33 P=O stretching
894.81 898.67 P-O stretching
537.08 539.97 O-P-O bending
473.44 470.55 PO4 stretching
The broad band that appears in the range from 3800
cm–1 to 2500 cm–1 is due to free O-H stretching of the
KDP. These functional groups arise at 3739 cm–1 and
3448 cm–1 in pure KDP and at 3744 cm–1 and 3427 cm–1
in doped KDP.
The peaks at 537 cm–1, 1100 cm–1 and 2489 cm–1 in
pure KDP and 539 cm–1 and 1096 cm–1 in doped KDP
are due to the O-P-O bending, P=O stretching and
O=P-OH asymmetric stretching of the KDP, respectively.
The O-P=O stretching, P-O stretching and PO4 stretching
are found at 1302 cm–1, 894 cm–1 and 473 cm–1 in pure
KDP and 1301 cm–1, 898 cm–1 and 470 cm–1 in doped
KDP, respectively.
We can clearly see from the comparison of the FTIR
spectrum of the two crystals that the presence of the do-
pant has led to a change in the intensity of the absorption
of the IR frequencies and a slight shift in some of the
frequencies. The strong similarities of the two graphs
reveal the fact that the peaks corresponding to pure KDP
crystal are predominant over those corresponding to the
K2CO3-added KDP crystal, which may be due to the
small amount of doped K2CO3 in the compound com-
pared to the KDP.
3.4 Laser damage threshold (LDT)
One of the most important considerations when
selecting a material for nonlinear optics applications is
its ability to withstand high power intensities.9
The laser damage threshold (LDT) of nonlinear
optical components depends on physical and chemical
factors, particularly imperfections, defects and the con-
centration and type of the impurities, etc.10
The LDT of the grown pure and 2M%-K2CO3-doped
KDP single crystals was carried out using a Nd:YAG
laser with the wavelength 1064 nm and shot-to-shot
mode, with an energy per pulse of 50 mJ, a repetition
rate of 10 Hz, a pulse duration of 7 ns, a beam waist of
0.0841 mm and a spot diameter of 2.3 mm. The laser
beam was focused on the sample with 1-m and 50-cm
focal-length lenses. A Tektronic 2430A digital oscillo-
scope was used to record the energy of every pulse and
save the data in a computer.
The damage threshold was calculated for the pure
KDP and the 2 M%-K2CO3-doped KDP crystals and the
value was found to be 12.1083 J/cm2 and 19.1782 J/cm2,
respectively. It shows that, adding 2M% K2CO3 to the
KDP solution increased the damage threshold of the
KDP single crystal, which can be attributed to the ability
of this additive to suppress the inclusions due to heavy-
metal impurities like Cr3+, Fe3+, and Al3+ that are present
in most of the commercially available chemicals.
4 CONCLUSIONS
In this work, pure KDP and 2M%-K2CO3-doped
KDP crystals were grown by the slow-evaporation solu-
tion technique from an aqueous solution at room tempe-
rature. Structural studies indicate that both the grown
A. ROUSTA, H. R. DIZAJI: GROWTH OF K2CO3-DOPED KDP CRYSTAL FROM AN AQUEOUS SOLUTION ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 695–698 697
Figure 4: FT-IR spectra of: a) pure KDP and b) 2 M% K2CO3 added
KDP single crystals
Slika 4: FT-IR-spektri: a) ~isti KDP in b) KDP monokristal z dodat-
kom 2 M% K2CO3
crystals have a tetragonal system with similar XRD
patterns. Optical transmission studies showed that using
2 M% K2CO3 as an additive increased the optical quality
of the KDP crystal compared to the pure crystal. Fourier
transform infrared analysis revealed the functional
groups of the samples. The laser-damage threshold of the
2 M%-K2CO3-added KDP crystal was found to be higher
than that of the undoped crystal, indicating the suitability
of K2CO3 as an additive to enable KDP to withstand high
power intensities.
Acknowledgments
This work is based on a Master of Science thesis that
was supported by the Islamic Azad University, Central
Tehran Branch.
The authors are thankful to Dr. M. Nikpour, assistant
Prof. of Chemistry Department and Mr. N. Madani
faculty member of Physics Department, Islamic Azad
University, Central Tehran Branch for their assistance.
5 REFERENCES
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and Technology, 3 (2013), 92–96, doi:10.4236/jcpt.2013.33015
2 S. Balamurugan, P. Ramasamy, Y. Inkong, P. Manyum, Effect of KCl
on the bulk growth KDP crystals by Sankaranarayanan-Ramasamy
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4 Y. Shangfeng, S. Genbo, L. Zhengdong, J. Rihong, Rapid growth of
KH2PO4 crystals in aqueous solution with additives, Journal of
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ijret.2013.0212125
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a Unidirectional EDTA Added KDP Single Crystal by the S-R
Method, Chinese Journal of Physics, 50 (2012), 652–658,
doi:10.6122/CJP
7 X. G. Xu, X. Sun, Z. P. Wang, Z. S. Shao, Z. S. Gao, Abnormal
optical properties in doped H3BO3 KDP crystals, Journal of Crystal
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01130-0
8 P. V. Dhanaraj, C. K. Mahadevan, C. K. Bhagavannarayana, P.
Ramasamy, N. P. Rajesh, Growth and characterization of KDP
crystals with potassium carbonate asadditive, Journal of Crystal
Growth, 310 (2008), 5341–5346, doi:10.1016/j.jcrysgro.2008.09.019
9 P. Rajesh, S. Sreedhar, K. Boopathi, S. Venugopal Rao, P. Rama-
samy, Enhancement of the crystalline perfection of directed KDP
single crystal, Current Applied Physics, 11 (2011), 1343–1348,
doi:10.1016/j.cap.2011.03.076
10 N. Balamurugan, P. Ramasamy, Investigation of the Growth Rate
Formula and Bulk Laser Damage Threshold KDP Crystal Growth
from Aqueous Solution by the Sankaranarayanan-Ramasamy (SR)
Method, Crystal Growth & Design, 6 (2006), 1642–1644,
doi:10.1021/cg050680n
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698 Materiali in tehnologije / Materials and technology 50 (2016) 5, 695–698
T. TAÑSKI et al.: SURFACE TREATMENT OF HEAT-TREATED CAST MAGNESIUM AND ALUMINIUM ALLOYS
699–706
SURFACE TREATMENT OF HEAT-TREATED CAST MAGNESIUM
AND ALUMINIUM ALLOYS
OBDELAVA POVR[INE TOPLOTNO OBDELANIH MAGNEZIJEVIH
IN ALUMINIJEVIH LIVNIH ZLITIN
Tomasz Tañski, Maciej Wiœniowski, Wiktor Matysiak, Marcin Staszuk, Rados³av Szklarek
Silesian University of Technology, Institute of Engineering Materials and Biomaterials, Konarskiego Str. 18A, 44-100 Gliwice, Poland
tomasz.tanski@polsl.pl
Prejem rokopisa – received: 2015-06-26; sprejem za objavo – accepted for publication: 2015-10-12
doi:10.17222/mit.2015.132
Modern coating systems deposited on surface layers of structural light materials are currently one of the most important issues
in up-to-date material engineering, where vacuum deposition techniques are often used to improve the mechanical and func-
tional properties of produced surface layers. Presented in this paper are gradient and monolithic coating types: Ti/Ti(C,N)/CrN,
Ti/Ti(C,N)/(Ti,Al)N, Ti/(Ti,Si)N/(Ti,Si)N, Cr/CrN/CrN, Cr/CrN/TiN and Ti/DLC/DLC deposited onto magnesium and
aluminium alloy substrates with the cathodic-arc-evaporation method (Arc PVD) and plasma-assisted process (PA CVD). Addi-
tionally, a thin metallic layer – in micrometers– (Cr and Ti) was deposited prior to the deposition of the final gradient coating to
improve its adhesion to the substrate. This work presents the investigation results concerning the obtained surface-layer micro-
structures and mechanical properties of the obtained bi-layer coatings (gradient/multicompound) deposited onto light-alloy
substrates using the chosen PVD and CVD methods, especially to meet the requirements needed for light-metal substrates – low
temperature and duration. The structure investigations of the deposited coating were performed using a scanning electron
microscopy (SEM) and glow discharge optical emission spectrometry (GDOES); the mechanical and functional properties were
examined using the ball-on-disk method for the wear-resistance determination, and microhardness tests were performed for the
functional usability of the coatings. The main finding is that the fracture morphology is characterized by a lack of columnar
structures in the obtained coatings. The metallographic examinations carried out proved that the coatings were deposited uni-
formly over the whole sample, onto the investigated substrate materials; the measured thickness is characteristic for the
produced coating type.It was also found that the particular layers adhere tightly to each other and to the light-metal substrate.
The investigation results of the up-to-date PVD methods, together with light alloys, led to obtaining new applications, especially
in the automobile and aviation industries.
Keywords: light alloys, PVD, CVD, structure, properties
Moderni sistemi nanosov na povr{inskih plasteh lahkih konstrukcijskih materialov so eden od najpomembnej{ih izzivov v
in`eniringu materialov, kjer se za izbolj{anje mehanskih in funkcionalnih lastnosti plasti na povr{ini pogosto uporabljajo tehnike
vakuumske depozicije. V ~lanku so predstavljeni gradientni in monolitni nanosi vrst: Ti/Ti(C,N)/CrN, Ti/Ti(C,N)/(Ti,Al)N,
Ti/(Ti,Si)N/(Ti,Si)N, Cr/CrN/CrN, Cr/CrN/TiN in Ti/DLC/DLC ki so bili nane{eni z metodo katodnega izparevanja v obloku
(Arc PVD) in s plazemskim postopkom (PA CVD). Dodatno je bila nane{ena tanka kovinska plast (Cr in Ti), debelina v mikro-
metrih, in sicer pred nanosom kon~nega gradientnega nanosa, da bi se izbolj{ala njegova oprijemljivost na podlago. ^lanek
predstavlja rezultate raziskave mikrostrukture povr{inskega nanosa in mehanske lastnosti dvoplastnega nanosa
(gradient/multicompound), nane{enega na podlago iz lahke zlitine, s pomo~jo izbranih metod PVD in CVD, da bi zagotovili
zahtevam podlage iz lahke kovine – nizka temperatura in kratko trajanje. Preiskave zgradbe nanosa so bile izvedene s pomo~jo
vrsti~ne elektronske mikroskopije (SEM) in z razelektritveno opti~no emisijsko spektrometrijo (GDOES), medtem ko so bile
mehanske in funkcionalne lastnosti, preiskane z uporabo metode kroglica na plo{~i za dolo~anje obrabne odpornosti ter
mikrotrdote za funkcionalno uporabnost nanosov. Glavna ugotovitev je, da v morfologiji preloma nanosov ni stebraste zgradbe.
Izvedene metalografske preiskave so pokazale, da so nanosi enakomerno nane{eni po vsej povr{ini preiskovane podlage,
izmerjene debeline so zna~ilne za to vrsto nanosov in ugotovljeno je tudi, da se posamezni nanosi med seboj tesno stikajo, tudi s
podlago iz lahke kovine. Rezultati raziskav ka`ejo, da uporaba sodobnih PVD metod, skupaj z lahkimi zlitinami, omogo~a nove
aplikacije, posebno v avtomobilski in letalski industriji.
Klju~ne besede: lahke zlitine, PVD, CVD, struktura, lastnosti
1 INTRODUCTION
Dynamic industry development introduces an escala-
tion of requirements concerning new needs and working
conditions, which facilitateand direct theprogress within
material engineering, especially in the case of fabrication
and examination of new materials.1–3 Properties of many
products and their elements depend not only on the pos-
sibility of transmitting mechanical loads through the
whole active intersection or the material’s physical and
chemical properties, but also on the structure and proper-
ties of the surface layer. The use of surface layers, fulfill-
ing thehigh requirements withsoft and cheap cores, is a
great way of reducing expenses. A wide range of avail-
able layers and ways of theirmodification facilitate the
design of thebest combination of core and layer proper-
ties possible. Modern surface-engineering techniques in-
cluding the use of a corrosion- and abrasive-resistant
hard material, despite the maintenance of accurate prop-
erties, should allow usto combine esthetic values and
ecological production.3
With many available techniques of improving engi-
neering materials, the physical vapordeposition (PVD),
Materiali in tehnologije / Materials and technology 50 (2016) 5, 699–706 699
UDK 621.7.015:621.78:669.721.5:669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)699(2016)
chemical vapordeposition (CVD) and hybrid methods
(which enable a full control of thechemical composition,
structure and properties using characteristics of particu-
lar methods like CVD, PVD and conventional thermo-
chemical treatments – thermal spraying + heat treatment,
nitriding or cyaniding + pulsed-laser deposition (PLD),
autocatalytic layer deposition + plasma-assisted process-
ing) are essential.4–20 Anotherimportant surface engineer-
ing technology, applied inlight-alloy processing, is the
laser treatment including remelting and alloying.3 These
techniques allow us to make layers with special proper-
ties (high hardness and tribological resistance combined
with constant substrate properties) and the thickness in a
range from tenths of a millimeter to even a few millime-
ter scan be achieved. A layer obtained with the laser-al-
loying or remelting technique has a different structure
and properties than those of the base or the alloying ele-
ments.3 The morphology of a quasi-composite layer is
homogeneous and exhibits a proper dispersion of the al-
loying elements into the whole depth except for a very
thin diffusion-saturation layer.
The aim of this research was to obtain a hard coat for
a soft core – such a material can resist different amounts
of load (depending on many factors) because of the coat-
ing and thanks to the soft core, the internal forces are
transferred and reduced inside. In some cases, the corro-
sion resistance is also observed, which is very desirable.
An important part of the investigationdone on the mate-
rial was the examination of the structures and mechani-
cal properties of gradient/monolithic coatings deposited
with the PVD and CVD methods onto magnesium and
aluminum casting alloys after the heat treatment.5,9,16,18
2 EXPERIMENTAL PART
The materials used for the investigation includedcast
magnesium and aluminium alloys, whosechemical com-
positions are presented in Table 1. The deposition of
coatings Ti/Ti(C,N)/CrN, Ti/Ti(C,N)/(Ti,Al)N,
Ti/(Ti,Si)N/(Ti,Si)N, Cr/CrN/CrN, Cr/CrN/TiN and
Ti/DLC/DLC was made within a device based on the
CAE PVD method in anatmosphere of Ar, N2 and C2H2;
moreover, the DLC coating wasdeposited using acety-
lene (C2H2) as the precursor and was produced with the
PA CVD method. A gradient change in thechemical
composition ofthe PVD coatings’ cross-sections was
achieved by changing the proportion of the reactive-gas
dose or a variation in thearc-source current. The DLC
coating was characterized byavariation in the silicon
(Me) concentration, demonstrating that a gradient layer
wasobtained. Silicon was supplied to the furnace cham-
ber from the gas phase, Ti/a-C:H-Me/a-C:H. Cathodes
containing pure metals (Cr, Ti) and alloys of TiAl and
TiSi (50:50 % amount fractions) were used for the depo-
sition of the coatings. The diameter of the used cathodes
was 65 mm. The temperature was controlled with
T. TAÑSKI et al.: SURFACE TREATMENT OF HEAT-TREATED CAST MAGNESIUM AND ALUMINIUM ALLOYS
700 Materiali in tehnologije / Materials and technology 50 (2016) 5, 699–706
Table 1: Chemical compositions oftheinvestigated alloys
Tabela 1: Kemijska sestava preiskovanih zlitin
Type of material Mass concentration of the elements, in mass fractions (w/%)
Al Zn Mn Si Mg Fe Cu Rest
Magnesium alloy – AZ91 9.09 0.77 0.21 0.04 89.8 0.011 – 0.079
Magnesium alloy – AZ61 5.92 0.49 0.15 0.04 93.3 0.007 – 0.093
Aluminium alloy – AlSi9Cu4 85.4 0.05 0.01 9.27 0.28 0.34 4.64 0.01
Aluminium alloy – AlSi9Cu 88.86 0.16 0.37 9.1 0.27 0.18 1.05 0.01
Table 2: Deposition parameters of the investigated coatings
Tabela 2: Parametri nana{anja preiskovanih nanosov
Coating parameters
Type of the achieved coating and the applied coating technique
PVD CVD
Ti/Ti(C,N)-gra-
dient/CrN
Ti/Ti(C,N)-gra-
dient/(Ti,Al)N
Cr/CrN-gradi-
ent/CrN
Cr/CrN-gradi-
ent/TiN
Ti/(Ti,Si)N-grad
ient/(Ti,Si)N
Ti/DLC-gradi-
ent/DLC
Base pressure (Pa) 5×10–3 5×10–3 5×10–3 5×10–3 5×10–3 1×10–3
Working pressure (Pa) 0.9/1.1-1.9/2.2 0.9/1.1-1.9/2.8 1.0/1.4-2.3/2.2 1.0/1.4-2.3/2.2 0.89/1.5-2.9/2.9 2
Argon flow rate (sccm)
80* 80* 80* 80* 80* 80*
10** 10** 80** 80** 20** –
10*** 10*** 20*** 20*** 20*** –
Nitrogen flow rate (sccm) 225→0** 0→225** 0→250** 0→250** 0→300** –
250*** 350*** 250*** 250***
Acetylene flow rate (sccm) 0→170** 140→0** – – 230
Substrate bias voltage (V)
70* 70* 60* 60* 70* 500
70** 70** 60** 60** 100**
60*** 70*** 60*** 100*** 100***
Target current (A) 60 60 60 60 60 -
Process temperature (°C) <150 <150 <150 <150 <150 <180
*during the metallic-layer deposition, **during the gradient-layer deposition, *** during the ceramic-layer deposition
thermocouples. To improve the adhesion of the coatings,
a transition Cr, Ti interlayer was deposited. The working
pressure during the deposition process was 2–4 Pa, de-
pending on the coating type. The distance between the
cathodes and the deposited substrates was 120 mm. Just
before the coating-deposition process, the specimens
were prepared with the standard procedures of grinding,
polishing and chemical cleaning using multi-stage wash-
ing in an ultrasonic cleaner and a cascade washer, then
dried in hot air. The next step was ion etching in the
chamber to clean the surfaces atthe atomic scale and to
activate them. The parameters used were asubstrate-po-
larization voltage of 800/200 V and a period of 20 min.
The conditions of the coating deposition are presented in
Table 2.
The investigations of the microstructures, micro-area
qualitative and quantitative chemical compositionswere
performed using a scanning electron microscope (SEM)
ZEISS Supra 35. The cross-sectional atomic composition
of the samples (coating and substrate) was obtained by
using the glow discharge optical spectrometer GDOS-
750 QDP from Leco Instruments.
The microhardness tests of the coatings were made
with a SHIMADZU DUH 202 ultra-microhardness
tester. The measurements were made with a 10 mN load,
to eliminate the substrate influence on the coating hard-
ness. Wear-resistance investigations were performed us-
ing the ball-on-disk method. A tungsten carbide ball with
a diameter of 3 mm was used as the counter part. The
tests were performed at room temperature overa defined
time using the following test conditions: a load of Fn =
5N, a rotation of the disk of 200 min–1 20.94 r/s, a wear
radius of 2.5 mm and a sliding speed of 0.05 m/s.
3 RESULTS AND DISCUSSION
In order to determine the structures and relationships
between the type of substrate and the types of hybrid lay-
ers, the metallographic investigation was done under the
technological conditions (the soft substrate – the gradient
intermediate layer able to easily change the concentra-
tion of one or a few components between the base and
the surface – and the external layer) of the cathodic-arc
deposition, Arc-PVD, and plasma-assisted chemical va-
por deposition, PA-CVD processes.
The layers obtained with the CAE-PVD technique
are heterogeneous, as many drop-shaped micro-sized
molecules exist in their structures. This fact leads to
changes in the mechanical, physical and chemical prop-
erties of the examined layers (Figures 1 to 6). The big-
gest surface heterogeneity, in comparison to the other ex-
amined layer surfaces, is visible within the Ti/Ti(C,N)/
(Ti,Al)N and Ti/Ti(C,N)/CrN systems, in which a lot of
precipitation of the evaporated, metal, clotty drops were
identified (Figures 1 and 2). The occurrence of this mor-
phological defect is related with the Arc-PVD process
characteristics. Depending on the process parameters, in-
cluding the kinetic energy transferred to the drops that
are crashed due to the metallic base, and the type of the
metal-vapor source used, particles varyingin shape and
size are observed. It was confirmed that the clotty drops
are spheroidal or irregular or they form agglomerates
that often include a few equal drops (Figures 1 to 6).
Moreover, characteristic hollows that form because of
the clotty drops falling out, were observed after the PVD
process wasfinished. On the basis of the metallographic
observations it was confirmed that the hollows do not
reach the surface. On the DLC layer obtained within the
PACVD process fine drops, mostly spheroidal, were
identified as well (Figure 5). The surface morphology of
the DLC layer is different from that obtained with classi-
cal high-temperature CVD processes – no micro-gaps or
wavy and globular-like surfaces were observed. The
smallest amount of morphologic surface defects was
obtained with the Cr/CrN/TiN layer (Figure 4).
A fractographic examination of magnesium- and alu-
minum-alloy samples with the applied layers, done with
a scanning electron microscope, showed a sharp transi-
T. TAÑSKI et al.: SURFACE TREATMENT OF HEAT-TREATED CAST MAGNESIUM AND ALUMINIUM ALLOYS
Materiali in tehnologije / Materials and technology 50 (2016) 5, 699–706 701
Figure 2: Surface topography of Ti/Ti(C,N)/(Ti,Al)N layer obtained
on MCMgAl6Zn1 cast magnesium alloy
Slika 2: Topografija povr{ine nanosa Ti/Ti(C,N)/(Ti,Al)N na
MCMgAl6Zn1 livni magnezijevi zlitini
Figure 1: Surface topography of Ti/Ti(C,N)/CrN layer obtained on
AlSi9Cu4 cast aluminum alloy
Slika 1: Topografija povr{ine nanosa Ti/Ti(C,N)/CrN na AlSi9Cu4
aluminijevi livni zlitini
tion zone between the base and the layer. The layers are
compact in structure with no visible delamination or de-
fects. They are placed uniformly and they adhere to the
base hermetically (Figures 7 to 10). Observations of the
fractures confirm that layers like Ti/Ti(C,N)/(Ti,Al)N
and Ti/Ti(C,N)/CrN are laminar with a visible transition
zone between the gradient and the anti-wear layers ob-
tained with different metal-vapor sources (Figure 7). On
the cross-section of the Cr/CrN/CrN, Ti/(Ti,Si)N/
(Ti,Si)N layers, in which identical sets of chemical ele-
ments of gradient and anti-wear layers were used, no vis-
ible differences were observed (Figure 8). In addition,
multi-layer carbon coats like Ti/DLC/DLC obtained with
the CVD method, in which the gradient in the middle
coating allows a variable silicon concentration, do not
exhibit any visible transition zone between individual
layers. Moreover, in the range of the thin adhesive layer
(whose task is to improve the adhesion of the base and
DLC layers) is was possible to identify a bright, continu-
ous layer of titanium, which was also confirmed with
theEDS spectrometry analysis (Figure 9). It was con-
firmed that the titanium nitride layer obtained with the
Cr/CrN/TiN system has a close to columned rise charac-
ter of crystallite that is characteristic for titanium
T. TAÑSKI et al.: SURFACE TREATMENT OF HEAT-TREATED CAST MAGNESIUM AND ALUMINIUM ALLOYS
702 Materiali in tehnologije / Materials and technology 50 (2016) 5, 699–706
Figure 4: Surface topography of Cr/CrN/TiN layer obtained on
MCMgAl9Zn1 cast magnesium alloy
Slika 4: Topografija povr{ine nanosa Cr/CrN/TiN na MCMgAl9Zn1
livni magnezijevi zlitini
Figure 6: Surface topography of Ti/(Ti,Si)N/(Ti,Si)N layer obtained
on AlSi9Cu1 cast aluminum alloy
Slika 6: Topografija povr{ine nanosa Ti/(Ti,Si)N/(Ti,Si)N na AlSi9Cu1
livni aluminijevi zlitini
Figure 3: Surface topography of Cr/CrN/CrN layer obtained on
MCMgAl9Zn1 cast magnesium alloy
Slika 3: Topografija povr{ine nanosa Cr/CrN/CrN na MCMgAl9Zn1
livni magnezijevi zlitini
Figure 7: Fracture of Ti/Ti(C,N)/CrN layer obtained on AlSi9Cu4 cast
aluminum alloy
Slika 7: Prelom nanosa Ti/Ti(C,N)/CrN na AlSi9Cu4 livni aluminijevi
zlitini
Figure 5: Surface topography of Ti/DLC/DLC layer obtained on
AlSi9Cu1 cast aluminum alloy
Slika 5: Topografija povr{ine nanosa Ti/DLC/DLC na AlSi9Cu1 livni
aluminijevi zlitini
nitride-based layers achieved with the Arc-PVD (Figure
10).
The chemical-composition investigation carried out
with GDOES and SEM confirmed the existence of the
chemical elements of the obtained layers in a depth of
1.4–3.4 μm (Figures 11 and 12). The maximum thick-
ness of the layers was measured to beas follows:
Ti/Ti(C,N)/CrN ~3.3μm; Ti/Ti(C,N)/Ti(Al,N) ~3.4μm;
Cr/CrN/CrN ~1.8μm; Cr/CrN/TiN ~1.7μm; Ti/Ti(Si,N)/
Ti(Si,N) ~1.6μm; Ti/DLC/DLC ~2.5μm. The variation in
the bonding-zone character – an increase in the chemi-
cal-element concentration of the base and a decrease in
the chemical-element concentration of the layer – leads
to the conclusion about the existence of a transitory dif-
fusion zone between the base material and the layer,
which improves their adhesion. Moreover, with the
GDOES examination, the decrease zone of the linear
concentration of the chemical elements of the layer was
T. TAÑSKI et al.: SURFACE TREATMENT OF HEAT-TREATED CAST MAGNESIUM AND ALUMINIUM ALLOYS
Materiali in tehnologije / Materials and technology 50 (2016) 5, 699–706 703
Figure 11: Variation in the concentration of Ti/Ti(C,N)/CrN layer
chemical elementsobtained on AlSi9Cu1 magnesium alloy
Slika 11: Spreminjanje koncentracije elementov v Ti/Ti(C,N)/CrN
nanosu na AlSi9Cu1 magnezijevi zlitini
Figure 9: Fracture of Ti/DLC/DLC layer obtained on AlSi9Cu1 cast
aluminum alloy
Slika 9: Prelom nanosa Ti/DLC/DLC na AlSi9Cu1 livni aluminijevi
zlitini
Figure 12: Variation in the concentration of Ti/(Ti,Si)N/(Ti,Si)N layer
chemical elements obtained on MCMgAl6Zn1magnesium alloy
Slika 12: Spreminjanje koncentracije elementov v Ti/(Ti,Si)N/
(Ti,Si)N nanosu na MCMgAl6Zn1 magnezijevi zlitini
Figure 10: Fracture of Cr/CrN/TiN layer obtained on MCMgAl9Zn1
cast magnesium alloy
Slika 10: Prelom nanosa Cr/CrN/TiN na MCMgAl9Zn1 livni mag-
nezijevi zlitini
Figure 8: Fracture of Ti/(Ti,Si)N/(Ti,Si)Nlayer obtained on
MCMgAl6Zn1 cast magnesium alloy
Slika 8: Prelom nanosa Ti/(Ti,Si)N/(Ti,Si)N na MCMgAl6Zn1 livni
magnezijevi zlitini
confirmed – it proves that the layers are gradient (Fig-
ures 11 and 12).
The layers obtained with Arc-PVD and PA CVD on
the base of Mg and Al alloys significantly increased the
microhardness compared to the base material (Figure
13). This phenomenon is caused by the chemical- and
phase-concentration change, various conditions, the type
of the method (PVD or CVD) and the combination of the
layers.
The type of base material – Mg or Al alloys – has the
least influence on the microhardness (Figure 13). A sig-
nificant rise in the microhardness after the precipitation
hardening, exceeding 100 % compared to the base mate-
rial, took place due totheCr/CrN/CrN, Cr/CrN/TiN and
Ti/(Ti,Si)N/(Ti,Si)N layers obtained with the cathodic
PVD process and N2 as the protective gas. The layers’
hardness does not exceed 2000 HV according to the
microhardness test results, while the Ti/Ti(C,N)/CrN and
Ti/Ti(C,N)/(Ti,Al)N layers obtained within the mixture
of CH4 and N2 as the protective gas are even harder than
2000 HV.
In the case of the 5N load used in the examination,
the average friction factor for theDLC coatings (done on
Al and Mg) achieved with a 0.05 m/s slide velocity is in
a range of 0.06–0.16 (Figure 15). Thisis a decrease in
the whole order of magnitude in comparison to the other
layers. This state is characteristic for the DLC layers that
consist of graphite, which works like a lubricant during
the abrasion process. Moreover, a high traverse speed,
which causes heat accumulation, is responsible for an
easier self-lubrication of the layer, resulting in a reduc-
tion in the friction coefficient (Figures 14 and 15). The
sliding distance for the DLC coatings obtained on Mg is
even 70 times higher than those measured for Cr/CrN/
CrN or diamond-like layers (Figure 14). The sliding-dis-
tance values for all the examined layers were in a range
of 6–630 m (Figure 14). When examining all the
ball-on-disk test results, it was confirmed that the value
T. TAÑSKI et al.: SURFACE TREATMENT OF HEAT-TREATED CAST MAGNESIUM AND ALUMINIUM ALLOYS
704 Materiali in tehnologije / Materials and technology 50 (2016) 5, 699–706
Figure 15: Dependence of layer friction coefficient on the coun-
ter-samplesliding distance achieved withtheball-on-disc method on: a)
Ti/(Ti(C,N)/(Ti,Al)N, b) Ti/(Ti(C,N)/CrN, c) Cr/CrN/TiN, d)
Cr/CrN/CrN, e) Ti/(Ti,Si)N/(Ti,Si)N, f) Ti/DLC/DLC, obtained on
aluminum and magnesium cast alloys
Slika 15: Odvisnost koeficienta trenja od razdalje drsenja, dobljene
pri metodi kroglica na plo{~i pri: a) Ti/(Ti(C,N)/(Ti,Al)N, b) Ti/
(Ti(C,N)/CrN, c) Cr/CrN/TiN, d) Cr/CrN/CrN, e) Ti/(Ti,Si)N/(Ti,Si)N,
f) Ti/DLC/DLC na livnih aluminijevih in magnezijevih zlitinah
Figure14: Dependence of sliding distanceto the coating damage on
the minimum and maximum friction coefficients intheball-on-disc test
used on PVD and CVD layers obtained on aluminum and magnesium
cast alloys
Slika 14: Odvisnost razdalje drsenja do po{kodbe nanosa na mini-
malni in maksimalni koeficient trenja pri preizkusu kroglica na plo{~i,
uporabljenem na PVD in CVD nanosih, na livnih aluminijevih in
magnezijevih zlitinah
Figure13: Microhardness examination results forcast magnesium
Mg-Al-Zn and aluminum Al-Si-Cu alloys after ageing, PVD and CVD
processing
Slika 13: Rezultati meritev mikrotrdote na liti magnezijevi Mg-Al-Zn
in aluminijevi Al-Si-Cu zlitini, po staranju ter PVD in CVD obdelavi
of the sliding distance for magnesium coatings is larger
than the sliding distance obtained with the alumi-
num-layer dual system; moreover, the results for the
DLC coatings are even 30 % better (Figure 14).
For the tribological-resistance examination of the ob-
tained layers, the charts demonstrating the rotation quan-
tity or displacement of the countersample before the
coating damage, depending on the friction coefficient
and/or displacement of the countersample along the ver-
tical axis were created. For all the registered friction co-
efficients dependingon the rotation quantity or the slid-
ing distance, similar characteristic curves, which can be
divided into two parts, were determined (Figure 15). In
the first part, a significant increase in the friction coeffi-
cient with the sliding-distance increase was determined.
It was accepted that it is the transient friction state. The
second part of the chart is similar to the stationary state.
Large changes in the friction coefficient measured during
the examination were caused by a spallation of the sam-
ple and counter-sample surfaces.
4 CONCLUSION
Gradient and monolithic coatings including
Ti/Ti(C,N)/CrN, Ti/Ti(C,N)/(Ti,Al)N, Ti/(Ti,Si)N/
(Ti,Si)N, Cr/CrN/CrN, Cr/CrN/TiN and Ti/DLC/DLC
were successfully deposited onto magnesium- and alumi-
num-alloy substrates using the cathodic-arc-evaporation
method (Arc-PVD) and plasma-assisted process (PA
CVD). The layers obtained with Arc-PVD and PA CVD
on the base of Mg and Al alloys significantly increased
the microhardness compared to the base material. The
type of base material had the least influence on the
microhardness. The fracture morphology was character-
ized by a lack of columnar structures in the obtained
coatings. The metallographic examinations proved that
the coatings were deposited uniformly over the whole
sample onto the investigated substrate materials and that
the measured thickness was characteristic for this coat-
ing type. It was also found that individual layers adhered
tightly to each other and to the light metal substrate. The
average friction factor of the DLC coatings (on Al and
Mg) achieved at a 0.05 m/s slide velocity was in a range
of 0.06–0.16, which was a decrease in the whole order of
magnitude in comparison to the other layers. Examining
all the ball-on-disk test results, it was confirmed that the
value of the sliding distance for the magnesium coatings
was higher than the sliding distance obtained with the
aluminum-layer dual system; moreover, the results for
the DLC coatings were even 30 % better. The investiga-
tion results obtained for the PVD method make it possi-
ble to obtain interesting solutions, which are very prom-
ising for the use in the automobile and aviation
industries, for the production of the elements whose high
functional properties like hardness and wear resistance
are crucial for long-term service.
Acknowledgements
This publication was co-financed within the frame-
work of the statutory financial grant providedby the Fac-
ulty of Mechanical Engineering of the Silesian Univer-
sity of Technology in 2015.
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706 Materiali in tehnologije / Materials and technology 50 (2016) 5, 699–706
K. GRUSZKA: ANALYSIS OF THE STRUCTURAL-DEFECT INFLUENCE ON THE MAGNETIZATION PROCESS ...
707–718
ANALYSIS OF THE STRUCTURAL-DEFECT INFLUENCE ON THE
MAGNETIZATION PROCESS IN AND ABOVE THE RAYLEIGH
REGION
ANALIZA VPLIVA STRUKTURNIH DEFEKTOV NA PROCES
MAGNETIZACIJE V IN NAD RAYLEIGH PODRO^JEM
Konrad Gruszka
Czestochowa University of Technology, Faculty of Production Engineering and Materials Technology, Institute of Physics,
Armii Krajowej Av. 19, 42-200 Czestochowa, Poland
kgruszka@wip.pcz.pl
Prejem rokopisa – received: 2015-06-30; sprejem za objavo – accepted for publication: 2015-09-16
doi:10.17222/mit.2015.154
The paper presents studies of the structural-defect influence on the magnetization process in low magnetic fields (H < 0.4 Hc)
and above the Rayleigh range. The investigated Fe62Co10Y8B20 alloy samples were obtained with the injection casting method
resulting in the amorphous-structure state, which was confirmed with XRD. The studies were conducted by analyzing the
disaccommodation of the magnetic-susceptibility process and using Kronmüller’s theory in the approach to the ferromagnetic
saturation area. On the basis of the obtained results, it was found that the main factor responsible for the processes of
magnetization in low magnetic fields are point defects, whereas in the case of high magnetic fields, the magnetization process
depends mainly on second-type pseudo-dislocation dipoles.
Keywords: metallic glasses, defects, disaccommodation, Kronmüller’s theory
^lanek predstavlja {tudij vpliva strukturnih defektov na proces magnetizacije v {ibkem magnetnem polju (H < 0,4 Hc) in nad
Rayleigh podro~jem. Vzorci preiskovane zlitine Fe62Co10Y8B20 so bili izdelani z metodo injekcijskega brizganja, kar je
povzro~ilo amorfno strukturo, ki je bila potrjena z rentgensko analizo (XRD). [tudije so bile izvedene z analizo procesa
neustreznosti magnetne ob~utljivosti in z uporabo Kronmülerjeve teorije pri pribli`evanju feromagnetno nasi~enemu podro~ju.
Na osnovi dobljenih rezultatov je bilo ugotovljeno, da so pri procesu magnetizacije glavni faktor to~kaste napake, medtem ko je
v primeru magnetizacije v mo~nem magnetnem polju proces odvisen predvsem od druge vrste psevdodislociranih dipolov.
Klju~ne besede: kovinska stekla, napake, neustreznost, Kronmüllerjeva teorija
1 INTRODUCTION
Iron-based amorphous materials are extensively stu-
died because of their excellent soft-magnetic proper-
ties.1–3 For this reason, they have been widely used in the
electrical industry, primarily as high-efficiency cores for
power transformers and chokes and also as coatings due
to their high corrosion resistance.4,5
In terms of topological structure, amorphous mate-
rials have properties similar to those of liquids. The
atoms forming the material are scattered in a chaotic
manner so that the observation of the long-range order is
not possible. When the cooling rate of a liquefied alloy is
sufficiently large, the kinetic energy of the atoms is taken
so fast that they are trapped at higher energy positions.
This results in the density and local chemical-compo-
sition fluctuations, and is the major cause of the areas
with a deficiency of atoms. This type of local voids are
called point defects (by analogy with the vacancies
occurring in a crystal structure). In the cases where
several point defects are located in a small local envi-
ronment, a concentration to two-dimensional defects
occurs and it is referred to as pseudo-dislocation dipoles.
Point defects and their conglomerates have a signifi-
cant impact on the process of magnetization in high
magnetic fields.6,7 The presence of structural defects,
which are the centers of internal stresses, causes a defor-
mation of local magnetization vectors. The presence of
structural defects in amorphous materials affects the
magnetization process in the area called the approach to
ferromagnetic saturation.8 In close proximity to a point
defect, the magnetization vectors are arranged in a
"streamline" way9,10 which is a potential cause of do-
main-wall anchors. A more complicated situation occurs
in the presence of a pseudo-dislocation dipole, where the
arrangement of individual vectors is not collinear and the
centers of the deformation of these vectors are hooked at
the dipole ends.11–13
Ferromagnetic materials, especially the ones based
on the Fe-Co-B composition are well known for their
great soft-magnetic properties, in particular a low coer-
cive field and magnetostriction while they are not expen-
sive and have a high sensitivity to alloying additives. In
the compounds of this type, iron is typically the major
part (more than a 50 % amount fraction) while Co,
which has similar properties, allows an increase in the
magnetization. Boron, due to a significant difference in
the atomic radius, improves the glass-forming ability
(GFA). In order to further increase the GFA, a small
Materiali in tehnologije / Materials and technology 50 (2016) 5, 707–718 707
UDK 620.1:67.017:621.318 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)707(2016)
amount of yttrium atoms are introduced. According to
the literature, a concentration of up to a 4 % amount
fraction has a positive effect on the GFA but exceeding
this value brakes down good soft-magnetic properties.
Using the studies on susceptibility-disaccommodation
phenomena and approach to the ferromagnetic saturation
area, it is possible to indirectly describe the structural
defects existing in a material. This paper focuses on
defects – the relations within a magnetization process
and shows that despite an yttrium addition of up to 8 %,
one can obtain reasonable results in terms of soft-mag-
netic parameters.
2 STUDIED MATERIAL AND METHODOLOGY
The investigated material comprises chemical ele-
ments of high purity (Fe – a 99.95 % amount fraction,
the remaining component elements – 99.99 % amount
fractions). A two-step preparation procedure was used.
Initially, the ingredients were melted using a plasma arc
(a working current of ~300 A) under a reduced pressure
in a protective atmosphere of argon. The samples were
melted several times to ensure a good homogeneity of
the constituent distribution. Then the resulting ingot was
melted, using an induction furnace, in a quartz tube and
injected into a copper water-cooled mold, which was also
under an argon atmosphere (at a pressure of 700 hPa). In
the radial cooling process, good-quality amorphous solid
samples with dimensions of 15 mm × 10 mm × 0.5 mm
were obtained. The structure of the obtained samples was
examined using an X-ray diffractometer (Bruker D8
Advance, Cu-K). The magnetic properties were deter-
mined with an analysis of the static hysteresis loop and
the initial magnetization curve, which were obtained
from the measurements of the magnetization as a func-
tion of the magnetic-field strength using a vibrating
sample magnetometer (model Lakeschore 7301). Measu-
rements of the initial permeability of the samples were
taken over a wide temperature range of 300–650 K. The
samples were suspended in a permalloy frame and
placed in a quartz vacuum tube with a thermocouple, and
the whole system was placed in an accumulative furnace.
The disaccommodation-aftereffect intensity defined as
Δ
1 1 1
1 0x t t
⎛
⎝
⎜ ⎞
⎠
⎟ = −
 ( ) ( )
was measured in time (t1 – t0) = 118 s
with the transformer method. The frequency of the ac
field was set to 50 Hz. The samples were demagnetized
by applying a sinusoidal decreasing magnetic field of
120 kHz.
3 RESULTS
Figure 1 shows the XRD of a sample after the solidi-
fication, measured for the 2 angle in a range from 30°
to 100°.
The diffraction pattern shows only the broad fuzzy
maximum, located in a 2 angle range from approxima-
tely 35° to 50°. This means that in a sample in the
as-quenched state, there is no long-range order and the
interactions between the atoms have a close- or me-
dium-range character. Such a shape of a diffraction
pattern is typical for the materials with an amorphous
structure. Then measurements of the static magnetic
hysteresis loop were carried out and the result is shown
in Figure 2.
The observed shape of the curve is characteristic for a
magnetic material with soft magnetic properties. Based
on the analysis of the static magnetic hysteresis loop, the
basic parameters for the magnetic sample were deter-
mined: magnetization saturation Ms = 1.27 (T), coercive
field Hc = 289 (A/m). Next, an analysis of the primary
magnetization curve in the approach to the ferromagnetic
saturation area was carried out. Figures 3 and 4 show
magnetization curves as a function of (μ0H)–2 and
(μ0H)1/2.
With respect to Kronmüller’s theory, as a result of the
analysis of the primary magnetization curve, it was
found that in a Fe62Co10Y8B20 alloy sample the magne-
tization process in the area called "Ewing’s knee" relates
to the change in magnetization vector directions caused
by the presence of free-volume conglomerates. The
linear-defect size (Ddip) must therefore be larger than the
exchange distance. Assuming that Ddip > lH, the a2/(μ0H)2
law of approach to the ferromagnetic saturation area is
K. GRUSZKA: ANALYSIS OF THE STRUCTURAL-DEFECT INFLUENCE ON THE MAGNETIZATION PROCESS ...
708 Materiali in tehnologije / Materials and technology 50 (2016) 5, 707–718
Figure 2: Static magnetic hysteresis loop for the Fe62Co10Y8B20 alloy
Slika 2: Stati~na magnetna histerezna zanka zlitine Fe62Co10Y8B20
Figure 1: X-ray diffraction pattern of Fe62Co10Y8B20 in the
as-quenched state
Slika 1: Rentgenska difrakcija Fe62Co10Y8B20 v kaljenem stanju
fulfilled. Thus, these defects are the main source of stress
in the magnetic-field strength from 0.084 (T) to 0.25 (T).
Above the (μ0H)–2 dependence, Holstein-Primakoff’s
process intensifies, indicating that the further process of
magnetization is associated with the damping of ther-
mally excited spin waves with an external magnetic field
(the b coefficient). The parameters obtained from the
analysis of the primary magnetization curve are summar-
ized in Table 1.
Table 1: Parameters obtained with the analysis of the primary mag-
netization curve
Tabela 1: Parametri, dobljeni z analizo primarne krivulje magnetiza-
cije
a2
(10–2 T2)
b
(10–2 T1/2)
Dsp
(10–2 meV nm2)
Aex
(10–12 J m–1)
lh
(nm)
0.056 6.59 40.70 1.64 3.59
The absence of the a1/2 and a1 factors indicate that the
point defects and linear defects of a smaller size had no
effect on the magnetization process in magnetic fields
greater than 0.4 Hc (above the Rayleigh region). This
does not mean, however, that those defects are not pre-
sent in the sample’s structure but only that they are insig-
nificant for the magnetization process. To investigate
their potential impact on the magnetization at low fields,
the susceptibility disaccommodation studies were con-
ducted. According to Kronmüller’s theory, it is possible
to calculate the exchange distance of pseudo-dislocation
dipoles (the lh parameter) and the exchange-constant
(Aex) parameter. The latter is responsible for transferring
magnetic interactions between the nearest neighbors (the
energy of aligned spins) and the former describes the
size of a defect’s influence zone. The Dsp parameter,
which describes the spin-wave stiffness (and, therefore,
the ability to transfer the spin torque) is more than
twelve times higher (Dsp/DFe = 12.96) than with pure Fe
(the largest percentage share in the alloy), which is found
to be between DFe = 2.8 meV nm2 at 4.2 K14 and DFe =
3.14 meV nm2 at room temperature.15 This parameter is
connected with the atomic packing density (a higher Dsp
means a higher surface density16) and it may indicate that
linear defects should rather be considered as swellings of
voids resulting in an increase in the local density around
the defects, while keeping the overall material density
low.
The thermal stability of the initial magnetic suscep-
tibility was also investigated. This parameter is very
important and often determines possible applications.
For the studied sample, the shape of the dependency is
shown in Figure 5.
Quite minor temperature-related changes in the value
of the initial magnetic susceptibility were observed. A
good stability and an almost linear growth may indicate
that, in magnetic terms, the material is homogeneous and
no other magnetic phases were formed (in accordance to
K. GRUSZKA: ANALYSIS OF THE STRUCTURAL-DEFECT INFLUENCE ON THE MAGNETIZATION PROCESS ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 707–718 709
Figure 5: Dependence of the low-field magnetic susceptibility as a
function of temperature for the Fe62Co10Y8B20 sample
Slika 5: Odvisnost magnetne ob~utljivosti v {ibkem polju od tempera-
ture vzorca Fe62Co10Y8B20
Figure 3: Magnetization as a function of (μ0H)–2 for the
Fe62Co10Y8B20 alloy
Slika 3: Magnetizacija kot funkcija (μ0H)–2 zlitine Fe62Co10Y8B20
Figure 4: Magnetization as a function of (μ0H)1/2 for the
Fe62Co10Y8B20 alloy. Holstein-Primakoff process
Slika 4: Magnetizacija kot funkcija (μ0H)1/2 zlitine Fe62Co10Y8B20.
Proces Holstein-Primakoff
the XRD). Considering the rather low temperatures
region, the conspicuous increase between 325 K and 400 K
can be elucidated mainly with the sample’s internal-
tension relaxation processes, together with a minor con-
tribution of the magnetic-anisotropy reduction. The sharp
decline observed near 600 K is associated with the
paramagnetic transition when the ferromagnetic order is
destroyed.
The time dependence of the initial susceptibility after
the demagnetization at various temperatures, or the
magnetic-susceptibility disaccommodation curve for the
investigated alloy, is presented in Figure 6.
As can be seen in Figure 6, the disaccommodation
intensity linearly rises up to the broad maximum located
at about 542 K. Above 542 K, the intensity decreases,
which is directly connected with the decrease in the
amount of the atom pairs reorienting in the point-defect
vicinity. A pair reorientation can occur in two cases: the
first one is associated with a reversible process, in which,
through energy delivery, atom pairs near the point
defects can swap their positions and return to the initial
state after the energy reduction, and in the second one,
this displacement is permanent due to the major changes
in the local space involving a minimum of three atoms.
Both processes can occur in the same time and the curve
observed in Figure 6 indicates the effect of the impo-
sition of these two phenomena. The visible kink in the
same temperature region as in the case of the stability of
the initial-susceptibility curve (Figure 5) is presumably
also caused by the relaxation processes. This phenome-
non must therefore at least partially occur through the
irreversible atomic-pair reorientation, clearly manifested
between 325 K and 400 K (Figure 6). As the relaxation
process of the material has to occur through the displace-
ment of atoms into positions that lead to a lower total
energy of the system, a part of these movements must
therefore lead to the reorientation of atomic pairs visible
as disaccommodation phenomena. Next, a numerical
analysis of the disaccommodation curve was done (Fig-
ure 7).
The isochronal magnetic-susceptibility disaccommo-
dation curve was numerically analyzed using the depen-
dence given with Equation (1):
Δ
1
3
3
1
1
x
I T
T
e
i
i
mi
i
l
pi pi
t
⎛
⎝
⎜ ⎞
⎠
⎟ = ⋅
⋅
−=
−
∫∑  

  



e
t
e
z
z
mi
z
ie e z−
⎛
⎝
⎜
⎞
⎠
⎟
− −
⎛
⎝
⎜⎜
⎞
⎠
⎟⎟
2
2
   d
(1)
where the mean relaxation time mi is given:
 mi
mi
pi
Q
k T T
= − −
⎛
⎝
⎜
⎞
⎠
⎟
⎛
⎝
⎜⎜
⎞
⎠
⎟⎟exp
1 1
 – the pre-exponential factor in the Arrhenius law, Ipi –
the intensity of the ith process at temperature Tpi, Qmi –
the mean activation energy,  – the distribution width
and z mi= ln /  .
The results of the analysis of the disaccommodation
curve in function of temperature, based on the above
formula done using the least squares method is presented
in Table 2.
Table 2: Results obtained from the numerical analysis of dis-
accommodation curve
Tabela 2: Rezultati, dobljeni iz numeri~ne analize neustrezne krivulje
a 1.095*10–7  (s)
b –2.230*10–5
Tp (K) 433.20 2.62×10-15
Ip 5.223*10–6
Q (eV) 1.378
 6.9368
Tp (K) 479.17 2.98×10-15
Ip 3.215*10–5
Q (eV) 1.5189
 5.0638
Tp (K) 542.87 5.26×10-15
Ip 2.52*10–5
Q (eV) 1.6942
 3.7668
The isochronal magnetic-susceptibility disaccommo-
dation curve was decomposed into three elementary
K. GRUSZKA: ANALYSIS OF THE STRUCTURAL-DEFECT INFLUENCE ON THE MAGNETIZATION PROCESS ...
710 Materiali in tehnologije / Materials and technology 50 (2016) 5, 707–718
Figure 7: Theoretical magnetic-susceptibility disaccommodation
curve (solid line) and experimental data for the Fe62Co10Y8B20 alloy
after solidification
Slika 7: Teoreti~na krivulja neprimerne magnetne ob~utljivosti (polna
linija) in eksperimentalni podatki za zlitino Fe62Co10Y8B20 po strje-
vanju
Figure 6: Magnetic-susceptibility disaccommodation curve for the
Fe62Co10Y8B20 alloy after solidification
Slika 6: Neprimerna krivulja magnetne ob~utljivosti zlitine
Fe62Co10Y8B20 po strjevanju
processes (peaks). The first peak with the maximum
localized at 433 K has the lowest activation energy (Q),
the highest width () and the highest intensity (Ip). Most
of the elementary-pair-reconfiguration processes there-
fore require a minimum activation energy of 1.38 eV. An
analysis of Table 2 reveals that the peak maximum
temperature (Tp) increases together with the activation
energy and relaxation time . At the same time, the
distribution width and process intensity decrease. This
phenomena is probably related with the atomic mass
(and also the radius) of the ingredients. Thus, atom pairs
with a higher mass or a bigger size require more energy
and time for the relaxation to occur. On this basis, it can
be concluded that the first elementary process is mostly
caused by boron (a radius of 82 pm, a weight of 10.8 u),
the second process mainly involves Fe (r. 126 pm, w.
55.9 u) and Co (r. 125 pm, w. 58.7 u) and the last pro-
cess, which requires the largest activation energy, is
caused due to an increasing involvement of yttrium (r.
180 pm, w. 88.9 u) atoms. Obviously, due to the smallest
size and weight of B, this element probably has the
largest share in each of the elementary processes, but it
should be noted that the pair reorientation involving
atoms with extreme size differences, tends to be
irreversible.
4 CONCLUSIONS
In the radial cooling process, a good-quality bulk
amorphous sample of the Fe62Co10Y8B20 composition was
prepared. Studies of its magnetic properties showed that
despite a notably large yttrium share, it is possible to
make the material to be rather soft (Hc = 289 A/m).
Susceptibility studies showed that in terms of magnetic
properties, the material is homogeneous. Simultaneously,
magnetic-susceptibility disaccommodation studies clear-
ly revealed a presence of point defects, which are res-
ponsible for the magnetization process in the Rayleigh
area. Therefore, point defects must also be distributed
uniformly across the entire volume of the material. The
decomposition of the susceptibility disaccommodation
curve in three elementary processes was sufficient to
completely describe the pair reorientation, showing the
dependence of the activation energy, intensity and rela-
xation time on the mass (and radius) of the elements
involved in the reorientation process. At H > 0.4 Hc, in
the approach to the ferromagnetic saturation area, the
initial-magnetization-curve analysis, with respect to
Kronmüller’s theory, led to the conclusion that in the
magnetization process, linear defects of the second type
play the main role.
A detailed analysis of the parameters obtained on the
basis of this theory showed that the defects of this type
can be considered as swellings of voids, leading to an
increase in the local density around them. On the other
hand, as it is known, long-term annealing below the
crystallization temperature leads to a decrease in dis-
accommodation phenomena17 associated with the stress
relaxation (due to the atom expansion) and defect diffu-
sion to the sample’s surface. There are no indications
that the linear defects are not subjected to the same
processes, leading to their decomposition into smaller
point defects. At the same time, no increase in the
intensity of disaccommodation phenomena is observed
(directly related to the amount of point defects). This
may suggest that the linear conglomerates of defects
disappear because of collapse-like processes and not due
to shredding. This is consistent with the conclusion
about an increase in the structural stresses around the
linear defects, pushing to fill the voids and, therefore,
reduce the system energy. Considering that the range of
the exchange interaction is smaller than the average size
of a linear defect, the stress relaxation due to the anneal-
ing process should lead to a reduction in the number of
defects, observed as a decrease in the Dsp parameter.
5 REFERENCES
1 M. G. Nabia³ek, P. Pietrusiewicz, M. J. Dospia³, M. Szota, K. B³och,
K. Gruszka, K. OŸga, S. Garus, Effect of manufacturing method on
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Gruszka, A. Dobrzañska-Danikiewicz, S. Walters, A. J. Bukowska,
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712 Materiali in tehnologije / Materials and technology 50 (2016) 5, 707–718
A. GRAJCAR, A. PLACHCIÑSKA: EFFECT OF SULPHIDE INCLUSIONS ON THE PITTING-CORROSION BEHAVIOUR ...
713–718
EFFECT OF SULPHIDE INCLUSIONS ON THE
PITTING-CORROSION BEHAVIOUR OF HIGH-Mn STEELS IN
CHLORIDE AND ALKALINE SOLUTIONS
VPLIV SULFIDNIH VKLJU^KOV NA JAMI^ASTO KOROZIJO
JEKEL Z VISOKO VSEBNOSTJO Mn V RAZTOPINAH KLORIDOV
IN ALKALIJ
Adam Grajcar, Aleksandra P³achciñska
Silesian University of Technology, Institute of Engineering Materials and Biomaterials, Konarskiego Street 18a, 44-100 Gliwice, Poland
adam.grajcar@polsl.pl
Prejem rokopisa – received: 2015-07-01; sprejem za objavo – accepted for publication: 2015-09-02
doi:10.17222/mit.2015.169
The corrosion behaviour of the 27Mn-4Si-2Al- and 26Mn-3Si-3Al-type austenitic steels were evaluated in chloride 3.5 % NaCl
and alkaline 0.1-M NaOH environments using potentiodynamic polarization tests. The type of non-metallic inclusions and their
pitting-corrosion behaviour were investigated. In the chloride solution, both the steels exhibited a lower corrosion resistance in
comparison to the alkaline solution. The high-Mn steels showed evidence of pitting and uniform corrosion, both in the chloride
and alkaline solutions. SEM micrographs revealed that the corrosion pits are characterized by various shapes and an irregular
distribution at the metallic matrix. Corrosion damage is more numerous in the chloride solution than in the alkaline solution.
EDS analyses revealed that the corrosion pits nucleated on MnS inclusions or complex oxysulphides. The chemical composition
of the steels (change in the Al and Si contents) does not affect the privileged areas of pit nucleation, whereas it influences the
electrochemical behaviour of the steels in the chloride solution.
Keywords: high-Mn steel, austenitic steel, non-metallic inclusion, corrosion resistance, pitting corrosion, potentiodynamic pola-
rization test
Korozijsko obna{anje avstenitnih jekel 27Mn-4Si-2Al in 26Mn-3Si-3Al je bilo ocenjeno v raztopini 3,5 % NaCl in v alkalni
raztopini 0,1 M NaOH, s pomo~jo potenciodinami~nih polarizacijskih preizkusov. Preiskovana je bila vrsta nekovinskih
vklju~kov in njihovo pona{anje pri jami~asti koroziji. V raztopini kloridov sta obe jekli, v primerjavi z alkalno raztopino, poka-
zali manj{o korozijsko obstojnost. Jekla z visoko vsebnostjo Mn so pokazala jami~asto in splo{no korozijo v obeh raztopinah,
tako v kloridni kot v alkali~ni. SEM-posnetki so pokazali korozijske jamice razli~nih oblik in njihovo neenakomerno razpo-
reditev po kovinski osnovi. Korozijske po{kodbe so bolj {tevilne v kloridni raztopini kot pa v alkalni raztopini. EDS-analize so
pokazale, da korozijske jamice nastajajo na vklju~kih MnS ali na kompleksnih oksisulfidih. Kemijska sestava jekel (spremembe
v vsebnosti Al in Si) ni vplivala na prednostna mesta nukleacije jamic, medtem ko je vplivala na elektrokemijsko pona{anje
jekel v raztopini kloridov.
Klju~ne besede: jekla z veliko vsebnostjo Mn, avstenitno jeklo, nekovinski vklju~ki, odpornost na korozijo, jami~asta korozija,
potenciodinami~ni polarizacijski preizkus
1 INTRODUCTION
Pitting corrosion is a type of localized corrosion. It
occurs mainly in the passive state of metals, in environ-
ments containing aggressive ions, i.e., chloride anions.
The pits are often invisible during the formation stage,
but their progressive local damage can lead to an element
perforation.1 It is well known that various factors – the
chemical composition, microstructure, heat treatment
and plastic deformation – affect the pitting potential.2–4
There are many reports that confirm the negative
impact of non-metallic inclusions on the corrosion resis-
tance of steel.5–7 High-manganese austenitic steels have
different types of inclusions, which form during melting
and casting. These steels contain Mn, which combines
with sulphur, and Si and Al additions with a high che-
mical affinity for oxygen (Al also to nitrogen).8–10 There-
fore, the presence of various sulphide and oxide inclu-
sions in these steels can be expected.11,12 I. J. Park et al.7
observed MnS, AlN, Al2O3, MnAl2O4 and other complex
inclusions in Mn-Al steels. ^. Donik et al.5 and A. Pardo
et al.6 reported that the Mn additions to stainless steels
have a detrimental effect on the pitting-corrosion
resistance in a NaCl medium. Manganese favours the
formation of MnS inclusions, which are vulnerable to the
initiation of corrosion pits. Moreover, its presence
drastically increases the corrosion current density of
steel and displaces the Ecorr values towards less noble
potentials. K. J. Park and H. S. Kwon13 found that the
size of the MnS inclusions increased with an increase in
the Mn concentration in Fe–18Cr–6Mn and
Fe-18Cr-12Mn steels. The shape, composition and
distribution of inclusions have significant effects on the
corrosion resistance too.
The high-manganese alloys belong to a new, ad-
vanced group of steels that combine successively high
strength and high plasticity due to the austenitic micro-
structure. Because of the homogeneous ductile micro-
Materiali in tehnologije / Materials and technology 50 (2016) 5, 713–718 713
UDK 67.017:549.3:620.193 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)713(2016)
structure they can be used for numerous elements of the
energy-absorbing structures of cars.14,15 These steels are
used as a cheaper substitute for austenitic stainless steels.
Manganese (austenite stabilizer) and aluminium can
replace expensive nickel and chromium additions. The
potential applications also include construction materials
to transport different liquid gases with various pH values.
However, their real application also depends on the co-
rrosion behaviour. Therefore, the effect of non-metallic
inclusions on the corrosion properties of the 27Mn-4Si-
2Al and 26Mn-3Si-3Al steels in two environments, i.e.,
3.5 % NaCl (neutral) and 0.1-M NaOH (alkaline), have
been investigated in this study using electrochemical
polarization tests.
2 EXPERIMENTAL PROCEDURE
The investigated materials were high-Mn austenitic
steels with the chemical composition shown in Table 1.
Both steels were treated using the same conditions. The
steel ingots were prepared by vacuum melting, then they
were hot-forged and roughly rolled to a thickness of
4.5 mm. The next step was their thermomechanical
processing, consisting of the hot rolling of flat samples
in three passes (relative reductions: 25 %, 25 % and
20 %) to a final sheet thickness of approximately 2 mm,
obtained at 850 °C. Subsequently, the samples were
rapidly cooled in water to room temperature.
The flat samples of the 27Mn-4Si-2Al and 26Mn-
3Si-3Al steels with a 0.38 cm2 exposed surface area were
prepared for the electrochemical tests in 3.5 % NaCl
(neutral) and 0.1-M NaOH (alkaline) solutions. The
samples were mechanically ground with SiC paper up to
1200 grit. Prior to the experiments, all the samples were
washed in distilled water and rinsed in acetone. The
solutions were prepared using deionised water. The
electrochemical cell comprised three electrodes. A stain-
less steel and a silver/silver chloride (Ag/AgCl) electrode
(SSE) were used as the counter and the reference
electrodes, respectively. The electrochemical measure-
ments were performed using an Atlas 0531
Electrochemical Unit potentiostat/galvanostat driven by
AtlasCorr05 software. The potentiodynamic polarisation
measurements were conducted at a scan rate of 1 mV/s.
The potentiodynamic scan data were collected to deter-
mine the electrochemical parameters: corrosion potential
Ecorr and corrosion current density Icorr.
The samples after the corrosion tests were polished
using Al2O3 with a granularity of 0.1 μm for the scanning
electron microscopy (SEM). To reveal the corrosion pits
the cover formed on the pit surface has to be removed.
Thus, the samples’ surfaces were polished to obtain a
uniform surface with pit-initiation sites. The corrosion
damage was examined based on SEM observations and
EDS techniques. Additionally, the depth of the corrosion
damage on the cross-sectioned specimens was evaluated
using a light microscope.
3 RESULTS AND DISCUSSION
Typical microstructures of the 27Mn-4Si-2Al and
26Mn-3Si-3Al steel specimens are shown in Figures 1a
and 1b, respectively. Both micrographs exhibit relatively
coarse austenite grains elongated according to the direc-
tion of hot rolling. The mean grain size is approximately
80 μm. The microstructures reveal the presence of
annealing twins, deformation effects and elongated
sulphide inclusions.
Potentiodynamic curves of the 27Mn-4Si-2Al and
26Mn-3Si-3Al steels registered in 3.5 % NaCl (neutral –
pH 7) and 0.1-M NaOH (alkaline – pH 14) solutions are
illustrated in Figures 2 and 3. The average calculated
values of the corrosion potential Ecorr and the corrosion
current density Icorr determined by the Tafel extrapolation
are shown in Table 2. Both steels show lower corrosion
resistance in the 3.5 % NaCl solution than in 0.1-M
NaOH solution. The corrosion current density registered
in the chloride solution was higher in comparison to the
alkaline solution (Table 2). The obtained data are
supported by the similar results of other authors16,17, who
reported that the high-Mn austenitic steels show a lower
A. GRAJCAR, A. PLACHCIÑSKA: EFFECT OF SULPHIDE INCLUSIONS ON THE PITTING-CORROSION BEHAVIOUR ...
714 Materiali in tehnologije / Materials and technology 50 (2016) 5, 713–718
Table 1: Chemical composition of investigated steels in mass fractions (w/%)
Tabela 1: Kemijska sestava preiskovanih jekel v masnih dele`ih (w/%)
Grade Mn Si Al S P Nb Ti N O Fe
27Mn-4Si-2Al 27.5 4.18 1.69 0.017 0.004 0.033 0.010 0.0028 0.0006 bal.
26Mn-3Si-3Al 26.0 3.08 2.87 0.013 0.002 0.034 0.010 0.0028 0.0006 bal.
Table 2: Average values of the electrochemical polarization data for the 27Mn-4Si-2Al and 26Mn-3Si-3Al steels obtained in the 3.5 % NaCl and
0.1-M NaOH solutions
Tabela 2: Srednje vrednosti podatkov elektrokemijske polarizacije 27Mn-4Si-2Al in 26Mn-3Si-3Al jekel, dobljene v raztopinah 3,5 % NaCl in
0,1 M NaOH
Grade Statistics
3.5 % NaCl 0.1 M NaOH
Ecorr/(mV) Icorr/(mA/cm2) Ecorr/(mV) Icorr/(mA/cm2)
27Mn-4Si-2Al
average value –788 0.090 –392 0.007
standard deviation 6.5 0.007 3.2 0.002
26Mn-3Si-3Al
average value –785 0.009 –395 0.005
standard deviation 11.2 0.005 4.3 0.003
corrosion resistance in the chloride medium than in the
alkaline solution.
In the 3.5 % NaCl solution, the 27Mn-4Si-2Al steel
specimens showed a much higher corrosion current
density (0.09 mA/cm2) than the 26Mn-3Si-3Al steel
(0.009 mA/cm2). This confirms our earlier results from
the potentiodynamic polarisation tests.3 It is related to
the higher Al and lower Si contents in 26Mn-3Si-3Al
steel in comparison to the steel containing 2 % Al (Table
1). It is reported18 that a silicon addition decreases the
corrosion resistance of steel. On the other hand, alumi-
nium improves the corrosion resistance due to its tenden-
cy to form a protective Al2O3 passive layer on the steel
surface in solutions of pH ~7 (Pourbaix diagrams).19 All
the specimens polarized in the 3.5 % NaCl solution show
Ecorr values shifted to less noble potentials (Table 2)
when compared to the specimens polarized in the 0.1-M
NaOH solution. The Ecorr shift was about 400 mV
towards the cathodic direction. The values of the corro-
sion-current density obtained in the 0.1-M NaOH were
quite similar for both steels. The 27Mn-4Si-2Al steel
specimens showed a corrosion current density of
approximately 0.007 mA/cm2, whereas it was 0.005
mA/cm2 for the second steel. The better corrosion
resistance of both steels in 0.1-M NaOH is related to the
fact that in alkaline solutions, manganese precipitates as
Mn(OH)2, which is slightly soluble in solutions with
pH>13, whereas in solutions of pH ~ 7 the manganese
dissolves as Mn2+ (Pourbaix diagrams).19
The morphology of the corrosion pits after the
electrochemical tests were studied using the SEM and
A. GRAJCAR, A. PLACHCIÑSKA: EFFECT OF SULPHIDE INCLUSIONS ON THE PITTING-CORROSION BEHAVIOUR ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 713–718 715
Figure 3: Potentiodynamic polarization curves of the: a) 27Mn-4Si-
2Al and b) 26Mn-3Si-3Al steels obtained in 0.1-M NaOH solution
Slika 3: Krivulje potenciodinami~ne polarizacije jekel v raztopini
0,1 M NaOH: a) 27Mn-4Si-2Al in b) 26Mn-3Si-3Al
Figure 2: Potentiodynamic polarization curves of the: a) 27Mn-
4Si-2Al and b) 26Mn-3Si-3Al steels obtained in 3.5 % NaCl solution
Slika 2: Krivulje potenciodinami~ne polarizacije jekel v raztopini
3,5 % NaCl: a) 27Mn-4Si-2Al in b) 26Mn-3Si-3Al
Figure 1: Austenitic microstructure of the thermomechanically
processed: a) 27Mn-4Si-2Al and b) 26Mn-3Si-3Al steels
Slika 1: Avstenitna mikrostruktura termomehansko izdelanih jekel:
a) 27Mn-4Si-2Al in b) 26Mn-3Si-3Al
EDS techniques. The SEM images of the corrosion
damage in both steels after the corrosion tests in 3.5 %
NaCl are shown in Figures 4 and 5. The corrosion pits
are characterized by various shapes and an irregular
distribution at the metallic matrix (Figures 4a and 5a).
They are formed both at the grain boundaries and within
the austenite grains. It is apparent that the pits are
initiated at non-metallic inclusions. The EDS analysis
revealed the variation of the chemical composition in the
interior of the individual pits. For instance, the chemical
composition of the particle inside the corrosion pit in
Figure 4b showed a high content of manganese and
sulphur (Figure 4d). This indicates that the privileged
A. GRAJCAR, A. PLACHCIÑSKA: EFFECT OF SULPHIDE INCLUSIONS ON THE PITTING-CORROSION BEHAVIOUR ...
716 Materiali in tehnologije / Materials and technology 50 (2016) 5, 713–718
Figure 5: a) SEM micrograph of the 26Mn-3Si-3Al steel surface,
b) the individual pit interior, c) EDS analysis from point C after
corrosion test in 3.5 % NaCl
Slika 5: a) SEM-posnetek povr{ine jekla 26Mn-3Si-3Al, b) izgled
posamezne jamice, c) EDS-analiza to~ke C po korozijskem preizkusu
v 3,5 % NaCl
Figure 4: a) SEM micrograph of the 27Mn-4Si-2Al steel surface, b)
the individual pit interior, c) EDS analysis from point D, d) EDS
analysis from point C after corrosion test in 3.5 % NaCl
Slika 4: a) SEM-posnetki povr{ine jekla 27Mn-4Si-2Al, b) izgled
posamezne jamice, c) EDS-analiza v to~ki D, d) EDS-analiza to~ke C
po korozijskem preizkusu v 3,5 % NaCl
places for the pit initiation are MnS inclusions. There are
many reports in the literature5-7,20 that confirm that MnS
inclusions are vulnerable for the initiation of corrosion
pits. Their presence increases the corrosion current
density and displaces the Ecorr values towards less noble
potentials. The chemical analysis of the corrosion
damage shown in Figure 4b also revealed the presence
of oxides containing Al and Si (Figure 4c).
Similar results were obtained for corrosion pits
created in the steel containing the higher Al content
(26Mn-3Si-3Al steel). The pits are preferentially ini-
tiated along the grain boundaries (Figure 5a). The EDS
analysis showed the presence of corrosion pits at parti-
cles with the high concentrations of manganese and
sulphur, too (Figures 5b and 5c). The resistance to
pitting corrosion strongly depends on the quantity, size
and type of non-metallic inclusions in the metallic
matrix.11 Park et al.7 found that the size of the MnS
inclusions increased with an increase in the Mn content
from 6 to 12 % in the high-Cr steel. This is why both
high-Mn steels contain a lot of corrosion damage.
The SEM images of both high-Mn steels after the
corrosion tests in 0.1-M NaOH show good agreement
with the results of the potentiodynamic tests. The obser-
vation of the steel surfaces (Figure 6a) confirmed a
substantial reduction in the amount of corrosion damage.
The EDS analyses of the individual pit in Figures 6b and
6c revealed a high content of Mn, S, Al, Si and O, which
indicates that complex oxysulphides containing Mn, Al
and Si are also preferential sites for pit formation in the
alkaline solution. According to I. J. Park et al.21 Al2O3
particles have a higher resistance to pit formation than
MnS particles.
The nature of the corrosion damage was evaluated on
cross-sectioned specimens. In the chloride solution, the
specimens showed evidence of uniform corrosion. In
addition to uniform corrosion, pitting corrosion was also
A. GRAJCAR, A. PLACHCIÑSKA: EFFECT OF SULPHIDE INCLUSIONS ON THE PITTING-CORROSION BEHAVIOUR ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 713–718 717
Figure 7: Light micrographs of the cross-section of 27Mn-4Si-2Al
steel potentiodynamically polarized in: a) chloride solution and b)
alkaline solution
Slika 7: Posnetek preseka jekla 27Mn-4Si-2Al potenciodinami~no
polariziranega v: a) kloridni raztopini in b) alkalni raztopini
Figure 6: a) SEM micrograph of the 27Mn-4Si-2Al steel surface,
b) the individual pit interior, c) EDS analysis from point C after
corrosion test in 0.1-M NaOH
Slika 6: a) SEM-posnetek povr{ine jekla 27Mn-4Si-2Al, b) izgled
posamezne jamice, c) EDS-analiza iz to~ke C po korozijskem preiz-
kusu v 0,1 M NaOH
observed (Figure 7). Other authors5–7,16 observed corro-
sion pits in different high-manganese steels after polari-
zation tests in chloride solution too. After the corrosion
tests in 3.5 % NaCl the maximum depth of the corrosion
pits in the steel containing 2 % Al was evaluated to be
15 μm (Figure 7a). Similar corrosion pits were also
identified for the 26Mn-3Si-3Al steel. The corrosion
damage formed in both steels in the alkaline medium is
characterized by a small depth of 5 μm (Figure 7b).
However, the quantity of corrosion damage was much
lower when compared to the samples investigated in the
chloride medium.
4 CONCLUSIONS
The morphology of corrosion damage after electro-
chemical tests in 3.5 % NaCl and 0.1-M NaOH supports
the data registered in potentiodynamic tests. The
high-Mn steels are characterized by a low corrosion
resistance, especially in a chloride solution, where
corrosion damage is more numerous. The corrosion pits
are characterized by various shapes and are distributed
irregularly both at grain boundaries and within the
grains. EDS analyses confirmed that the corrosion pits
nucleated preferentially on the MnS inclusions and
complex oxysulphides containing Mn, Al and Si. The
low density of corrosion damage in the alkaline solution
is related to the fact that Mn precipitates as Mn(OH)2,
which is slightly soluble in solutions of pH>13, whereas
in solutions of pH ~ 7 the manganese dissolves as Mn2+.
The concentration of the individual alloying elements
was not strongly related to the corrosion behaviour of the
steels in 0.1-M NaOH, in contrast to the 3.5 % NaCl
solution. The increased contents of Mn and Si and the
smaller content of Al are reflected in the lower corrosion
resistance of the 27Mn-4Si-2Al, as registered during the
potentiodynamic tests.
Acknowledgment
This work was financially supported with statutory
funds of the Faculty of Mechanical Engineering of the
Silesian University of Technology in 2015.
5 REFERENCES
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hot-working conditions on a structure of high-manganese steel,
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steels in the context of dominant stress mechanism, Defect and
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C-Mn-Si-Al-Nb high-manganese steel during the thermomechanical
processing, Materials Science Forum, 638-642 (2010), 3224–3229,
doi:10.4028/MSF.638-642.3224
11 A. Grajcar, U. Galisz, L. Bulkowski, Non-metallic inclusions in. high
manganese austenitic alloys, Archives of Materials Science and
Engineering, 50 (2011), 21–30
12 A. Grajcar, M. Ró¿añski, M. Kamiñska, B. Grzegorczyk, Study on
non-metallic inclusions in laser-welded TRIP-aided Nb microalloyed
steel, Archives of Metallurgy and Materials, 59 (2014) 3, 1163–1169,
doi:10.2478/amm-2014-0203
13 K. J. Park, H. S. Kwon, Effects of Mn on the localized corrosion
behavior of Fe–18Cr alloys, Electrochimica Acta, 55 (2010),
3421–3427, doi:10.1016/j.electacta.2010.01.006
14 M. Jab³oñska, G. Niewielski, R. Kawalla, High manganese TWIP
steel – technological plasticity and selected properties, Solid State
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on the corrosion properties of a Fe–Mn–Al–Si steel and an
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A. GRAJCAR, A. PLACHCIÑSKA: EFFECT OF SULPHIDE INCLUSIONS ON THE PITTING-CORROSION BEHAVIOUR ...
718 Materiali in tehnologije / Materials and technology 50 (2016) 5, 713–718
A. AYDAY: INFLUENCE OF Na2SiF6 ON THE SURFACE MORPHOLOGY AND CORROSION RESISTANCE ...
719–722
INFLUENCE OF Na2SiF6 ON THE SURFACE MORPHOLOGY AND
CORROSION RESISTANCE OF AN AM60 MAGNESIUM ALLOY
COATED BY MICRO ARC OXIDATION
VPLIV Na2SiF6 NA MORFOLOGIJO POVRŠINE IN KOROZIJSKO
ODPORNOST MAGNEZIJEVE ZLITINE AM60, PREKRITE Z
MIKROOBLO^NO OKSIDACIJO
Aysun Ayday
Sakarya University, Faculty of Engineering, Department of Metallurgical and Materials Engineering, 54187 Sakarya, Turkey
aayday@sakarya.edu.tr
Prejem rokopisa – received: 2015-07-01; sprejem za objavo – accepted for publication: 2015-09-09
doi:10.17222/mit.2015.176
Oxide coatings were formed by micro arc oxidation (MAO) on an AM60 magnesium alloy substrate. The effects of Na2SiF6 in
an electrolytic solution on the micro arc oxidation process and the structure and mechanical properties of the oxide coatings
were investigated. The results showed that the MAO coating produced in the electrolyte with Na2SiF6 was thicker and more uni-
form than that produced in the electrolyte without Na2SiF6. The pore diameter of the MAO coatings was reduced by the addition
of Na2SiF6, while the coating density and surface roughness were increased. The coating formed in the electrolytic solution with
or without the Na2SiF6 had a higher surface hardness than the AM60 alloy and the results of the corrosion behavior for including
Na2SiF6 showed better resistance than that formed in the solution without Na2SiF6.
Keywords: magnesium alloy, micro arc oxidation (MAO), Na2SiF6, corrosion
Oksidne prevleke nastajajo pri oksidaciji v mikroobloku (MAO) podlage iz magnezijeve zlitine AM60. Preiskovan je bil vpliv
Na2SiF6 v elektrolitni raztopini na proces oksidacije v mikroobloku in na mehanske lastnosti oksidne prevleke. Rezultati so
pokazali, da je oksidna prevleka MAO, izdelana v elektrolitu z Na2SiF6, debelej{a in bolj enakomerna, kot ~e je izdelana v
elektrolitu brez Na2SiF6. Premer por v MAO prevleki se je zmanj{al z dodatkom Na2SiF6, medtem ko sta gostota prevleke in
hrapavost povr{ine narasli. Prevleka, nastala v elektrolitski raztopini, z ali brez Na2SiF6, ima ve~jo trdoto povr{ine kot AM60
zlitina. Rezultati obna{anja pri koroziji, vklju~no z Na2SiF6, ka`ejo na bolj{o odpornost kot pri prevleki, nastali v raztopini brez
Na2SiF6.
Klju~ne besede: magnezijeva zlitina, oksidacija v mikro obloku (MAO), Na2SiF6, korozija
1 INTRODUCTION
Magnesium (Mg) alloys have been used in many in-
dustrial applications due to their high specific strength,
low density and excellent mechanical properties. In re-
cent years, Mg alloys are widely used in automotive pro-
duction, with their low density, good castability and stiff-
ness.1–6 However, the poor corrosion resistance of Mg
alloys is restricting their applications. That is why it is
essential for magnesium alloy products to be protected
with a surface treatment.1,4,7 There are many techniques
to improve the corrosion resistance of Mg alloys, such as
electroless plating, conversion films, laser surface melt-
ing and organic coatings. Micro arc oxidation (MAO) is
another efficient method to improve the properties of Mg
alloys by producing ceramic films on their surface.1,4,8
The MAO coatings have a strong adhesion to the Mg
substrate, controllable thickness and other excellent
properties, such as corrosion resistance, thermal shock
resistance. However, the properties of MAO coatings are
affected by the processing parameters, such as the com-
position of the electrolyte, voltage, current density, time,
etc.1,9 In this work, micro arc oxidation films have been
coated on a Mg alloy with and without the Na2SiF6 in an
electrolytic solution and the structure and corrosion re-
sistance of the oxide coatings were investigated. The
properties of the coatings were characterized by scan-
ning electron microscopy (SEM), X-ray diffraction
(XRD). The results were compared and correlated to un-
derstand the influence of the Na2SiF6 in the electrolytic
solution on the coating-formation process, properties and
corrosion behavior.
2 EXPERIMENTAL PART
2.1 Material and coating process
AM60 magnesium alloy was used as the substrate
material in this study. The chemical composition of
AM60 is given in Table 1.
Table 1: Chemical composition of AZ91D magnesium alloy (in mass
fractions, w/%)
Tabela 1: Kemijska sestava magnezijeve zlitine AZ91D (v masnih
dele`ih, w/%)
Al Mn Si Fe Mg
5.93 0.18 0.02 (max.) 0.013 Balance
Materiali in tehnologije / Materials and technology 50 (2016) 5, 719–722 719
UDK 620.193:621.793:67.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)719(2016)
The samples for all the tests were cut into cylinders
with dimensions of 30 mm × 10 mm × 10 mm and
mechanically polished with 600- and 1200-grit emery
papers, rinsed with distilled water and dried in warm air.
The MAO process of the sample was coated in alkali
silicate electrolyte solution, which consisted of Na2SiO3
in distilled water with NaOH. After that 1%, 2%, and 4%
Na2SiF6 was added to the electrolytic solution. The
effects of the Na2SiF6 in the electrolytic solution on the
MAO process and the structure and mechanical pro-
perties of the oxide coatings were investigated. The
electrolyte composition is given in Table 2. The surface
roughness (Ra) of the MAO coatings was detected using
a Mahr, Perthometer M1 surface roughmeter. The
thicknesses of the coatings were measured using an SEM
and the values of the conductivity for the electrolytes
prepared with the base electrolyte and different con-
centrations of Na2SiF6 were measured and are shown in
Table 2.
The oxide coatings were produced at a constant an-
odic voltage of 370 V for 30 min. The temperature of the
electrolyte was kept at approximately 30 °C using a stir-
ring and cooling system and the current density was var-
ied in the range of 0.6–2 A cm–2. The samples were
rinsed in water and dried in hot air after the MAO pro-
cess was finished.
2.2 Microstructure
The surface morphologies of the AM60 samples
coated by MAO were characterized with scanning elec-
tron microscopy (SEM). The phase components of the
coated samples were analyzed with X-ray diffraction
(XRD) using Cu-K radiation.
2.3 Hardness test
The hardnesses of the AM60 and coated samples
were measured using an FUTURE TECH-CORP.FM-
700 microhardness tester at a load of 100 g for loading
time of 10 s. The average of three measurements was re-
ported.
2.4 Corrosion test
The immersion corrosion test was carried out in 10 %
of mass fractions of NaOH solution for 10 d in an open
system, the corrosion products were cleaned in distilled
water with an ultrasonic cleaner, all the samples were
weighed with a JA5003N electronic balance (accuracy: 1
mg) before and after the immersion test, and the corro-
sion rate was calculated from the weight-loss data. The
PH of the solution was around 12±0.5.
3 RESULTS AND DISCUSSION
The SEM microstructures of the AM60 alloy after the
MAO treatment for different electrolyte compositions are
shown in Figure 1. It is clear that an increase in Na2SiF6
in the electrolytic solution changed the surface
morphologies of the MAO coatings. The MAO coating
processed for AM60-% 0, as shown in Figures 1a and
1b, exhibits a relatively uniform surface appearance with
large pores. Figures 1c to 1f show the morphologies of
the MAO coatings when adding 1 % and 4 % Na2SiF6,
respectively, and the coatings are much rougher when
compared with Figure 1b. Na2SiF6 can change the solu-
tion’s properties, such as the solution conductivity, which
A. AYDAY: INFLUENCE OF Na2SiF6 ON THE SURFACE MORPHOLOGY AND CORROSION RESISTANCE ...
720 Materiali in tehnologije / Materials and technology 50 (2016) 5, 719–722
Table 2: Concentration of the electrolyte solution
Tabela 2: Koncentracija elektrolitske raztopine
Sample code Na2SiO3(g/L)
NaOH
(g/L)
Na2SiF6
(g/L)
Conductivity
(mS/cm)
Roughness
(μm)
Average thickness
(μm)
AM60-%0 15 5 – 14.6 2.986 31.2±5
AM60-%1 15 5 1 15.1 3.532 32.6±5
AM60-%2 15 5 2 15.3 3.565 46.63±5
AM60-%4 15 5 4 17.4 3.920 47.63±5
Figure 1: SEM images after the MAO treatment for: a), b) AM60–0 %,
c), d) AM60–1 %, e), f) AM60–4 %
Slika 1: SEM-posnetki po MAO-obdelavi: a), b) AM60–0 %,
c), d) AM60–1 %, e), f) AM60– 4 %
plays an important role in determining the morphology
and thickness. The sizes of certain pores decrease obvi-
ously with an increase of the Na2SiF6 solution, which is
considered to be related to the increasing electrolyte con-
ductivity.10 According to Table 2, the electrolyte conduc-
tivity increased an increase in the concentration of
Na2SiF6.
Table 2 reveals the roughness and average thickness
of the MAO coatings on the AM60 alloy. It can be seen
that the thickness and roughness increase with the con-
centration of the Na2SiF6, especially after 2%. The coat-
ing properties such as thickness and porosity are influ-
enced by the final voltage, which is closely related to the
solution conductivity. The electrical conductivity of the
electrolytes increases with an increase of the Na2SiF6
concentration. The higher Na2SiF6 concentration corre-
sponds to a higher current and thus a more intensive mi-
cro-arc discharge will occur on the surface.11,12
Before the coating, the average micro-hardness value
is about 60±5 HV0.1 for the AM60 alloy. After the MAO
coating, the surface hardness increases with increasing
Na2SiF6 concentration, and nearly all the coated samples
have a hardness of approximately 479±5 HV0.1
(AM60-%1). The surface hardness increases eight times
when compared with the uncoated sample
The phases identified through the analysis of the
XRD patterns for the uncoated AM60 alloy, AM60-%0
and AM60-%4 samples are presented in Figure 2. It is
clear that the bulk material is formed of Mg and
Al0.56Mg0.44 phases. However, the MAO coatings formed
of Mg, Mg2SiO4 (Forsterite), SiO2 (Silicon Oxide) and
MgO (Periclase) phases. In addition, it can be seen from
Figure 2b that the Mg2SiO4 (Fosterite) is the minor com-
mon phase that is present in all the coatings. This phase
is formed due to the composition of Si in the Na2SiO3
(present as a constituent of the electrolyte) in the coating
in the form of Mg2SiO4. Different phases were not seen
on AM60-%4 sample surface when adding Na2SiF6.
Figure 3 shows the corrosion rate variation with the
immersion time of the MAO coatings in 10 % mass frac-
tions of NaOH solution. It is clear that three characteris-
tics behaviors occur in the corrosion test, as can be seen
from the curves. First of all, the corrosion rate values of
the samples were negative at the initial periods. The
mass gain phenomenon can result from the re-oxidation
and attachment of the corrosion products. Then, the mass
loss happened after immersion for about 5 d. After a
long immersion period the coating layer and corrosion
products began to exfoliate from the samples’ surfaces.
Thirdly, the corrosion rates of the MAO-coated samples
were lower than the as-cast sample (AM60) for the
whole immersion test. This indicates that the samples
having the MAO coating with Na2SiF6 a have higher cor-
rosion resistance.
A. AYDAY: INFLUENCE OF Na2SiF6 ON THE SURFACE MORPHOLOGY AND CORROSION RESISTANCE ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 719–722 721
Figure 3: Variation of corrosion rate with immersion time in 10 % of
mass fractions of NaOH
Slika 3: Spreminjanje hitrosti korozije s ~asom potapljanja v 10 %
masnem dele`u raztopine NAOH
Figure 2: XRD patterns for: a) uncoated AM60, b) AM60-%0 and
AM60-%4
Slika 2: Rentgenogram za: a) AM60 brez prevleke, b) AM60 0 % in
AM60 4 %
4 CONCLUSIONS
Oxide coatings were produced on the AM60 alloy by
micro arc oxidation in different solutions with and with-
out Na2SiF6. The coatings produced with Na2SiF6 were
thicker than the ones produced without Na2SiF6 for the
same parameters. The pores on the surface decrease with
an increasing Na2SiF6 concentration and the surface be-
comes rougher. The hardness improves nearly eight
times when compared with uncoated sample. The corro-
sion resistance of the samples coated in the Na2SiF6 elec-
trolyte solution can be attributed to the more uniform and
compact structure of this coating, which acts as a barrier
to the transfer of corrosive ion from the aggressive solu-
tion into the coating. The AM60-%4 sample shows the
best corrosion resistance in 10 % mass fractions of
NaOH solution.
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C.-E. PELIN et al.: MECHANICAL PROPERTIES OF POLYAMIDE/CARBON-FIBER-FABRIC COMPOSITES
723–728
MECHANICAL PROPERTIES OF
POLYAMIDE/CARBON-FIBER-FABRIC COMPOSITES
MEHANSKE LASTNOSTI KOMPOZITNE TKANINE IZ
POLIAMID/OGLJIKOVIH VLAKEN
Cristina-Elisabeta Pelin1,2, George Pelin1,2, Adriana ªtefan1, Ecaterina Andronescu2,
Ion Dincã1, Anton Ficai2, Roxana Truºcã3
1National Institute for Aerospace Research "Elie Carafoli" Bucharest- Materials Unit, 220 Iuliu Maniu Blvd, 061126 Bucharest, Romania
2University Politehnica of Bucharest, Faculty of Applied Chemistry and Materials Science, 1-7 Gh. Polizu St., 011061 Bucharest, Romania
3S.C. METAV Research & Development S.A., 31 C.A. Rosetti St., 020011 Bucharest, Romania
bancristina@gmail.com, ban.cristina@incas.ro
Prejem rokopisa – received: 2015-07-01; sprejem za objavo – accepted for publication: 2015-09-10
doi:10.17222/mit.2015.171
This paper presents the production of carbon-fiber-fabric-reinforced laminated composites based on a polyamide 6 matrix using
a multiple-stages technique that involves polymer dissolution in formic acid followed by fabric impregnation and high-tempera-
ture pressing. The polyamide/solvent ratio’s influence on the interface and mechanical properties is discussed, analyzing three
PA6 weight contents of (10, 20, and 30) % in a formic acid solvent. The mechanical behavior of the obtained laminated compo-
sites is evaluated using tensile and 3-point bending tests and the fracture cross-section is analyzed using microscopy
investigation techniques in order to evaluate the fiber-matrix interface and the composite fracture mechanism. The results show
that the best mechanical performance is obtained when using a solution of 20 % mass fraction of polyamide in formic acid, as
this leads to the formation of a uniform polymer layer that is able to completely embed the fibers that constitute the fabric and
create a strong mechanical interface within the composite.
Keywords: polyamide 6, carbon fiber, mechanical properties, polymer/solvent ratio, mechanical interface
^lanek predstavlja izdelavo laminatnega kompozita na osnovi poliamida 6, oja~anega s tkanino iz ogljikovih vlaken, z uporabo
ve~stopenjske tehnike, ki vklju~uje raztapljanje poliamida v mravljin~ni kislini ter impregnacijo tkanine in stiskanje pri visoki
temperaturi. Razlo`en je vpliv razmerja poliamid/topilo na stik in mehanske lastnosti, z analizo treh masnih vsebnosti PA6 (10,
20, 30) % v mravljin~ni kislini. Mehansko obna{anje dobljenega laminiranega kompozita je ocenjeno z nateznim preizkusom in
s 3-to~kovnim upogibnim preizkusom, presek preloma pa je analiziran z mikroskopsko tehniko, da bi ocenili stik z vlaknato
osnovo in mehanizem preloma kompozita. Rezultati ka`ejo, da je najbolj{a mehanska zmogljivost dose`ena pri uporabi
raztopine z 20 % masnim dele`em poliamida v mravljin~ni kislini, ker to povzro~i nastanek enakomernega polimernega sloja, ki
lahko popolnoma obda vlakna tkanine in ustvari mo~an mehanski stik v kompozitu.
Klju~ne besede: poliamid 6, ogljikovo vlakno, mehanske lastnosti, razmerje polimer/topilo, mehanski stik
1 INTRODUCTION
In recent decades the use of composite structures in
both aeronautic and automotive applications has in-
creased tremendously. Nowadays, composites represent
approximately 50 % of the structure of the Airbus A350
XWB and the Boeing 787 Dreamliner, resulting in 20–25
% reduction in fuel consumption.1 Most composite struc-
tures are based on thermoset matrix fiber composites, but
nowadays there is an increasing interest in replacing the
thermoset with a thermoplastic matrix. This trend is due
to problems arising from thermoset composites, such as
the high costs of raw materials, high energy consump-
tion, extended processing times, non-visible damage,
complex repair procedures, recycling difficulties and sig-
nificant amounts of scrap.2–3 A thermoplastic matrix of-
fers facile recycling possibilities, lower costs and more
flexible processing routes3–6, making them a promising
solution for the shortcomings of thermoset composites.
Commercial carbon-fiber-fabric-based composites for
the transport industry obtained using melt and solvent
impregnation with thermoplastic matrixes use PEEK,
PPS or PEI7–8, which are high-performance polymers that
require over 300 °C for the processing temperature, lead-
ing to high costs. This study focuses on an engineering
polymer, with a high potential, i.e., polyamide 6. Most
literature studies present polyamide 6 composites rein-
forced with short fibers (carbon, glass or aramid) pro-
cessed by melt extrusion9–12, a few studies present fab-
ric-reinforced polyamide 6 composites13–14, because of
the technological processing difficulties arising from fab-
ric-impregnation issues generated by the molten polymer
and fiber wet out.15 Moreover, the data concerning the
optimum solution viscosity range and the optimum poly-
mer/solvent ratio when using the solvent-impregnation
method is very briefly discussed, although its importance
is underlined16–18, as the rheological properties of the ma-
trix are important for establishing the composite’s pro-
cessing parameters. This paper presents the production
of carbon-fiber-fabric-reinforced polyamide 6 matrix
laminated composites using polymer dissolution fol-
lowed by fabric solvent impregnation, solvent removal
Materiali in tehnologije / Materials and technology 50 (2016) 5, 723–728 723
UDK 67.017:621.315.614:677.494.675 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)723(2016)
and thermal pressing. The processing and final properties
of fabric-reinforced composites depend on several fac-
tors, one of them being the viscosity of the matrix used
to impregnate the fabric. In this context, the novelty of
the study is represented by an optimization study con-
cerning the PA6/formic acid (polymer/solvent) ratio’s di-
rect influence on the mechanical properties of the final
composites, analyzing the solution’s viscosity and its di-
rect effect on the fiber/matrix interface and, conse-
quently, on the failure mode of the composites. The me-
chanical behavior of the obtained laminated composites
was evaluated using tensile and 3-point bending tests, the
fracture cross-section being analyzed using microscopy
techniques to evaluate the fiber-matrix interface and the
composite fracture mechanism and to establish the opti-
mum polymer content for the impregnation solution. The
results indicate that at an optimum polymer/solvent ratio
of 20 % of mass fractions, the obtained materials possess
the highest tensile and flexural properties.
2 EXPERIMENTAL SECTION
2.1 Materials
The matrix was polyamide 6 (PA6) pellets supplied
by SC ICEFS Sãvineºti, while the solvent was formic
acid 85 % analytical grade, purchased from Chemical
Company, Romania. The reinforcing agent was a car-
bon-fiber-fabric twill weave (FC) produced by Chemie
Craft, France, 3 K warp, with 193 g/m2 fabric areal
weight and 1.7 g/cm3 fiber density.
2.2 Obtaining process
The procedure resembles a process involving fabric
impregnation with thermoset resins, but it is adjusted to
allow the use of polyamide pellets as the raw material.
The dried polymer pellets were dissolved in 85 % formic
acid in three different concentrations, i.e., 1(0, 20 and
30) % (polymer weight/solvent volume), under mecha-
nical stirring for 4 h. Each ply of the carbon-fiber fabric
was impregnated with the solution and stacked up in
groups of five layers. The solvent was removed at room
temperature for 48 h and additional traces were removed
at 80–100 °C for 8 h. Each laminated composite was
pressed using a CARVER hot platens press, following an
established temperature program, with a linear tempe-
rature increase from 25 °C to 230 °C and 5 min dwell
periods at (230, 235, 240 and 245) °C. Because the
fabric layers were semi-impregnated with polymer after
the solvent’s removal, the high-temperature pressing did
not generate impregnation difficulties that appear during
standard polymer-melt impregnation because of the fiber
wet-out difficulties.19 The cooling took place under
pressure down to room temperature. There were obtained
laminated plates differing in the PA6 content solution
used, referred to as PA6(10%)/5FC, PA6(20%)/5FC and
PA6(30%)/5FC, with an average fiber volumetric ratio of
66 %, that were processed into tensile and flexural shape
specimens.
2.3 Testing and characterization
The viscosity of the different concentration solutions
used for the impregnation was measured using an
Ubbelohde capillary-tube viscometer (Cannon CT-1000).
The PA6 matrix was subjected to FTIR spectroscopy
analysis (ThermoiN10 MX Mid Infrared FT-IR Micro-
scope) operated in ATR mode and scanning electron mi-
croscopy (QUANTA INSPECT F). The carbon-fiber
laminates were mechanically tested (INSTRON 5982
machine) in tensile conditions according to SR EN ISO
527-220 at a 5 mm/min tensile rate on dumbbell speci-
mens and flexural tests, according to SR EN ISO 1412521
at 2 mm/min, conventional deflection and span length on
rectangular specimens. The fracture cross-section was
analyzed using SEM and the fracture mode was evalu-
ated using optical microscopy (Meiji 8500).
3 RESULTS AND DISCUSSION
3.1 Viscosity measurement
The solutions containing different contents of dis-
solved polymer used to impregnate the fabric has to have
optimum viscosity, as it distributes the polymer through
the fibers.22 The effect of the PA6 content dissolved in
formic acid on the kinematic viscosity of the solution
was studied at room temperature. Figure 1 presents the
kinematic viscosity values of the solutions with three dif-
ferent PA6 weight contents, after 4 h of mechanical stir-
ring. The viscosity increases dramatically with polymer
content, from 34.1 mm2/s for 10 % PA6 in formic acid to
173 mm2/s for 20 % and 898 mm2/s for 30 %. As the
polymer content is increased, the viscosity increases ex-
ponentially by approximately five times compared to the
C.-E. PELIN et al.: MECHANICAL PROPERTIES OF POLYAMIDE/CARBON-FIBER-FABRIC COMPOSITES
724 Materiali in tehnologije / Materials and technology 50 (2016) 5, 723–728
Figure 1: Kinematic viscosity as a function of polyamide 6 weight
content in an 85 % formic acid solution
Slika 1: Odvisnost kinemati~ne viskoznosti od vsebnosti poliamida 6
v raztopini 85 % mravljin~ne kisline
previous value. The dramatic increase for the 30 % con-
tent is due to the fact that this value is very close to the
homogenous phase-formation limit of a polyamide/for-
mic acid system.23–27 The rheological properties affect
the impregnation degree and the fiber/matrix interface. If
the impregnating solution viscosity is too low, it will
pass through the fabric, resulting in polymer coverings
that are too thin, while a too high viscosity will not en-
sure uniform and large fiber-matrix contact surfaces, as
any penetration through the fibers will be difficult.28
3.2 FTIR spectroscopy
FTIR spectroscopy was performed on polyamide
films obtained after solvent removal from the three dif-
ferent solution concentrations to evaluate the chemical
structure and the eventual solvent traces. Figure 2 pres-
ents the spectra of the three PA6 samples compared to a
pure PA6 pellet, all the spectra showing the characteristic
peaks of polyamide 6.29,30 There are no significant
changes in the spectra of the samples compared with the
pure pellet, as no supplementary peak appears, and there
are no traces of unremoved solvent, confirming that the
polyamide’s chemical structure was not altered by the
solvent’s presence and it was fully dissolved. The small
modification of the bands from (689.5, 1201.4 and
1464.7) cm–1 can be assigned to the restructuring of the
polymer as a result of solubilization followed by crystal-
lization.31 The 1200 cm–1 and 1465 cm–1 peaks corre-
spond to the amide V and CH2 bending vibrations, re-
spectively, in polyamide  or  form, which is
commonly obtained when the processing of PA6 in-
volves slow cooling32, as is the case here.
C.-E. PELIN et al.: MECHANICAL PROPERTIES OF POLYAMIDE/CARBON-FIBER-FABRIC COMPOSITES
Materiali in tehnologije / Materials and technology 50 (2016) 5, 723–728 725
Figure 4: SEM images of the fracture cross-section of laminated composites: a), b) PA6(10%)/5FC, c) PA6(20%)/5FC
Slika 4: SEM-posnetki preseka preloma laminiranega kompozita: a), b) PA6 (10 %)/5FC, c) PA6 (20 %)/5FC
Figure 2: FTIR spectra of the pure PA6 pellet and dried samples
based on different PA6 contents dissolved in 85 % formic acid: PA6
(10 %), PA6 (20 %), PA6 (30 %)
Slika 2: FTIR-spektri ~istega peleta PA6 in posu{enih vzorcev z
razli~no vsebnostjo PA6, raztopljene v 85 % mravljin~ni kislini: PA6
(10 %), PA6 (20 %), PA6 (30 %)
Figure 3: SEM images of matrix samples used to form the laminated composites: a) PA6(10 %), b) PA6(20 %), c) PA6(30 %)
Slika 3: SEM-posnetki osnove vzorcev, uporabljenih za laminiran kompozit: a) PA6 (10 %), b) PA6 (20 %), c) PA6 (30 %)
3.3 SEM analysis
SEM analyses were performed on the matrix samples
in the form of films to evaluate the morphology and ho-
mogeneity. Figure 3 presents the images of dried sam-
ples obtained after the dissolution of (10, 20 and 30) %
PA6 into formic acid. As the PA6 content increased up to
30 %, there are visible areas of undissolved polymer; this
higher concentration value being close to the weight con-
tent limit of PA6 in formic acid23–26, it probably needs
different process parameters for a complete dissolution
(e.g., a longer homogenization time).
SEM investigations were performed on the fracture
cross-section of the tensile tested laminates. Figure 4 il-
lustrates the samples with a matrix obtained by dissolv-
ing 10 % PA6 (Figure 4a and 4b) and 20 % PA6 (Figure
4c). In PA6(10%)/5FC (Figure 4a) there are several ar-
eas where a thin polymer layer uniformly covers the fi-
bers of the fabric, but there are also a few areas where
the polymer layer is partially detached from the fiber sur-
face and there are uncovered fibers (Figure 4b). The dif-
ference is significant for PA6(20%)/5FC, in which case
the polymer is distributed in a solid layer that covers the
entire fiber surface, but its thickness is not as uniform as
in PA6(10%)/5FC. Each carbon fabric ply is covered
with its own polymer layer. The polymer ductile fracture
is distinguished from the fiber’s fragile fracture.
3.4 Mechanical testing
Table 1 illustrates the average values of the strength
and elasticity modulus, obtained after tensile and flexural
testing of the obtained composites based on five carbon
fabric plies. The highest mechanical performance in both
the tensile and flexural testing is presented by the sample
based on PA6(20%). PA6(20%)/5FC has a 60 % higher
tensile strength compared to PA6(10%)/5FC, while the
flexural strength is approximately 70% higher. These
samples also have a superior rigidity, showing a modulus
of elasticity that is higher by 35–40 %. The
PA6(30%)/5FC samples showed lower tensile and flex-
ural strengths and tensile moduli compared with
PA6(20%)/5FC, but higher than the PA6(10%)/5FC,
while the flexural modulus had the lowest average value
for the entire series. These inconsistent results in
PA6(30%)/5FC are most likely due to the high viscosity
of the impregnating solution that was probably not able
to uniformly distribute on the surface of the fiber fabric,
supplemented by possible undissolved polymer sites,
which although they were melted during thermal press-
ing, could also lead to non-uniform polymer layer sites at
the microscopic level. These issues generate
inhomogeneity and stress-concentration sites that de-
crease the rigidity.33–34
Table 1: Mechanical properties of composites based on PA6 (dis-
solved in different contents relative to the solvent) and five plies of
carbon fiber fabric
Tabela 1: Mehanske lastnosti kompozita na osnovi PA6 (raztopljene
razli~ne koli~ine v topilu) in petih plasti tkanine iz ogljikovih vlaken
Sample
Tensile
strength
(MPa)
Young’s
modulus
(GPa)
Flexural
strength
(MPa)
Young’s
flexural
modulus
(GPa)
PA6(10%)/5FC 339.2 45.5 436.7 38.3
PA6(20%)/5FC 540.5 63.6 732.6 51.4
PA6(30%)/5FC 505 56.3 571.6 38.1
Overall, the mechanical results show that all the ob-
tained carbon fabric/PA6 laminated samples are superior
to the ones exhibited by long-carbon-fiber-reinforced
PA635–37 that show a tensile strength between 240 and
300 MPa, a flexural strength between 330 and 500 MPa,
and Young’s modulus in the range 25–40 GPa for tensile
and 20–30 GPa for flexural. Also, PA6(20%)/5FC me-
chanical properties are comparable with those presented
by carbon-fiber-fabric-epoxy composites38 with extended
applications in aeronautics.
3.5 Fractography
Optical microscopy images were recorded on the
fracture region of the laminated samples to establish the
fracture mechanisms that led to the failure and to evalu-
ate their behavior when tested under tensile and flexural
loads. Figure 5 presents the fracture region of represen-
tative specimens from each sample tested in tensile. The
identified fracture mechanisms are marked, all of them
being classified as classical mechanisms presented by 39–40:
(1) crack propagation, (2) layer de-bonding, (3) fiber
breakage, (4) fiber pull-out, (5) ply breakage. The
PA6(10%)/5FC and PA6(20%)/5FC fractures resemble
as three layers are broken by fiber breakage in the same
area, the fracture causing layer de-bonding to the next
C.-E. PELIN et al.: MECHANICAL PROPERTIES OF POLYAMIDE/CARBON-FIBER-FABRIC COMPOSITES
726 Materiali in tehnologije / Materials and technology 50 (2016) 5, 723–728
Figure 5: Fracture regions of tensile-tested specimens: a) PA6
(10 %)/5FC, b) PA6 (20 %)/5FC, c) PA6(30%)/5FC
Slika 5: Podro~ja preloma nateznih preizku{ancev: a) PA6
(10 %)/5FC, b) PA6 (20 %)/5FC, c) PA6 (30 %)/5FC
layers, with more pronounced de-bonding in
PA6(10%)/5FC. PA6(30%)/5FC presented the most
destructive failure, involving several mechanisms
including fiber pull out and interlayer crack propagation,
leading to delaminated areas. This can be explained by
the high viscosity of the solution that led to non-uniform
matrix layers, creating stress-concentration sites. In
flexural testing (Figure 6), PA6(10%)/5FC and
PA6(20%)/5FC did not present any layer rupture until
conventional deflection, but the PA6(30%)/5FC
presented the rupture of one external layer that
de-bonded on a longer length. It is important to mention
that although PA6(10%)/5FC presented in general the
lowest mechanical performance, its fracture mode was
not as destructive as for PA6(30%)/5FC. The optical
images complement both the mechanical test results and
the SEM studies.
4 CONCLUSIONS
The study presents the production of carbon-fab-
ric-reinforced laminated composites based on the engi-
neering polymer polyamide 6 as the matrix using a mul-
tiple-stages technique that involves fabric solvent
impregnation with a formic acid solution that contains
different contents of dissolved polymer, solvent removal
and high-temperature pressing. The aim of the study was
to evaluate the polymer/solvent (PA6/formic acid) ratio’s
influence on the mechanical interface within the compos-
ite and on the mechanical properties. Three polymer con-
tents in formic acid were used (10, 20 and 30) %
weight/volume, with the rheology tests showing that the
solution viscosity increases exponentially. A lower con-
tent solution led to a slightly uniform polymer layer that
covered the fibers of the fabric, ensuring large contact ar-
eas, but because the thin layer was too weak, the lami-
nates’ mechanical performance was lower. The highest
content solution (30 %) had an extremely high viscosity
and most likely did not distribute uniformly, probably
generating stress-concentration sites that resulted in a
very destructive failure during the mechanical testing.
The study concludes that the optimum content is
20 % polymer dissolved into the solvent, as it leads to
medium-viscosity solution that supports polymer pene-
tration through the fibers and forms polymer layers with
suitable thickness and uniformity on the fiber surface.
This ensures high contact areas, a strong fiber-matrix in-
terface and, consequently, an optimum fiber-matrix load
transfer, leading to high mechanical performances in ten-
sile and bending and failure modes that do not exhibit a
high degree of delamination. The results show that at an
optimum polymer/solvent ratio, these composites can
represent potential solutions as materials for high me-
chanical performance in aeronautics and automotive ap-
plications.
Acknowledgments
This work was funded by the Sectoral Operational
Programme Human Resources Development 2007-2013
of the Ministry of European Funds through the Financial
Agreement POSDRU/159/1.5/S/132397 and by Roma-
nian Ministry of Education through the PN-II-PT-PCCA-
168/2012 project: “Hybrid composite materials with
thermoplastic matrices doped with fibres and disperse
nano fillings for materials with special purposes”.
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tions, RTP 287 A Nylon 6 (PA) Carbon Fiber, Minnesota, USA 2014
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tive study of mechanical properties of woven of carbon fiber twill
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ference on Composite Materials – BCCM2, São José dos Cam-
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posite by microscopic technique, Bachelor’s Degree Thesis, Depart-
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Current status and future issues, Proc. of the 13th European confer-
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728 Materiali in tehnologije / Materials and technology 50 (2016) 5, 723–728
K. DVOØÁK et al.: EVALUATION OF THE GRINDABILITY OF RECYCLED GLASS ...
729–734
EVALUATION OF THE GRINDABILITY OF RECYCLED GLASS IN
THE PRODUCTION OF BLENDED CEMENTS
OCENA SPOSOBNOSTI DROBLJENJA RECIKLIRANEGA STEKLA
PRI PROIZVODNJI ME[ANIH CEMENTOV
Karel Dvoøák1, Du{an Dolák1, Dalibor V{ianský2, Petr Dobrovolný1
1Brno University of Technology, Faculty of Civil Engineering, Veveøí 331/95, 602 00 Brno, Czech Republic
2Masaryk University, Faculty of Science, Kotláøská 267/2, 611 37 Brno, Czech Republic
dolak.d@fce.vutbr.cz
Prejem rokopisa – received: 2015-07-01; sprejem za objavo – accepted for publication: 2015-09-14
doi:10.17222/mit.2015.184
The replacement of primary raw materials in cement production is a current topic. Potentially usable raw materials include recy-
cled glass. The disadvantage of glass is its tendency to aggregate. The conventional method for the production of blended ce-
ment is separate grinding and then homogenizing the components. However, in this case aggregated fine glass in the cement
composites acts only physically and mechanically as a filler rather than as an active pozzolan. An interesting option to prevent
the formation of aggregates formed from glass is co-grinding. This procedure is not very common in practice. Various ingredi-
ents of blended cements have widely different grindabilities, and it is therefore better to grind them separately. The aim is to
compare the co-grinding and separate grinding of a combination of Portland clinker and recycled glass. The grindability was
tested on clinker, glass, and blended cements prepared by co-grinding and by separate grinding. The results of the experiment
show that by co-grinding the components of blended cement with the addition of 20 % of recycled glass as a pozzolan, a syn-
ergy effect caused by the various mechanical properties of the components occurs. The aggregation of grains is less significant
than during separate grinding and it leads to a better grinding effect. This knowledge can by utilized in the design and process-
ing of new blended cements. Also, co-milling the glass-cement system can eliminate the stage of homogenization, and, therefore
save energy.
Keywords: grindability, Portland clinker, recycled glass
Tema ~lanka je nadome{~anje primarnih surovin pri proizvodnji cementa. Potencialno uporabno surovino predstavlja reciklirano
steklo. Pomanjkljivost stekla je, da ima nagnjenost k sprijemanju. Obi~ajna metoda proizvodnje me{anega cementa je lo~eno
drobljenje in homogenizacija sestavin. V tem primeru drobnozrnato steklo v cementnih me{anicah deluje samo fizikalno in
mehansko kot polnilo in ne kot aktiven pozolan. Dodatno mletje je pomembno za prepre~evanje nastanka skupkov stekla.
Vendar pa ta postopek ni tako pogost v praksi. Razli~ne sestavine cementne me{anice se razli~no drobijo in je zato bolje, da se
jih drobi lo~eno. Namen {tudije je primerjati dodatno mletje in lo~eno mletje kombinacije Portland klinkerja in recikliranega
stekla. Mletje je bilo preizku{eno na klinkerju, steklu in me{anici cementov, pripravljenih z dodatnim mletjem in z lo~enim
mletjem. Rezultati preizkusov so pokazali, da se pri isto~asnem mletju me{anic cementov z dodatkom 20 % recikliranega stekla
kot pozolana, pojavi sinergijski pojav zaradi razli~nih mehanskih lastnosti sestavin. Zdru`evanje zrn je manj izrazito kot pa pri
lo~enem mletju in povzro~i bolj{i u~inek mletja. To dejstvo je mogo~e uporabiti pri na~rtovanju in izdelavi novih me{anic
cementov. Torej se lahko z isto~asnim mletjem sistema steklo-cement, odpravi fazo homogenizacije in s tem prihrani energijo.
Klju~ne besede: sposobnost mletja, portlandski klinker, reciklirano steklo
1 INTRODUCTION
Secondary raw materials represent an ever more fre-
quent replacement for primary raw materials in the pro-
duction of building materials. The area of cement pro-
duction is no exception. In current practice, blended
cements are applied increasingly more often. In these ce-
ments, the Portland clinker is replaced by hydraulically
active compounds or agents with pozzolanic proper-
ties.1–3 Glass is chemically and mineralogically very
close to traditional pozzolans. Therefore, various types
of recycled glass may be potentially interesting raw ma-
terials for the production of blended cements. Various
authors have described the behavior of finely ground
glass in cement composites.4,5 However, due to a consid-
erable ability to agglomerate, the recycled glass used as
an additive for the cement composite is not reactive
enough and acts only physically-mechanically as a filler.4
The common production process for blended cements is
separate grinding of the individual components and their
subsequent homogenization. This procedure is common
in the production of blast-furnace slag cements. In this
case the procedure is advantageous because Portland ce-
ment clinker and blast-furnace slag have very different
grindabilities and it is therefore preferable to grind them
separately, and subsequently to homogenize.6 As noted
above, fine glass powder exhibits a significant ability for
aggregation, which greatly complicates the homogeniza-
tion with Portland cement. Therefore, the traditional ap-
proach of separate grinding and subsequent homogeniza-
tion seems to be less suitable in the case of a
glass-cement system. An interesting option to prevent
the formation of agglomerates with pure glass is
co-grinding of the glass and clinker. The content of SiO2
in recycled glass is only in amorphous form and the
hardness is 7 on the Mohs scale. The standard alite
Materiali in tehnologije / Materials and technology 50 (2016) 5, 729–734 729
UDK 620.1:621.926:621.315.612:621.742.48 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)729(2016)
clinker contains four main minerals, and their weighted
average hardness is between 6 and 7 on the Mohs scale.7
However, the recycled glass is much more fragile, which
means the grindability of both components could be very
similar. Various authors have chosen different methods to
assess grindability. Most methods are based on an evalu-
ation of the ratio of the energy consumption and refine-
ment of the material.8–12 An interesting approach is to
evaluate the grindability by using particle size distribu-
tion curves.6 The aim was to assess whether in the case
of a cement-glass system, co-grinding of the component
is more advantageous than separate grinding with subse-
quent homogenization. The selected approach was to
monitor and compare the grindability of the individual
components as well as the mix. The method of monitor-
ing the impact of the constant grinding time on the parti-
cle size distribution curves and specific surface area was
chosen for the experiment.
2 MATERIALS AND METHODS
For monitoring the grindability, recycled glass and
Portland cement were used. The chemical composition
of the recycled glass was determined by traditional
chemical analysis. The modified Chapelle test method13
was used for the pozzolanic activity determination. The
modified Chapelle test consists of the reaction of
pozzolan and freshly annealed CaO in an aquatic envi-
ronment at 93 °C for 24 h. The reaction takes place in a
tightly sealed stainless-steel vessel and the suspension is
stirred by an electromagnetic stirrer. The result is
expressed as the amount of Ca(OH)2 bound in mg per 1 g
of pozzolan. The density of the recycled glass was deter-
mined by automatic pycnometer Micromeritics AccuPyc
II 1340. For the measurement of the specific surface
according to Blaine, an automatic PC-Blaine-Star device
was used to measure the cell capacity of 7.95 cm3. The
determination was performed three times to eliminate
errors. The morphology of the particles was determined
by scanning electron microscope (SEM). A Tescan
MIRA 3 XMU SEM with a secondary-electron detector
was used. The Portland cement was prepared in a labo-
ratory ball mill by co-grinding of the Portland clinker
from cement plant Hranice and the chemo-gypsum
Pregips in the ratio 95/5. Milling was carried out to the
same specific surface area that was measured on the
recycled glass. The chemical composition, the density,
the specific surface area, and morphology of the particles
were also determined. The blended cement was prepared
by co-milling Portland clinker, gypsum and recycled
glass in the ratio 76/4/20. Milling was carried out to the
same specific surface area that was measured on the
recycled glass. As in the previous case, chemical compo-
sition, density, specific surface area and morphology of
the particles were also determined. The milling in this
phase of the experiment was always carried out at a total
dose of 5 kg in a Brio OM 20 ball mill at a speed of
40 min–1. The grinding of the recycled glass, the Portland
cement and the blended cement for determining the
grindability was performed in a Fritsch Pulverisette 6
planetary mill at 500 min–1. A steel vessel of 500 mL and
25 steel grinding balls of 20 mm diameter and a mass of
180 g of material were always used. The grinding times
were (1, 2, 3 and 5) min. Then, the particle size
distribution was performed on each of these samples
using a Matest Air jet sieve. The sieves mesh size were
(0.010, 0.020, 0.041, 0.063, 0.090 and 0.125) mm. The
Blaine specific surface area and morphology of the
particles by SEM were determined on all the samples
ground for 5 min. A simple calculation using weighted-
average values of the surface areas separately milled
components was made to facilitate the grindability com-
parison, Equation (1):
SBC = 0.8 · SPC + 0.2 · SGR (1)
Where SBC is the theoretical specific surface area of
the blended cement, SPC is the specific surface area of the
Portland cement and SRG is the specific surface area of
the recycled glass. Subsequently, a sample of the blended
cement was prepared by homogenization of the sepa-
rately ground components in the same proportions. Ho-
mogenization of the sample was performed using a labo-
ratory homogenizer for 1 h. On the resulting samples the
specific surface area according to Blaine was deter-
mined. The results were compared with the calculation.
3 RESULTS
Chemical compositions of the clinker and gypsum
are summarized in Tables 1 and 2.
Table 1: Partial chemical composition of clinker
Tabela 1: Parcialna kemijska sestava klinkerja
Component SiO2 CaO Al2O3 Fe2O3 SO3 Others
Content (%) 20.29 65.33 5.21 5.04 0.79 3.34
Table 2: Partial chemical composition of gypsum
Tabela 2: Parcialna kemijska sestava mavca
Component CaSO4·2H2O H2O CaSO4 Others
Content (%) 84.00 11.00 2.40 2.60
The chemical composition of the selected clinker is
typical for Portland clinkers. In the case of gypsum, it is
highly pure with a relatively high humidity; therefore, it
should be dried for the cement preparation to reduce the
humidity to under 5 % according to ^SN 721206.
Chemical composition of the recycled glass is sum-
marized in Table 3 and its pozzolanic activity is indi-
cated in Table 4.
Table 3: Partial chemical composition of the recycled glass
Tabela 3: Parcialna kemijska sestava recikliranega stekla
Component SiO2 CaO Al2O3 K2O Na2O Others
Content (%) 69.25 8.09 0.83 0.41 16.44 4.98
K. DVOØÁK et al.: EVALUATION OF THE GRINDABILITY OF RECYCLED GLASS ...
730 Materiali in tehnologije / Materials and technology 50 (2016) 5, 729–734
Table 4: Pozzolanic activity of recycled glass with different specific
surface area
Tabela 4: Pozolanska aktivnost recikliranega stekla z razli~no
specifi~no povr{ino
Specific surface area (m2 kg–1) 244
Pozzolanic activity (mg Ca(OH)2/g pozzolan) 1112
The chemical and mineralogical compositions of the
recycled glass resemble a classic pozzolan. The sample
of recycled glass with a specific surface area of 244 m2
kg–1 reached a pozzolanic activity of 1112 mg Ca(OH)2/1
g in a modified Chapelle test.
An overview of the properties of the raw materials on
the grindability are included in Table 5. All the
pre-grinding was made on a ball mill to ensure roughly
the same surface area as the recycled glass.
Table 5: Overview of input materials
Tabela 5: Pregled vhodnih materialov
Material Compo-nents (%)
Pre-
ground
Density
(kg/m3)
Specific
surface
area
(m2 kg–1)
Portland
cement
Clinker 95
Yes 3081 250
Gypsum 5
Glass Glass 100 No 2458 244
Blended
cement
Clinker 76
Yes 2952 247Gypsum 4
Glass 20
All the input materials were then ground in a
planetary mill for (1, 2, 3 and 5) min with a rotational
speed of 500 min–1. Each sample was then examined
with a sieve analysis. The results are summarized in Fig-
ures 1 to 5.
From the size distribution curves of the input materi-
als it is evident that although the specific surface area is
similar, the recycled glass is much coarser. This is
caused by the different shapes of the grains in the recy-
K. DVOØÁK et al.: EVALUATION OF THE GRINDABILITY OF RECYCLED GLASS ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 729–734 731
Figure 4: Size distribution after 3 min
Slika 4: Razporeditev velikosti po 3 min
Figure 1: Size distribution of input materials
Slika 1: Razporeditev velikosti vhodnih materialov
Figure 2: Size distribution after 1 min
Slika 2: Razporeditev velikosti po1 min
Figure 3: Size distribution after 2 min
Slika 3: Razporeditev velikosti po 2 min
cled glass, Portland cement and blended cement, as seen
in Figures 6 to 8. The curves of the particle size distribu-
tion of the grinded materials indicate that the grindability
of the recycled glass and the Portland cement is similar
for selected time intervals. However, by co-grinding
these materials, the grinding effect was stronger.
The images taken by scanning electron microscopy
before and after the grinding of the components are
K. DVOØÁK et al.: EVALUATION OF THE GRINDABILITY OF RECYCLED GLASS ...
732 Materiali in tehnologije / Materials and technology 50 (2016) 5, 729–734
Figure 8: SEM image of RG: a) before and b) after grinding
Slika 8: SEM-posnetek RG: a) pred in b) po mletju
Figure 6: SEM image of BC: a) before and b) after grinding
Slika 6: SEM-posnetek BC: a) pred in b) po mletju
Figure 5: Size distribution after 5 min
Slika 5: Razporeditev velikosti po 5 min
Figure 7: SEM image of PC: a) before and b) after grinding
Slika 7: SEM-posnetek PC: a) pred in b) po mletju
shown in Figures 6 to 8. BC is an abbreviation for
blended cement, PC for Portland cement and RG for re-
cycled glass.
The Blaine specific surface area was determined on
all the samples ground for 5 min. The results of the mea-
surement and the calculated specific surface area are
summarized in Table 6.
SBC = 0.8 · SPC + 0.2 · SGR
SBC = 0.8 · 500 + 0.2 ·517 = 503.4
Table 6: Change of specific surface area before and after 5 min of
grinding
Tabela 6: Sprememba specifi~ne povr{ine, pred in po 5 min mletju
Component Cement Glass
Blended
c (to-
gether)
Blended
c (sepa-
rate)
Blended
c (calcu-
lation)
Specific sur-
face area be-
fore g (m2/kg)
250 244 247 – –
Specific sur-
face area after
g (m2/kg)
500 517 532 504 503.4
Specific surface area of separate grinded blended ce-
ment corresponds with the calculation. In the case of
co-grinded blended cement the specific surface area is
considerably higher.
4 DISCUSSION
The chemical and mineralogical compositions of the
recycled glass resemble a classic pozzolan. The sample of
recycled glass with a specific surface area of 244 m2 kg–1
reached a pozzolanic activity of 1112 mg Ca(OH)2/1 g in
a modified Chapelle test. The pozzolan activity of the
chosen recycled material can be rated as high, because
classic pozzolans such as fly ash reach values of 700 mg
to 850 mg.14 Therefore, it can be stated that this is a pro-
mising pozzolanic material. The chemical composition
of the selected clinker is typical for Portland clinkers
with a large amount of tricalcium silicate.7 In the case of
gypsum, it is a highly pure by-product gypsum from the
production of titanium dioxide.
The grain morphology of the Portland cement and
glass, which were adjusted to the same initial surface
area and were used as the input for the grindability tests,
are significantly different. Unlike Portland cement, recy-
cled glass consists of grains with a substantially sharp-
edged morphology. This grain shape can be explained by
the high fragility and amorphous structure of the glass.15
From the measured values of the balances of the raw ma-
terials on the sieves, the statement can be made that with
a low surface area and larger grain size, Portland cement
is milled the best. Glass is indeed fragile, but has a
higher tendency to aggregate the particles.15,16 This phe-
nomenon can affect the outcome of the determination of
the particle size distribution during the early stages of
grinding. As shown in Figures 2 to 5, with increasing
surface area, blended cement is ground more intensively
than recycled glass or Portland cement. Increased effi-
ciency co-milling is caused by the different mechanical
properties of the clinker and the glass.7,15 When recycled
glass is milled separately, it preserves the delicate char-
acter and this leads to its rapid disintegration. Neverthe-
less, the distinctive ability of aggregation and agglomera-
tion negatively affects the final particle size distribution,
as evidenced in Figure 8. In the case of clinker, the abil-
ity to aggregate and agglomerate is lower.17 Because of
the chemical and mineralogical compositions of the
grains they are more able to compensate for the impacts
of the grinding elements. This affects the particle size
distribution. The co-milling of cement and recycled glass
leads to a better milling effect, since the above-described
phenomena are compensated, plus the clinker grains are
functioning as an additional grinding medium. This is re-
flected not only in the resulting particle size distribution
obtained on specific surfaces, but also in a better homo-
geneity of mixed cement. The synergistic effect of
co-milling in the case of cement glass was proved by the
simple calculation model of weighted averages for the
results of the separately ground materials’ surface areas.
The result of the calculation, 503.4 m2 kg–1, correlated
well with the experimentally measured value of the spe-
cific surface area of the blended cement composed of
separately milled components, i.e., 504 m2 kg–1. By joint
grinding of the blended cement in same grinding condi-
tions, a much larger specific surface area has been
achieved, i.e., 532 m2 kg–1.
5 CONCLUSION
The pozzolanic activity of the fine recycled glass is
relatively high. It reaches higher levels of pozzolanic ac-
tivity than traditional ash, and on a significantly lower
surface area. The distribution of particles of recycled
glass and Portland cement measured by sieve analysis
with separate grinding is similar. The synergistic effect
of co-milling was demonstrated in comparison with a
blended cement prepared by the homogenization of sepa-
rately ground materials. This phenomenon is caused by
the fact that the negatives associated with separate grind-
ing of the individual materials are suppressed. Another
advantage of co-milling the glass-cement system is en-
ergy saving, by eliminating the stage of homogenization.
Based on the results obtained, recycled glass appears as a
potentially useful pozzolan for the preparation of
blended cements by co-milling with cement.
Acknowledgment
This work was financially supported by project num-
ber: 15-08755S "Study of effects of samples preparation
on inorganic binders final properties" and project No.
LO1408 "AdMaS UP – Advanced Materials, Structures
and Technologies", supported by Ministry of Education,
K. DVOØÁK et al.: EVALUATION OF THE GRINDABILITY OF RECYCLED GLASS ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 729–734 733
Youth and Sports under the National Sustainability
Programme I.
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Comminution Circuits, Chapter 12 – Determination of the bond work
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Labo. P. et Ch., 93, (1978), 61–65
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Pavnlík, R. ^erný, Application of a-SiO2 Rich Additives in Cement
Paste, Applied Mechanics and Materials, 749 (2015), 362–367,
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734 Materiali in tehnologije / Materials and technology 50 (2016) 5, 729–734
J. SZYMANSKA et al.: RHEOLOGICAL PROPERTIES OF ALUMINA CERAMIC SLURRIES FOR CERAMIC ...
735–738
RHEOLOGICAL PROPERTIES OF ALUMINA CERAMIC
SLURRIES FOR CERAMIC SHELL-MOULD FABRICATION
REOLO[KE LASTNOSTI GO[^E IZ GLINICE ZA IZDELAVO
KERAMI^NIH KALUPOV
Joanna Szymañska, Pawe³ Wiœniewski, Marcin Ma³ek, Jaros³aw Mizera
Warsaw University of Technology, Faculty of Materials Science and Engineering, Woloska Street 141, 02-507 Warsaw, Poland
joanna.szymanska.pl@gmail.com
Prejem rokopisa – received: 2015-07-01; sprejem za objavo – accepted for publication: 2015-09-15
doi:10.17222/mit.2015.188
This research is about the properties of ceramic slurries prepared from hydrous nano-alumina-based binder and a corundum
matrix used for fabricating the prime coat of ceramic shell moulds. Solid-state alumina powders with different granulations were
used. The modification of the technological properties of the prepared slurries was based on additions of a polyacrylic binder
with different amounts of polymer with respect to the alumina for different powder ratios. The slurries were prepared and tested
in a mechanical mixer. During the slurry preparation (within 96 h), the plate weight, Zahn cup 4# viscosity and dynamic
viscosity were controlled. The morphology and chemical properties of corundum powders and polymer were characterized with
SEM and powder-grain-size distribution. The obtained results of the corundum-based ceramic slurries indicate that the
application of a polymeric binder with various concentrations based on nano-alumina oxides causes different properties in
comparison to the other commonly used binders.
Keywords: ceramic slurries, investment casting, shell moulds, alumina powder
Raziskava obsega lastnosti kerami~nih go{~, pripravljenih iz nanoglini~nega veziva na vodni osnovi in korundne osnove,
uporabljenih za prvo prevleko pri izdelavi kerami~nih tankostenskih form. Uporabljen je bil prah glinice v trdnem stanju, z
razli~no zrnatostjo. Spreminjanje tehnolo{kih lastnosti pripravljene go{~e je temeljilo na dodatku poliakrilnega veziva z razli~no
vsebnostjo polimerov, glede na glinico pri razli~nih razmerjih praha. Go{~e so bile pripravljene in preizku{ene v mehanskem
me{alniku. Med pripravo go{~e (v okviru 96 h), je bila kontrolirana te`a plo{~, viskoznost 4# z Zahn potopnim viskozimetrom
in dinami~na viskoznost. Morfologija in kemijske lastnosti korundnih prahov in polimerov so bile dolo~ene s SEM in z
razporeditvijo velikosti zrn. Dobljeni rezultati kerami~nih go{~ na osnovi korunda ka`ejo, da se z uporabo polimernega veziva in
razli~ne koncentracije nanoglinice, dose`e razli~ne lastnosti v primerjavi s standardno uporabljanimi vezivi.
Klju~ne besede: kerami~na go{~a, precizijsko litje, tankostenska forma, glinica v prahu
1 INTRODUCTION
The investment-casting process is commonly applied
in the manufacturing of the materials for the aviation,
energy and military industries. The limiting components
(flight safety parts) such as aircraft turbine blades
characterized by complicated shapes are cast with the
Bridgman method.1 A commonly applied technique is
the lost-wax processing including the use of ceramic
shells. It determines the precise shape, dimensional
accuracy, appropriate structure and metallurgic purity of
designed parts. So far, ceramic shell moulds were
fabricated on the basis of colloidal silica. However, the
presence of SiO2 in the prime coat during the Ni- or
Co-superalloy casting causes a reaction with the liquid
metal at a high temperature, inducing an oxidation of the
reactive metal such as Hf. Such an adverse phenomenon
reduces the quality of the properties of cast parts,
affecting its exploitation time.2
The basic components for a ceramic slurry are
binders and fillers in the form of ceramic powders and
supportive materials. A commonly used binder is hydro-
lyzed tetraethylorthosilicate together with organic com-
pounds of silicon.3 However, pure ethyl silicate does not
have the binding capacity. Water-based binders dry more
slowly than alcohol-based ones. Consequently, there is a
time elongation enabling the control of the surface
smoothness, permeability, strength and dimensional
stability of the model.4,5
A proper selection of powder for ceramic shell
moulds and their parameters such as the kind, shape and
size of particles affect the final characteristics of the cast
elements.6 Ceramic powders present a thermal resistance,
a slight thermal expansion and a lack of polymorphic
transitions.
Deflocculants, softeners and surfactants mainly deter-
mine the rheological properties of ceramic slurries.3
It was found that a nano-Al2O3-based binder does not
react with Ni-alloy components. Moreover, such a binder
demonstrates a higher melting point and a larger surface
area than other inorganic solvents. It is also characterized
by an improved dispersion of the particles in water, thus
allowing the control of rheological properties by
preventing the sedimentation of heavy particles in a cera-
mic slurry.6,7 This is why such a binder can be applied
instead of the colloidal silica-based binder.
Materiali in tehnologije / Materials and technology 50 (2016) 5, 735–738 735
UDK 67.017:666.122.3:666.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)735(2016)
The main aim of the following research was to
examine and define the properties of a nano-aluminum-
oxide-based binder and a corundum matrix using a
polymeric binder with various concentrations.
2 MATERIALS AND EXPERIMENTAL
METHODS
The subject of this research was a powder of Al2O3
with granulation of 0–30 and 200 mesh (Treibacher)
characterized by the average size of 11.79 and 45.00 μm,
respectively. Solvent, binder and a hydrous polymer were
dispersed in colloidal Al2O3 with a particle size of 16 nm
(Imerys, Evonik). The additive material applied to
modify the rheological properties of the slurry was a
poly acrylic polymer (Imerys, Evonik).
Ceramic slurries with a solid phase content of 72.5 %
by weight and polymer amounts of (6, 10, 15) % mass
fractions with respect to the alumina for different powder
ratios of 35:65 and 65:35 (200:030 mesh) were prepared
in a mechanical mixer within 96 h with a speed of
160 min–1. During the slurry preparation, the pH (with
the use of a pH meter), plate weight and Zahn cup 4#
viscosity were checked every 24 h. These measurements
are fundamental for the investment-casting industry.
After 96 h of mixing, rheological properties such as
dynamic viscosity were also defined with a Brookfield
DV-II rheometer with the spindle rotating in a speed
range of 1–200–1 min–1. All the measurements were
taken in an air-conditioned lab at 21 °C.
To characterize the morphology of the corundum
powders and the polymer, SEM images were taken with
a Hitachi SU70 scanning electron microscope and a BSE
detector at a voltage of 5 kV. A particle-size test was
done using a Horiba LA-950 laser diffraction device
(Hitachi, Japan).
The plate test was based on immersing the plate
(7.5×7.5 cm) in the moulding mass and estimating its
weight after 120 s.
3 RESULTS AND DISCUSSION
The morphology of the powder based on the SEM
analysis of #200 and #0–30 indicated typical structures
of molten powders with angular-shaped particles.
The obtained results shown on Figure 1 prove that
the lowest plate values correspond to 6 % of mass frac-
tions of the polymer content for a powder ratio of 35:65.
The largest ones were noticed for the slurries with 6 and
15 % of mass fractions of the polymer content at a pro-
portion of 65:35. The highest plate stability was obtained
for the slurries with 15 % of mass fractions of polymer
addition (65:35) and 10 % of mass fractions of polymer
content (35:65). The values of the plate weight con-
trolled on the last days of the measurements were in a
range of 1.7–2.4 g. The measurements of the plate weight
revealed a correlation with the polymer content: a 6 % of
mass fraction of the polymer addition resulted in the
highest weight value, equal to 2.40 g; this value was
slightly lower in the case of a 15 % of mass fraction of
the polymer content and the lowest for a 10 % of mass
fraction of the polymer amount.
Zahn Cup 4# measurements showed (Figure 2) the
lowest values (13–15 s) for the slurry with the 6 % of
mass fraction of the polymer at the 35:65 ratio. The
highest viscosity was noticed for the slurry with the
15 % of mass fraction at a powder ratio of 65:32. In this
case, there was also a rapid viscosity change from 32 s
(noticed on the first day) to 21 s after 96 h. The viscosity
was stable during the whole ceramic-slurry preparation
process for 6 (at 65:35) and 10 % of of mass fractions (at
35:65).
The obtained results shown on Figure 3 indicate sta-
bility of all the measured slurries within the measure-
ment time. The lowest values were noticed for the
slurries with 10 % of mass fractions of the polymer
content at the 65:35 ratio and for 6 of % of mass
fractions of the polymer addition at the 35:65 powder
proportion. The thickness values estimated after 96 h
oscillated from 0.12 to 16 mm. Moreover, the slurries
J. SZYMAÑSKA et al.: RHEOLOGICAL PROPERTIES OF ALUMINA CERAMIC SLURRIES FOR CERAMIC ...
736 Materiali in tehnologije / Materials and technology 50 (2016) 5, 735–738
Figure 2: Relation between Zahn cup 4# viscosity and stirring time
for ceramic slurries with 72.5 % of mass fractions of solid content for
different polymer contents at 35:65 and 65:35 powder ratios (200:030)
Slika 2: Razmerje med Zahn viskoznostjo 4# in ~asom me{anja kera-
mi~ne go{~e z 72,5 % trdnega masnega dele`a, pri razli~ni vsebnosti
polimera in razmerju prahov 35:65 in 65:35, mre`a (200:030)
Figure 1: Relation between plate weight and stirring time for ceramic
slurries with solid-content mass fraction of 72.5 % for different
polymer contents at 35:65 and 65:35 powder ratios (200:030)
Slika 1: Razmerje med te`o plo{~e in ~asom me{anja go{~e z
vsebnostjo 72,5 % trdnega masnega dele`a, pri razli~nih vsebnostih
polimera in razmerju prahov 35:65 in 65:35 (200:030)
with 6 and 10 % of mass fractions of the polymer
content were characterized as similar according to the
viscosity level.
The measurements of the ceramic-slurry dynamic
viscosity are shown on Figures 4 do 6 where the rela-
tionship between the shear rate and viscosity is pre-
sented. As seen on the diagrams, additions of different
concentrations of polymers to the ceramic slurries of
Al2O3, with two powder ratios, determine their viscosity.
The obtained results indicate that the application of 10 %
of mass fractions of polymer at the 35:65 powder ratio
causes the largest increase in the dynamic viscosity
where the maximum value is 763 MPa s. The most
effective was the addition of 6 % of mass fraction of
polymer at the 35:65 powder ratio resulting in the lowest
dynamic-viscosity value of 321.12 MPa s.
4 CONCLUSION
The Al2O3 powder characterized by irregularly
shaped particles with sharp edges demonstrates the
ability to agglomerate, resulting in a non-uniform par-
ticle-size distribution. The Zahn cup viscosity (7.35 s) is
slightly larger in comparison to water viscosity (5.83 s),
thus the Al2O3 particle dispersion in the binder is faci-
litated.
In addition, a relatively large content of the solid
phase in a slurry reduces the coat shrinkage during the
drying process and enhances its strength. The properties
of the coating surface may be improved by increasing the
plate weight. An addition of a poly acrylic polymer at
the lowest content to the alumina powders with various
granulation values allows a regulation of the rheological
properties of the ceramic slurry towards more effective
ceramic shell-mould fabrication.
The investigated slurries show standard features in
the investment-casting process on an industrial scale.
They are prospective for future shell-mould fabrication.
Acknowledgement
The financial support from the Structural Funds for
the Operational Programme Innovative Economy (IE
OP) provided by the European Regional Development
Fund – Project "Modern material technologies in aero-
J. SZYMAÑSKA et al.: RHEOLOGICAL PROPERTIES OF ALUMINA CERAMIC SLURRIES FOR CERAMIC ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 735–738 737
Figure 6: Relation between viscosity and shear rate of Al2O3 ceramic
slurries with 72.5 % of mass fractions of solid content for 15 % of
mass fractions of polymer content at two powder ratios, 35:65 and
65:35 (200:030 mesh)
Slika 6: Odvisnost med viskoznostjo in stri`no hitrostjo Al2O3
kerami~ne go{~e z 72,5 % masnim dele`em trdnega pri 15 % masnega
dele`a polimera, pri dveh razmerjih zrnatosti prahov 35:65 in 65:35,
mre`a (200:300)
Figure 4: Relation between viscosity and shear rate of Al2O3 ceramic
slurries with 72.5 % of mass fractions of solid content for 6 % of mass
fractions of polymer content at two powder ratios, 35:65 and 65:35
(200:030 mesh)
Slika 4: Odvisnost med viskoznostjo in stri`no hitrostjo Al2O3
kerami~ne go{~e z 72,5 % dele`em trdnega pri 6 % masnega dele`a
polimera, pri dveh razmerjih zrnatosti prahov 35:65 in 65:35, mre`a
(200:030)
Figure 5: Relation between viscosity and shear rate of Al2O3 ceramic
slurries with 72.5 % of mass fractions of solid content for 10 % of
mass fractions of polymer content at two powder ratios, 35:65 and
65:35 (200:030 mesh)
Slika 5: Odvisnost med viskoznostjo in stri`no hitrostjo Al2O3
kerami~ne go{~e z 72,5 % masnim dele`em trdnega pri 10 % masnega
dele`a polimera, pri dveh razmerjih zrnatosti prahov 35:65 in 65:35,
mre`a (200:030)
Figure 3: Coating thickness (H) dependence on time for the slurries
with 72.5 % of mass fractions of solid content for (6, 10, 15) % of
mass fractions of polymer content at two powder ratios, 35:65 and
65:35 (200:030 mesh)
Slika 3: Debelina nanosa (H) v odvisnosti od ~asa pri go{~i z 72, 5 %
masnim dele`em trdnega in pri vsebnosti (6, 10 , 15) % masnega
dele`a polimera, pri dveh razmerjih prahov 35:65 in 65:35, mre`a
(200:030)
space industry", Nr POIG.01.01.02-00-015/08-00, is
gratefully acknowledged.
5 REFERENCES
1 S. Roskosz, Relationship between mould’s technology and structure
of investment cast nickel based superalloys, In¿ynieria Materia³owa,
29 (2008) 4, 375–379, doi:bwmeta1.element.baztech-article-
BPL8-0006-0070
2 H. Matysiak, J. Ferenc, J. Michalski, Z. Lipiñski, G. Jakubowicz, K.
J. Kurzyd³owski, Porosity and strength of ceramic shell moulds used
in investment casting process by Bridgman method, In¿ynieria
Materia³owa, 32 (2011) 1, 17–21, doi:bwmetal.element.baztech-
cle4cb00-4f01-476e-8401-81df9ad0e967
3 J. Raabe, E. Bobryk, Functional Ceramics, Oficyna Wydawnicza
Politechniki Warszawskiej, 1997
4 R. Haratym, Investment casting processes for ceramic shell moulds,
Warszawa 1997
5 S. Jones, C. Yuan, Advances in shell moulding for investment cast-
ing, Journal of Materials Processing Technology, 135 (2003) 2–3,
258–265, doi:10.1016/S0924-0136(02)00907-X
6 M. Zagórska, P. Wiœniewski, H. Matysiak, K. Kwapiszewska, J.
Ferenc-Dominik, J. Michalski, K. J. Kurzyd³owski, The influence of
polymer binder, based on nano-Al2O3 dispersion, on the properties of
ceramic slurries used in the investment casting, Euromat, 2011
7 M. R. Ismael, R. D. Dos Anjos, R. Salomao, V. C. Pandolffelli,
Colloidal silica as a nanostructured binder for refractory castables,
Refractories Applications and News, 11 (2006) 4, 16–20
J. SZYMAÑSKA et al.: RHEOLOGICAL PROPERTIES OF ALUMINA CERAMIC SLURRIES FOR CERAMIC ...
738 Materiali in tehnologije / Materials and technology 50 (2016) 5, 735–738
D. KÝRSEVER et al.: EFFECT OF MECHANICAL ACTIVATION ON THE SYNTHESIS OF A MAGNESIUM ...
739–742
EFFECT OF MECHANICAL ACTIVATION ON THE SYNTHESIS
OF A MAGNESIUM ALUMINATE SPINEL
VPLIV MEHANSKE AKTIVACIJE NA SINTEZO
MAGNEZIJ-ALUMINATNEGA [PINELA
Derya Kýrsever, Nilgün Kaya Karabulut, Nuray Canikoðlu, Hüseyin Özkan Toplan
Sakarya University, Metallurgy and Materials Engineering, 54187 Sakarya, Turkey
dkirsever@sakarya.edu.tr
Prejem rokopisa – received: 2015-07-08; sprejem za objavo – accepted for publication: 2015-09-09
doi:10.17222/mit.2015.209
A magnesium aluminate spinel powder (72 % Al2O3 & 28 % MgO) was prepared with mechanical activation. Samples were
sintered in a temperature range of 1400–1750 °C. The final sintered products were characterized with densification, phase and
microstructural analyses and a hardness measurement to evaluate the influence of mechanical activation on the synthesis of a
magnesium aluminate spinel.
Keywords: mechanical activation, magnesium aluminate spinel, ceramic, sintering, densification, mechanical properties
Prah magnezij aluminatnega {pinela (72 % Al2O3 & 28 % MgO) je bil pripravljen z mehansko aktivacijo. Sintranje vzorcev je
bilo izvedeno v temperaturnem obmo~ju 1400–1750 °C. Na kon~nih sintranih vzorcih je bila dolo~ena zgostitev, opravljena je
bila analiza faz in mikrostrukture ter meritev trdote, da bi ocenili vpliv mehanske aktivacije na sintezo magnezij- aluminatnega
{pinela.
Klju~ne besede: mehanska aktivacija, magnezij-aluminatni {pinel, keramika, sintranje, zgo{~evanje, mehanske lastnosti
1 INTRODUCTION
Magnesium aluminate spinel (MgAl2O4, MA) is a
widely used refractory material due to its high-tempera-
ture properties, mechanical resistance, thermal-shock re-
sistance and high corrosion resistance to acidic and basic
slags.1,2 MA spinel has a high melting point (2135 °C),
high hardness (16.1 GPa), relatively low density (3.58
g/cm3), high strength (180 MPa) at room and at elevated
temperatures, high chemical inertness, a low thermal-ex-
pansion coefficient (9 × 10–6/°C between 30 °C and
1400 °C) and high thermal-shock resistance.3Also, MA
spinel refractories are very attractive due to their envi-
ronmental friendliness, contrary to magnesium chromite
refractories. However, aspinel formation is accompanied
by a 5–7 % volume expansion which does not allow it to
densify in a single-stage firing.4 Therefore, the synthesis
of spinel and fabrication of spinel refractories were not
feasible with commercial methods due to the difficulty
with sintering.1,2 High-purity spinel was synthesized
mostly with hydrothermal techniques, sol-gel, spray
plasma, cool drying, controlled hydrolysis, co-precipi-
tation, mechanical activation and the aerosol method.1
Mechanical activation is a method, which can induce
changes tothe solid-state properties, such as the distor-
tion of the structure, accompanied by the accumulation
of energy and the formation of active centers on the
newly formed surfaces.5 Different processes can remark-
ably influence the reactivity of solids. Mechanical treat-
ments are particularly important as long as they can help
to produce changes to the texture and structure of the
solids. In many cases, these alterations tothe structure
cause certain modifications to the phases formed due to
the thermal treatment of the solids, which were
mechanochemically treated.6
The main aim of this study was to prepare a magne-
sium aluminate spinel by firing between 1400–1750 °C
and to analyze the effect of mechanical activation. Inves-
tigations of the phases, crystal morphology and densifi-
cation of the fired products were carried out. In addition,
the hardness values of the samples for different sintering
temperatures were studied.
2 EXPERIMENTAL DETAILS
Al2O3 (72 % of mass fractions) and MgO (28 % of
mass fractions) powders were ball milled with alumina
balls in a polyethylene bottle for 1 h. The mixture was
made in a high-energy planetary ball mill (Fristch) at a
rotation speed of 600 min–1. The ball-to-powder weight
ratio was adjusted to 20. Milling of the precursor was
carried out for 1 h. Activated and non-activated powders
were uniaxially pressed to form pellets at 255 MPa. The
pellets were sintered in the temperature range of
1400–1750 °C for 1 h. An X-ray diffraction analysis was
performed using a Rigaku Ultima X-ray diffractometer.
A Joel 6060 LV scanning electron microscope was used
for the morphological analysis of the non-activated and
activated powders and sintered samples. The hardness
Materiali in tehnologije / Materials and technology 50 (2016) 5, 739–742 739
UDK 67.017:621.762:669.721:661.862’027 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)739(2016)
measurements of the samples were done using Leica
Microsysteme GmbH. The apparent porosity and bulk
density of the sintered samples were measured with the
liquid-displacement method using Archimedes’principle.
Water absorption was alsoinvestigated.
3 RESULTS AND DISCUSSION
Figure 1 shows SEM micrographs of non-activated
and activated powder mixtures. The non-activated pow-
der mixture has well-defined faces and edges. However,
the particle size decreases and the particle shape be-
comes round with mechanical activation.
Figure 2 shows XRD patterns of the non-activated
and activated powder mixtures of MgO and Al2O3. As a
result, Mg(OH)2, Al2O3 and MgO peaks are observed.
SEM micrographs of the fractured surfaces of all the
samples sintered in the temperature range of
1400–1750 °C for 2 h are shown in Figure 3. It can be
seen that all the samples appear to be relatively dense
D. KÝRSEVER et al.: EFFECT OF MECHANICAL ACTIVATION ON THE SYNTHESIS OF A MAGNESIUM ...
740 Materiali in tehnologije / Materials and technology 50 (2016) 5, 739–742
Figure 1: SEM micrographs of powder mixtures: a) non-activated, b) activated for 1 h
Slika 1: SEM-posnetka me{anice prahu: a) neaktiviran, b) aktiviran 1 h
Figure 3: SEM micrographs of all sintered samples prepared from: a), b), c), d), e) non-activated and f), g), h), i), j) activated powder mixtures
Slika 3: SEM-posnetki vseh sintranih vzorcev pripravljenih iz: a), b), c), d) , e) neaktivirana in f), g), h), i), j) aktivirana me{anica prahu
Figure 2: XRD patterns of non-activatedand activated powder mix-
tures of MgO and Al2O3(A: Al2O3, B: Mg(OH)2, P: MgO)
Slika 2: Rentgenograma neaktivirane in aktivirane me{anice prahu
MgO in Al2O3 (A: Al2O3, B: Mg(OH)2, P: MgO)
with the increasing sintering temperature and mechanical
activation. After the sintering at 1700 °C, spinel grain
growth was found for the non-activated and activated
samples. However, porosity levels seem relatively higher
for the non-activated samples.
Figure 4 shows the XRD patterns of the non-acti-
vated and activated samples after the sintering in the
temperature range of 1400–1750 °C for 2 h. These con-
firm that the Mg-Al spinel is the only phase of the acti-
vated samples. However, Al2O3 peaks are also present at
1400 °C for the non-activated samples.
Figure 5 summarizes the bulk density and apparent
porosity of all the sintered samples. A general trend in
the increasing bulk density and decreasing apparent po-
rosity with the increasing sintering temperature was ob-
served. The mechanically activated samples showed rela-
tively higher density values compared to those of the
non-activated samples. On the other hand, the mechani-
cally activated samples obtained a lower apparent poros-
ity than the non-activated samples. The sintered bulk
density and apparent porosity ofthenon-activated and ac-
tivated samples are 3.76 g/cm3 and 3.1 g/cm3, respec-
tively.
The reduction of the particle size decreases the dis-
tance between the vacancy sites (or betweenthegrain
boundaries) and enhances the vacancy diffusion to the
surface, thus increasing the densification. The reduction
of the particle size is obtained with mechanical activa-
tion.7
Figure 6 shows the water absorption of all the sin-
tered samples. It can be seen that the water absorption
decreases with the sintering temperature and mechanical
activation. For the activated samples, there is no relative
water absorption above 1600 °C.
Figure 7 shows the hardness resultsfor the non-acti-
vated and activated samples with respect to the sintering
temperature. The highest hardness value of an activated
sample is 1623 HV at 1650 °C. Also, the hardness values
of the activated samples arehigher than those for the
D. KÝRSEVER et al.: EFFECT OF MECHANICAL ACTIVATION ON THE SYNTHESIS OF A MAGNESIUM ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 739–742 741
Figure 4: XRD patterns of: a) non-activated and b) activated samples sintered at different temperatures for 2 h(M: MgAl2O4, A: Al2O3)
Slika 4: Rentgenogrami: a) neaktivirani in b) aktivirani vzorci sintrani 2 h na razli~nih temperaturah (A: Al2O3, M: MgAl2O4)
Figure 6: Hardness measurements forall the sintered samples
Slika 6: Meritve trdote vseh sintranih vzorcev
Figure 5: Bulk-density and apparent-porosity plots for all the sintered
samples
Slika 5: Diagrama gostote osnove in navidezne poroznosti vseh sin-
tranih vzorcev
non-activated samples. These results arerelated to
thedensification and a low porosity level.
4 CONCLUSIONS
Magnesium aluminate spinel samples were synthe-
sized using mechanical activation. The activated samples
resulted in higher density and hardness values in com-
parison with the non-activated samples. In addition,
theactivated samples exhibited dense grains and a low
porosity. So, mechanical activation can facilitate a sin-
gle-stage sintering process and greatly influence the
costs of the production.
5 REFERENCES
1 P. Orosco, L. Barbosa, M. C. Ruiz, Synthesis of magnesium alu-
minate spinel by periclase and alumina chlorination, Materials Re-
search Bulletin, 59 (2014), 337–340, doi:10.1016/j.materresbull.
2014.07.026
2 P. G. Lampropoulou, C. G. Katagas, Effects of zirconium silicate and
chromite addition on the microstructure and bulk density of magne-
sia–magnesium aluminate spinel-based refractory materials, Ce-
ramics International, 34 (2008), 1247–1252, doi:10.1016/j.ceramint.
2007.03.015
3 I. Ganesh, Fabrication of magnesium aluminate (MgAl2O4) spinel
foams, Ceramics International, 37 (2011), 2237–2245, doi:10.1016/
j.ceramint.2011.03.068
4 H. S.Tripathi, B. Mukherjee, S. Das, M. K. Haldar, S. K. Das,
A.Ghosh, Synthesis and densification of magnesium aluminate
spinel: effect of MgO reactivity, Ceramics International, 29 (2003),
915–918, doi:10.1016/S0272-8842(03)00036-1
5 E. Turianicová, A. Obut, ¼. Tu~ek, A. Zorkovská, Ý. Girgin, P. Balá`,
Interaction of natural and thermally processed vermiculites with gas-
eous carbon dioxide during mechanical activation, Applied Clay Sci.,
88–89 (2014), 86–91, doi:10.1016/j.clay.2013.11.005
6 S. Koç, N. Toplan, K. Yýldýz, H. Ö. Toplan, Effects of mechanical
activation on the non-isothermal kinetics of mullite formation from
kaolinite, J. Therm. Anal. Calorim., 103 (2010), 791–796,
doi:10.1007/s10973-010-1154-5
7 R.Sarkar, S.Kumar Das, G.Banerjee, Effect of attritor milling on the
densification of magnesium aluminate spinel, Ceramics Interna-
tional, 25 (1999), 485–489, doi:10.1016/S0272-8842(98)00065-0
D. KÝRSEVER et al.: EFFECT OF MECHANICAL ACTIVATION ON THE SYNTHESIS OF A MAGNESIUM ...
742 Materiali in tehnologije / Materials and technology 50 (2016) 5, 739–742
Figure 7: Water absorption of all the sintered samples
Slika 7: Absorpcija vode vseh sintranih vzorcev
K. ZUPAN et al.: PHASE AND MICROSTRUCTURE DEVELOPMENT OF LSCM PEROVSKITE MATERIALS ...
743–748
PHASE AND MICROSTRUCTURE DEVELOPMENT OF LSCM
PEROVSKITE MATERIALS FOR SOFC ANODES PREPARED BY
THE CARBONATE-COPRECIPITATION METHOD
RAZVOJ KRISTALNIH FAZ IN MIKROSTRUKTURE LSCM
PEROVSKITNIH MATERIALOV ZA SOFC ANODE,
PRIPRAVLJENIH S KARBONATNO METODO KOPRECIPITACIJE
Klementina Zupan, Marjan Marin{ek, Tina Skalar
University of Ljubljana, Faculty of Chemistry and Chemical Technology, Ve~na pot 113, 1000 Ljubljana, Slovenia
klementina.zupan@fkkt.uni-lj.si
Prejem rokopisa – received: 2015-07-21; sprejem za objavo – accepted for publication: 2015-10-09
doi:10.17222/mit.2015.232
Most SOFC development has been based on nickel yttria-stabilized zirconia anodes. Such materials have excellent catalytic
properties for fuel oxidation, high electrical conductivity, good mechanical strength and an appropriate thermal expansion coef-
ficient compatible with other cell components. Unfortunately, cermet anodes based on doped zirconia exhibit some disadvan-
tages, e.g., the catalysing side reaction of carbon deposition during hydro-carbon fuel oxidation and a susceptibility to sulphur
poisoning. Perovskite-type compounds based on lanthanum-strontium-manganese-chromium oxide (LSCM) can serve as an al-
ternative material. Since the optimal perovskite composition is still not known, La1–xSrxMnyCr1.yO3± (x from 0 to 0.3 and y from
0.4 to 0.6) ceramics were prepared with the co-precipitation method. Crystalline phase formation was followed by X-ray powder
diffraction and Rietveld refinement. Quantitative microstructure analysis of the samples sintered at various temperatures was
performed on SEM micrographs using Axiovision 4.8 software.
Keywords: co-precipitation, oxide LSCM anode, phase development, microstructure
Ve~ina razvoja visokotemperaturnih gorivnih celic je temeljila na anodnih materialih na osnovi niklja in cirkonijevega dioksida,
stabiliziranega z itrijem. Ta ima odli~ne katalitske lastnosti pri reakciji oksidacije goriva, visoko elektri~no prevodnost, dobro
mehansko trdnost in temperaturni razteznostni koeficient, skladen z ostalimi komponentami celice. @al so ti materiali med
delovanjem podvr`eni ne`elenim reakcijam izlo~anja ogljika in zastrupljanja z `veplom, zato jih posku{amo nadomestiti z
oksidnimi spojinami perovskitnega tipa, z lantan-stroncij-mangan-krom oksidom (LSCM). Optimalna sestava teh materialov {e
ni znana, zato smo z metodo soobarjanja pripravili keramiko La1–xSrxMnyCr1.yO3±  (x od 0 do 0,3 in y od 0,4 do 0,6). Z
rentgensko pra{kovno analizo in Rietveldovim prilagajanjem smo spremljali razvoj kristalnih faz. Z analizo SEM posnetkov
vzorcev po sintranju pri razli~nih temperaturah smo mikrostrukture pripravljenih materialov kvantitativno ovrednotili z uporabo
programa Axiovision 4.8.
Klju~ne besede: kopercipitacija, oksidna LSCM anoda, razvoj faz, mikrostruktura
1 INTRODUCTION
Fuel cells can be considered as devices that electro-
chemically convert fuels into electricity or, more pre-
cisely, batteries with permanent fuel supplies. Solid-ox-
ide fuel cells (SOFCs), based on an ion-conducting
electrolyte, have several advantages over other types of
fuel cells, including their potential fuel flexibility and
very high chemical-to-electrical conversion efficiency
due to the absence of Carnot limitations. Further energy
gains can be achieved in SOFC systems when cogene-
rated heat is used for the internal reforming of methane
or other hydrocarbon fuels directly on the anode.1
Porous Ni/YSZ-based materials are conventionally
used as SOFC anodes due to their high electrical conduc-
tivity, activity for electrode electro-chemical oxidation,
stability under reduced environmental conditions, appro-
priate thermal expansion and a chemical compatibility
with other cell components.2 Despite the many advan-
tages they possess, there are some drawbacks, such as a
low tolerance to sulphur impurities3 and a tendency to
coke when hydrocarbons are used as fuels4.
Alternative cermet anode materials have been
extensively studied, such as Cu-CeO2-YSZ. Researchers
have demonstrated their operation using various fuels.5,6
Recently, another alternative approach aiming to prepare
all-oxide anode materials has been proposed in order to
develop electrodes that exhibit catalytic, electron- and
ion-conducting properties. Many problems with all-oxide
anodes have been overcome with the introduction of
novel perovskite-structure materials with the general
formula ABO3–. The A element is typically lanthanide,
while the B element is the transition metal. In principle,
catalytic activity toward fuel oxidation, electron- and
ion-conductivity can be tailored by a wide range of
doping elements. The oxidation states of the A-site and
B-site cations determine the oxygen vacancy
concentration .2 Among various perovskites, a
particularly complex metal oxide with the composition
La0,75Sr0,25Mn0,5Cr0,5O3– has attracted much attention as a
Materiali in tehnologije / Materials and technology 50 (2016) 5, 743–748 743
UDK 67.017:661.8’02:621.3.032.22:537.533 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)743(2016)
promising anode material, due to its good catalytic activ-
ity, excellent redox stability, reduced carbon deposition
susceptibility and improved sulphur-poisoning stability.7
An introduction of alkaline earth ions, i.e., Mg2+, Ca2+,
and Sr2+, into the A site of lanthanum chromite can en-
hance the electrical conductivity by two orders of magni-
tude.7 L. Deleebeck et al.9 first demonstrated that the
removal of Sr from La1–xSrxMn1–yCryO3± improves the
thermo-chemical stability and the electronic conductivity
in a humidified H2 atmosphere. In their second study,
they reported that the catalytic activity toward H2 oxida-
tion decreases with increasing Cr content (y = 0.4–0.6),
while the relatively high Sr content (x = 0.2) shows a
lower catalytic activity.10 The optimal LSCM material
composition is not yet known.
Various chemical routes to prepare LSCM powders
have been reported, including the solid-state reaction,7
the chelating method,11 gel casting,12 and combustion
synthesis.13–16 However, the synthesis of single-phase
LSCM composed of fine powders requires a further im-
provement. Co-precipitation is a promising and simple
chemical method to prepare well-defined and less-ag-
glomerated perovskite powders. It was reported that the
most significant synthesis parameter for LSCM prepara-
tion via co-precipitation is the pH value of the reaction
mixture, which should be maintained slightly below 8 in
order to ensure that all the cations precipitate.17 In addi-
tion to an appropriate chemical composition of the
LSCM material, the electrode performance in an operat-
ing SOFC is also essentially dependent on the electrode
microstructure, final porosity and potential presence of
secondary phases.
In this work, we applied the "reverse strike" carbon-
ate co-precipitation method for batch
La1–xSrxMn1–yCryO3± perovskite preparation in which the
Sr content and the Cr-to-Mn molar ratio were varied. The
aim of this work is to describe the relationship between
the microstructure parameters and the LSCM composi-
tion using various analytical techniques.
2 EXPERIMENTAL PROCEDURE
La1–xSrxMn1–yCryO3± (x = 0, 0.1, 0.2 or 0.3 and y =
0.4, 0.5 or 0.6) oxides were prepared using co-precipita-
tion synthesis (Table 1). La(NO3)3⋅6H2O (99 %),
Sr(NO3)2 (98 %), Cr(NO3)3⋅9H2O (98.5 %) and
Mn(NO3)2⋅4H2O (98 %), all from Alfa Aesar, were used
as the source of metal ions. The carbonate precursors
were prepared using the "reverse strike" method in which
a mixed metal nitrate solution is added to a precipitant
carbonate solution to achieve a more uniform cation dis-
tribution by instantaneous precipitation. A total of
600 mL of 0.125-M aqueous solution of (NH4)2CO3 was
poured into a jacket glass reactor (1.25 L); 0.5-M metal
nitrates solutions were prepared, as were adequate vol-
umes of each LSCM component regarding the desired fi-
nal LSCM composition. The solutions were mixed to-
gether and dripped into a stirring precipitant solution.
The precipitating solution was kept at 60 °C under a CO2
protective atmosphere to prevent manganese oxidation
during synthesis. The pH inside the jacket glass reactor
was kept at 7.8±0.1 by the periodic addition of ammonia
(25 %, aq.). Afterwards, the precipitate was filtered off
under a CO2 environment and washed three times
(50 mL) with a 0.125-M solution of (NH4)2CO3, dried for
6 h at 110 °C and finally calcined at 1000 °C in an air at-
mosphere.
Table 1: Compositions and sample notations
Tabela 1: Sestave in poimenovanje vzorcev
Sample composition Sample name
La0,7Sr0,3Cr0,5Mn0,5O3–d La7Cr5
La0,8Sr0,2Cr0,5Mn0,5O3–d La8Cr5
La0,9Sr0,1Cr0,5Mn0,5O3–d La9Cr5
La1Cr0,5Mn0,5O3–d La10Cr5
LaCr0,6Mn0,4O3 La10Cr6
La0,9Sr0,1Cr0,6Mn0,4O3 La9Cr6
La0,9Sr0,1Cr0,4Mn0,6O3 La9Cr4
LaCr0,4Mn0,6O3 La10Cr4
The synthesized powders were milled in an agate
mortar and un-axially pressed into pellets (100 MPa) and
sintered at various temperatures (1250 °C, 1300 °C, 1400 °C
and 1500 °C) for 1 h. The calcined and sintered samples
were analysed with a PANalytical X’Pert PRO MPD
apparatus. For the determination of the microstructure,
the sintered tablets were polished (diamond pastes of 3
μm and 0.25 μm), thermally etched, and subsequently
analysed with a FE-Zeiss ULTRA Plus SEM. The
quantitative analyses of the microstructures were per-
formed on digital images (images were digitized into
pixels with 255 different grey values) using Axiovision
4.8 image-analysis software.
3 RESULTS AND DISCUSSION
The carbonate co-precipitation route is an appropriate
method for the preparation of complex metal oxides,
such as La1–xSrxMn1–yCryO3± (LSCM). When the "re-
versed strike" co-precipitation method is used and the
mixed solution of metal ions drips into the concentrated
precipitant solution, the various cations within each
droplet precipitate almost instaneously.18 Lanthanum car-
bonate precipitated at pH > 4.2, manganese carbonate at
pH > 5 and strontium carbonate at pH > 7.3. Chromium
precipitates as a hydroxide in a very narrow pH range
from 6.6 to 7.3; however, Cr(OH)3 starts to dissolve at a
pH value of 7.9.17 Therefore, the precipitation of mixed
metal oxide should be carried out carefully in the tiny pH
range from 7.3 and 7.9. If we take into account only sim-
ple carbonate and hydroxide species (Equations (1) to
(4)) for the calculation for the stoichiometric amount of
ammonium carbonate as a precipitant agent, we can con-
clude that an excess of 50 % is used during the precipita-
K. ZUPAN et al.: PHASE AND MICROSTRUCTURE DEVELOPMENT OF LSCM PEROVSKITE MATERIALS ...
744 Materiali in tehnologije / Materials and technology 50 (2016) 5, 743–748
tion process. Thus, the super-saturation ratio in the case
of the LSCM synthesis calculated from the molar ratio of
ammonium carbonate and total metal ions is 1.5.
3(NH4)2CO3 + 2La
3+ 	 La2(CO3)3 + 6NH4
+ (1)
(NH4)2CO3 + Sr
2+ 	 SrCO3 + 2NH4
+ (2)
3(NH4)2CO3 + Cr
3+ + H2O 	 Cr(OH)3 + 6NH4
+ +
+ 3HCO3
– (3)
(NH4)2CO3 + Mn
2+ 	 MnCO3 + 2NH4
+ (4)
According to Figure 1, the perovskite LSCM phase
formation for all the samples is practically complete after
calcination at 1000 °C for 1 h (the perovskite peaks are
denoted with a letter "P"). The main perovskite phase is
quite well crystallised. In the sample La7Cr5, with the
highest Sr content and equal amounts of chromium and
manganese, the XRD analysis revealed a small amount
of strontium secondary phase SrCrO4 (denoted with the
letter "S"). It is described in the literature that in the hu-
midified hydrogen atmosphere SrCrO4 further transforms
into the Ruddlesden-Popper phase Sr2CrO4.9 Addi-
tionally, the X-ray powder diffraction indicates that in
samples with the lowest Cr content (0.4), a lanthanum-
rich secondary phase La2CrO6 is formed (denoted with
the letter "L"). Varying the Sr content in this Cr-poor
sample reveals that in the Sr-free and Sr = 0.1 samples
the La2CrO6 content determined according to the Riet-
veld refinement is 3.1 % and 7.6 %, respectively. Multi-
ple RTG peaks observed in some patterns are a conse-
quence of a perovskite lattice superstructure. This lattice
superstructure is formed due to the octahedron tilting,
which has its origin in the random Sr-incorporation into
the perovskite structure.19,20 By doping lanthanum-man-
ganite with Sr and Cr, the symmetry of the structure is
lowered due to the different sizes of the introduced cat-
ions compared to the original ions. Consequently, octa-
hedrons defined by a central cation (B-site cation) and
the surrounding oxygen ions are slightly tilted and the
repeating structure pattern is defined by eight original
unit cells.
The secondary phases are highly undesired in the fi-
nal LSCM since they result in the so-called layered
perovskite structure with additional layers of Sr-oxide,
La-oxide, or a mixture of both separating the LSCM at
temperatures around 1100 °C and in a H2 atmosphere.
Furthermore, the secondary phases decrease the thermo-
chemical stability, catalytic activity and electrical con-
ductivity of the LSCM.9,12
K. ZUPAN et al.: PHASE AND MICROSTRUCTURE DEVELOPMENT OF LSCM PEROVSKITE MATERIALS ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 743–748 745
Figure 2: X-ray diffraction patterns of all samples after calcination at
1250 °C
Slika 2: Rentgenogram vseh vzorcev po kalcinaciji pri 1250 °C
Figure 1: X-ray diffraction patterns of prepared samples after calcina-
tion at 1000 °C
Slika 1: Rentgenogram pripravljenih vzorcev po kalcinaciji pri
1000 °C
Figure 3: SEM micrographs of all samples sintered at 1250 °C and
sample La9Cr6 sintered at 1250 °C, 1300 °C, 1400 °C and 1500 °C
Slika 3: SEM-posnetki vseh vzorcev po sintranju pri 1250 °C in vzor-
ca La9Cr6, sintranega pri 1250 °C, 1300 °C, 1400 °C in 1500 °C
After the sintering at 1250 °C, the only phase present
in all the samples is the LSCM perovskite, as shown in
Figure 2. The absence of secondary phases indicates that
they re-dissolve in the main LSCM phase when the
sintering temperature is increased. This discovery also
gives a very effective tool for controlling the amount of
secondary phases in LSCM or even eliminates them
completely. Due to the possibility of eliminating the sec-
ondary phases from the LSCM, it is reasonable to con-
clude that the co-precipitation method offers an essential
advantage over the synthesis processes that are based on
the solid-state reactions in which local inhomogeneity in
chemical composition are quite common.
One of the major challenges in applying LSCM mate-
rial as an anode in ceramic fuel cells is achieving the
continuity of the electrode material as well as the conti-
nuity of the pores. Good contact between the particles is
critical for forming continuous paths throughout the
formed anode, reaching a high conductivity. The sinter-
ing behaviour of all the samples after sintering at 1250 °C
is demonstrated in Figure 3. Since all the microstructure
parameters that are important for an exact anode analysis
are sometimes difficult to deduce simply from the SEM
micrographs, a detailed quantitative microstructure
analysis of sintered samples is performed. For statisti-
cally reliable data in each case, 5 to 10 different regions
were analyzed. The results of the quantitative micro-
structure analysis are summarized in Table 2. The
parameters d , dx, dy and 
 are represented as the diameter
of the area-analogue circle – Dcircle, the intercept lengths
in the x and y directions – FERETX, FERETY and the
Shape factor fcircle. Spore and FERETMAX are determined as
the pore areas and maximum intercept lengths of the
pore while rel and  are determined from the geometric
densities of the tablets and the theoretical densities that
were calculated using the Rietveld refinement.
SEM micrographs sintered at 1250 °C reveal that the
microstructure parameters greatly depend on the sample
composition. When comparing samples with equal
amounts of Cr and Mn with various Sr contents we can
observe that the presence of strontium promotes sinter-
ing. It is also evident that the grains are larger in samples
K. ZUPAN et al.: PHASE AND MICROSTRUCTURE DEVELOPMENT OF LSCM PEROVSKITE MATERIALS ...
746 Materiali in tehnologije / Materials and technology 50 (2016) 5, 743–748
Table 2: Results of quantitative microstructure analysis of the sintered samples
Tabela 2: Rezultati kvantitativne analize mikrostrukture sintranih vzorcev
T/°C /% dy/μm dx/μm d/μm 
 Spore/μm2 FERETMAX/μm
teor /
g cm–3 rel /%
La10Cr5
1250 51.9 0.30 0.31 0.28 0.67 0.44 1.02
6.71
48.1
1300 50.9 0.43 0.44 0.41 0.74 0.86 1.66 49.1
1400 40.1 1.00 1.01 0.92 0.73 1.31 2.03 59.9
1500 27.4 1.85 1.91 1.70 0.70 1.27 1.82 72.6
La9Cr5
1250 44.6 0.52 0.53 0.48 0.70 0.77 1.69
6.60
55.4
1300 41.5 0.60 0.61 0.55 0.71 0.72 1.60 58.5
1400 25.4 1.10 1.11 1.02 0.74 0.58 1.26 74.6
1500 9.9 2.89 2.92 2.67 0.69 0.39 0.65 90.1
La8Cr5
1250 43.7 0.50 0.49 0.45 0.70 0.36 1.05
6.56
56.3
1300 40.3 0.62 0.60 0.55 0.67 0.20 0.76 59.7
1400 27.0 1.22 1.25 1.14 0.74 0.77 1.46 73.0
1500 16.0 2.54 2.65 2.35 0.72 0.65 1.12 84.3
La7Cr5
1250 45.9 0.56 0.58 0.53 0.72 0.37 0.80
6.45
54.1
1300 42.1 0.86 0.86 0.79 0.76 0.74 1.04 57.9
1400 31.2 1.23 0.86 1.14 0.75 0.66 1.47 68.8
1500 21.8 1.94 2.00 1.79 0.73 0.39 0.87 78.2
La10Cr6
1250 54.2 0.44 0.45 0.42 0.72 1.12 1.42
6.72
45.8
1300 48.1 0.76 0.77 0.72 0.75 1.33 1.92 51.9
1400 37.9 1.20 1.21 1.13 0.77 1.33 2.06 62.1
1500 23.4 1.97 1.96 1.84 0.76 1.94 2.31 76.6
La9Cr6
1250 54.0 0.24 0.23 0.22 0.79 0.37 1.07
6.57
46.0
1300 50.5 0.39 0.39 0.36 0.79 3.10 2.73 49.5
1400 40.7 0.67 0.69 0.63 0.76 2.54 2.94 59.3
1500 38.6 1.70 1.73 1.60 0.78 1.35 2.02 61.4
La10Cr4
1250 44.5 0.40 0.41 0.38 0.78 0.59 1.11
6.73
55.5
1300 39.2 0.77 0.78 0.72 0.71 0.36 1.01 60.8
1400 31.6 1.21 1.19 1.11 0.72 0.83 1.42 68.4
1500 24.0 2.15 2.09 1.94 0.70 1.11 1.47 76.0
La9Cr4
1250 54.1 0.43 0.43 0.40 0.72 2.63 2.76
6.59
45.9
1300 50.7 0.69 0.67 0.64 0.73 1.65 2.43 49.3
1400 34.1 1.43 1.44 1.33 0.76 2.30 2.65 65.9
1500 22.5 2.68 2.60 2.44 0.75 1.13 1.18 77.5
that contain Sr than in the Sr-free sample. Although no
secondary phases are detected in the Sr-free sample
(La10Cr5), the microstructure is composed of regions
with larger and smaller grains. All Sr-free samples with a
lower Cr-content behave similarly. The presence of a
secondary phase SrCrO4 in the sample with the highest
Sr content (La7Cr5) evidently does not have a significant
effect on the LSCM grain size distribution. In contrast,
the presence of a lanthanum-rich secondary phase
La2CrO6 in samples with a lower Cr content (La10Cr4
and La9Cr4) causes non-homogenous microstructures
with the denser regions containing larger grains and the
less dense regions containing smaller grains. In the
phase-pure LSCM sample with a Cr content y = 0.6
(La9Cr6), a very interesting microstructure is formed af-
ter sintering at 1250 °C. In this sample, the formation of
well-connected fine grains is observed.
The results of a quantitative microstructural analysis
are in good agreement with optical observations. The rel-
ative sintered density increases and the porosity de-
creases with the increasing sintering temperature. Sam-
ples with a Sr content x = 0.1, 0.2 or 0.3 and equal
contents of Cr and Mn sinter at the lowest sintering tem-
perature of 1250 °C to somewhat higher densities than
the Sr-free sample (La10Cr5). At the same sintering tem-
perature, rel reaches 48.1 %, 55.4 %, 56.3 % and 54.1 %
for the samples La10Cr5, La9Cr5 La8Cr5 and La7Cr5,
respectively. The addition of strontium to the perovskite
also results in grain growth, during which grains reach
an average size of 0.28 μm at 1250 °C in a Sr-free sam-
ple, while for the highest Sr content in sample x = 0.3 the
average grain size grew to 0.53 μm. At the highest sinter-
ing temperature (1500 °C), rel. reaches 90.1 %, 72.6 %,
84.3 % and 78.2 % for the samples La9Cr5, La10Cr5,
La8Cr5 and La7Cr5, respectively. From this fact, it can
be deduced that the addition of Sr to some amount x =
0.1 accelerates sintering, while adding Sr to perovskite
above a certain concentration x = 0.2 and 0.3 supresses
the densifying process. At a lower Sr concentration,
SrCrO4 (according to phase diagram SrO–Cr2O3) forms a
liquid phase due to eutectic and peritectic reactions21 that
promote sintering.22 In principle, a higher Sr-content in-
creases the amount of SrCrO4 phase which reacts to the
liquid phase and secondary solid phase through a peri-
tectic transformation. This secondary solid phase hinders
sintering. A higher sintering temperature also results in
pronounced grain growth in which originally sub-micro-
metre grains grow to almost ~2.7 μm in size at 1500 °C.
With this pronounced growth the grains become less
similar to an ideal sphere, which is manifested as a slight
decrease in the shape factor. Similar behaviour was ob-
served for the sintering of a combustion-derived LSCM
ceramic.23
In the LSCM phase, with a pure sample (La9Cr6)
with Cr content y = 0.6 and a low Sr-addition, the forma-
tion of well-connected grains is observed after sintering
at 1250 °C. With increasing sintering temperature, the
densification process normally advances and the grains
grow; however, the grain growth is somehow less pro-
nounced than that in other samples. The average grain
size for sample La9Cr6 sintered at various temperatures
1250 °C, 1300 °C, 1400 °C and 1500 °C is 0.22 μm, 0.36
μm, 0.63 μm and 1.6 μm, respectively. Furthermore, for
the sample La9Cr6, a calculation of the average
FERETMAX of pores versus the grain diameter gives the
highest value among all the samples. Since the average
pore diameter is comparable at sintering temperature
1250 °C in all samples, this value somehow indicates a
low average LSCM grain size and pore appearance in the
sample, which contribute mainly to the open porosity.
This fact together with the absolute value of porosity is
very important from the practical point of view if such
material is to be used as an anode layer in the operating
SOFC. In order to form a continuous phase of pores in
sintered samples, the porosity should be at least 30
vol.%, while the pore appearance should contribute to
open porosity to keep the LSCM anode layer permeable
for gases.
Several very important findings arise from the quanti-
tative microstructure analysis of LSCM samples. Re-
garding sintering optimisation, a lower Sr content (x =
0.1) with a somewhat higher chromium content (y = 0.6)
(sample La9Cr6) leads to proper microstructure forma-
tion at 1250 °C, where the grain-to-grain contact area is
enlarged, making it progressively easier to find a solid
continuous path of LSCM throughout the sample. At the
same time, the appropriate porosity is preserved and the
average grain size is the smallest, thus enlarging the in-
terface area where gaseous reactants meet the electro-
catalytic solid surface in a potential fuel cell. With addi-
tional information from the literature10 regarding LSCM
catalytic activity toward H2 oxidation, it can be con-
cluded that La0,9Sr0,1Cr0,6Mn0,4O3 is the most appropriate
LSCM chemical composition, which will also ensure the
desired microstructure characteristics at the relatively
low sintering temperature of 1250 °C.
4 CONCLUSIONS
La1–xSrxMn1–yCryO3± perovskite materials (x from 0
to 0.3 and y from 0.4 to 0.6) were prepared using the "re-
verse strike" carbonate co-precipitation method, which
has been shown to be an appropriate method since it al-
lows good control over the reaction system and the prep-
aration of LSCM materials with various compositions.
After calcination of the precipitated mixed carbon-
ate-hydroxide precursors at 1000 °C, the main crystalline
phase in all the samples is LSCM perovskite. In the sam-
ple with the highest Sr content (x = 0.3) and equal
amounts of chromium and manganese (y = 0.5), a stron-
tium secondary phase SrCrO4 was detected, while in
samples with a lower Cr content (y = 0.4) a lanthanum-
rich secondary phase La2CrO6 is formed. After sintering
at 1250 °C, the secondary phase re-dissolved into the
perovskite.
K. ZUPAN et al.: PHASE AND MICROSTRUCTURE DEVELOPMENT OF LSCM PEROVSKITE MATERIALS ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 743–748 747
Microstructure parameters for the LSCM ceramics
greatly depend on the sample composition. In samples
with equal contents of Cr and Mn, a slight addition of Sr
(x = 0.1) accelerates the sintering, while adding Sr to the
perovskite above a concentration of x = 0.2 supresses the
densifying process. In samples with a lower Cr content
(x = 0.4), the presence of a lanthanum-rich secondary
phase La2CrO6 causes non-homogenous microstructures
to form.
Samples with a lower Sr content (x = 0.1) and a
somewhat higher chromium content (y = 0.6) (La9Cr6)
lead to appropriate microstructure formation at sintering
temperatures as low as 1250 °C. Such a composition and
sintering temperature are also recognized as the most ap-
propriate parameters for suitable LSCM material.
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combustion method for solid oxide fuel application, Journal of Al-
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K. ZUPAN et al.: PHASE AND MICROSTRUCTURE DEVELOPMENT OF LSCM PEROVSKITE MATERIALS ...
748 Materiali in tehnologije / Materials and technology 50 (2016) 5, 743–748
V. CERNY, R. DROCHYTKA: ARTIFICIAL AGGREGATE FROM SINTERED COAL ASH
749–753
ARTIFICIAL AGGREGATE FROM SINTERED COAL ASH
UMETNI AGREGAT IZ SINTRANEGA PEPELA PREMOGA
Vit Cerny, Rostislav Drochytka
Brno University of Technology, Faculty of Civil Engineering, Veveri 95, 602 00 Brno, Czech Republic
cerny.v@fce.vutbr.cz
Prejem rokopisa – received: 2015-07-22; sprejem za objavo – accepted for publication: 2015-09-23
doi:10.17222/mit.2015.235
Fly ash is one of the most commonly used secondary raw materials in the Czech Republic. It is used predominantly for
re-cultivation, roads and additions to cement or plasters. The use of fly ash in the technology of sintering-based artificial
aggregate was tested in the 1980s. However, the production was stopped for various technological and economic reasons.
Nowadays, possibilities for the production of artificial aggregate are tested with fly ash produced in the Czech Republic and
could be promising for future technologies, mainly with respect to the stability of production. The paper presents part of the
study describing the possibilities of using microsilica and Fe2O3 for the optimization of a raw-material mix and fly-ash body.
Three samples of high-temperature lignite combustion fly ash were selected from prospective sources in the Czech Republic.
The fly ash was mixed with 5 % or 10 % additions. Then, samples were fired at temperatures of 1150 °C and 1200 °C. After
firing, the physico-mechanical properties of the fly-ash bodies and microstructure were evaluated. The results imply that the
addition of microsilica unambiguously improves the quality of the fly-ash body. The addition of Fe2O3 did not take part in the
formation of the melt and weakened the fly-ash body’s structure.
Keywords: fly ash, ash body, firing, microsilica, sintering, iron trioxide
Lete~i pepel sodi k najpogosteje uporabljanim sekundarnim surovinam na ^e{kem. Predvsem se uporablja za rekultivacijo, za
ceste in za cementne omete. Uporaba lete~ega pepela v tehnologiji izdelave sintranih agregatov je bila preizku{ena
v osemdesetih letih prej{njega stoletja. Vendar je bila proizvodnja ustavljena iz razli~nih tehnolo{kih in ekonomskih razlogov.
Dandanes se preizku{ajo mo`nosti za izdelavo umetnih agregatov iz lete~ega pepela nastalega na ^e{kem, kar je lahko
perspektivno za bodo~e tehnologije, predvsem iz stali{~a stabilnosti izdelave. ^lanek predstavlja del {tudije o mo`nosti uporabe
mikrosilike in Fe2O3 za optimizacijo me{anice surovin na osnovi lete~ega pepela. Trije vzorci lete~ega pepela pri
visokotemperaturnem zgorevanju lignita, so bili izbrani iz obetavnih virov, ki so na voljo na ^e{kem. Lete~i pepel je bil zme{an
s 5 % ali 10 % dodatka. Nato so bili vzorci `gani pri temperaturi 1150 °C in 1200 °C; po `ganju so bile dolo~ene
fizikalno-mehanske lastnosti teles iz lete~ega pepela in ocenjena je bila mikrostruktura. Rezultati ka`ejo, da dodatek mikrosilike
nedvomno izbolj{a kvaliteto kosov iz lete~ega pepela. Dodatek Fe2O3 ni sodeloval pri nastanku taline in je oslabil zgradbo kosov
lete~ega pepela.
Klju~ne besede: lete~i pepel, kos lete~ega pepela, `ganje, mikrosilika, sintranje, `elezov trioksid
1 INTRODUCTION
Power plants producing electrical energy create ener-
getic by-products during the combustion of pulverized
lignite. The dominant proportion of the by-products is
fly ash. Ecological and economic reasons motivate tech-
nological innovations focusing on using solid waste. The
production of artificial fly-ash aggregate is a suitable
construction material, which could be made from fly ash.
Artificial, sintered, fly-ash aggregate is one of the few
construction materials that can be produced only with fly
ash. European and worldwide trends of newly developed
building technologies accentuate the demand for the
production of high-quality, sintered, artificial, fly-ash
aggregate. If the character of the fly ash is optimal, no
further treatment is needed. However, not every sample
of fly ash has an optimal composition. The quality of fly
ash has an influence on the composition of the mix, the
technological parameters and the quality of the produced
aggregate. For enhancing the properties of sintered
fly-ash aggregate, additions commonly used in the cera-
mics industry could be used. The main candidates are
microsilica and oxides of iron.
Microsilica is an amorphous SiO2, which is common-
ly obtained from separators during the production of
ferrosilicon and silicium in electric arc furnaces. The
application of microsilica in fire-resistant materials has
been known for more than 40 years. The main task of
microsilica in fire-resistant ceramic materials is the reac-
tion in the system of binders, including the reaction me-
chanism at various temperatures. Various temperatures
can be critical as regards the reactions. Microsilica con-
sists of spheroidal particles with a mean diameter of
around 0.15 microns. These spheroidal particles are a
construction unit of primary agglomerates, which are
bound to one another by strong bonds. The large specific
surface and the wide distribution of microsilica increases
the effectiveness of the encapsulation of the grains and
the functionality of the fire-resistant ceramic materials
compared to a narrow fraction. Microsilica is usually the
finest part of the system with a specific surface of around
20 m2 g. The surface properties and possible impurities
are important for a determination of the properties of the
Materiali in tehnologije / Materials and technology 50 (2016) 5, 749–753 749
UDK 67.017:622.411.52:621.762.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)749(2016)
final product. Microsilica can add more than 50 % to the
total surface area of the particles in the mix.1
Oxides of iron influence the formation of a ceramic
body in various atmospheres2, temperatures and firing
cycles in furnaces There are three main forms of iron
used in ceramics: red iron trioxide (Fe2O3), black iron
monoxide oxide or magnetite (FeO or Fe3O4) and yellow
hydrated iron oxide (FeO(OH)). Iron trioxide is con-
siderably influenced by a reducing atmosphere, in which
it can act as a melting agent at high temperatures. The
presence of iron reduces the temperature of melting.3
From the point of view of a reduction of the firing tem-
perature, oxides of iron can be used as effective melting
agents for the production of sintered fly-ash aggregate.
The optimal mix of appropriate raw materials for the
production of artificial aggregate is mixed with water
and granulated on a cylindrical or plate granulator, which
makes the appropriate shape of the sintered fly-ash
aggregate. After granulation, the bodies are fired at a
maximum temperature of around 1200 °C. After suffi-
cient firing and cooling, the mix is crushed and screened
for the final fractions.
For the production of aggregate of good quality, the
granulometry of the input materials, their structure and
chemical composition are also important. The produced
sintered fly-ash aggregate is used for filtration layers,
back filling and lightweight concrete.4–6
2 EXPERIMENTAL PART
Three samples representing prospective sources in
the Czech Republic were selected for verification of the
appropriateness of high-temperature lignite combustion
fly ash for the production of artificial aggregate.
From the physico-mechanical and physico-chemical
parameters we selected the loss on ignition (CSN 72
0103) showing unburned residues, the bulk density (CSN
72 2071), the specific surface area and the rest on the
0.045-mm sieve (CSN 72 2072-6). We also conducted
the chemical and mineralogical analyses.
The parameters listed in Table 1 imply that fly ash
produced in the Czech Republic has a minimal unburnt
content and fulfills the requirements of the standard
(CSN 72 2072-6, 2013) for a maximum loss on ignition
of 15 %. For self-firing, it will be necessary to mix the
fly ash with pulverized coal to achieve an optimal value
of 8 % combustibles by weight. The greater coarseness
of the fly ash FA3 evaluated with respect to the rest on
the 0.0045-mm sieve and specific surface is evident. This
fly ash also has a larger content of iron and lime, which
can cause a reduction of the melting temperature of a
batch. All the values of the high-temperature fly-ash
samples fulfill the requirements of the Standard CSN 72
2072-6, 2013 for a minimal value of 800 kg m–3.
As is clear from Table 2 FA1 has the highest per-
centage of mullite and a lower content of amorphous
phase in comparison with FA2 and FA3. The minera-
logical analysis further confirms that FA2 contains more
minerals with a lower melting point and a low content of
mullite.
Table 2: Mineralogy of tested ashes, in mass fractions (w/%)
Tabela 2: Mineralogija preizku{enih pepelov, v masnih odstotkih
(w/%)
Sample
Quartz Mullite Hematite Magnetite Amorphousphase
SiO2 Al6Si2O13 Fe2O3 Fe3O4 –
FA1 7.0 39.3 1.2 0.1 39.5
FA2 7.2 19.1 2.5 3.1 55.2
FA3 7.8 32.3 – 0.2 58.1
As an addition for the experimental work, microsilica
(96 % SiO2) and Fe2O3 (powder, 97 % cleanliness) were
selected. As a reference, samples from pure fly ash were
used, which were then modified with a 5 % or 10 %
addition. The mixtures were mixed with water to reach
the limit of fluidity. Samples of size 20 mm × 20 mm ×
100 mm were made, which were next day dried at 60 °C
for 2 h and then fired in a muffle kiln. The firing was
characterized by an initial temperature of 25 °C and the
rate of firing the muffle kiln of 300 °C/h and an isother-
mal dwell at 1150 °C (resp. 1200 °C) for 10 min. After
firing and natural cooling, the specimens were taken out
of the kiln and placed in a desiccator. After thermal sta-
bilization, their density (CSN EN 1015-10), compressive
strength (CSN EN 14617-15) and water-absorbing
capacity (CSN EN 1097-6) were determined. Then, the
samples were analyzed with a scanning electron micro-
scope using a sensing element in an environmental form.
Primarily, the structure was analyzed and the influence
of the addition on the fly ash body.
3 RESULTS AND DISCUSSION
After firing in a laboratory furnace and sufficient
cooling to laboratory temperature 20±5 °C, the following
V. CERNY, R. DROCHYTKA: ARTIFICIAL AGGREGATE FROM SINTERED COAL ASH
750 Materiali in tehnologije / Materials and technology 50 (2016) 5, 749–753
Table 1: Main parameters of the tested ashes
Tabela 1: Glavni parametri preizku{enih pepelov
Sample
Loss on
ignition
(%)
Bulk density
(compacted)
(kg/m3)
Rest on the sieve
0.045 mm
(%)
Specific
surface
(m2/kg)
Chemical composition (%)
SiO2 Al2O3 Fe2O3 SO3 CaO
FA1 1.19 990 58.5 329 47.7 28.2 5.6 0.13 1.1
FA2 1.15 1010 72.0 234 50.0 23.4 14.5 0.26 3.4
FA3 1.07 1110 53.1 299 54.6 29.5 5.5 0.10 1.8
physico-mechanical parameters of the test specimens
were determined.
Figure 1 shows the results of the determination of
density. In general, it can be stated that the density grows
slightly with an increasing temperature of firing. The
samples based on FA1 showed the highest values of den-
sity, while the samples based on FA3 showed the lowest
density.
A determination of the compressive strength (Figure
2) showed that the strengths at the firing temperature
1200 °C were considerably higher than the strengths at
1150 °C. The addition of silica had a very positive effect
on the strength of the fly ash body – samples with as
little as 5 % showed considerably higher strengths. In
contrast, the addition of Fe2O3 weakened the fly-ash
body and the measured strengths were often lower than
those of the reference samples based on pure fly ash.
An evaluation of the results of the determination of
the water-absorbing capacity (Figure 3) in some cases
shows the influence of firing temperature on the quality
of the fly-ash body. The water-absorbing capacity of the
samples fired at 1150 °C was often much higher. This
can be caused by insufficient sintering of the fly-ash
body and a higher proportion of open porosity. The
influence of the type of addition on the results of the
water-absorbing capacity is most marked for the samples
based on FA with the addition of Fe2O3.
To clarify the results determined during the evalua-
tion of the physico-mechanical parameters, the samples
V. CERNY, R. DROCHYTKA: ARTIFICIAL AGGREGATE FROM SINTERED COAL ASH
Materiali in tehnologije / Materials and technology 50 (2016) 5, 749–753 751
Figure 3: Water absorption of fired samples
Slika 3: Absorpcija vode `ganih vzorcev
Figure 1: Density of fired samples
Slika 1: Gostota `ganih vzorcev
Figure 5: Detail of structure of fly-ash body with 5 % of mass
fractions of Fe2O3
Slika 5: Detajl strukture kosa lete~ega pepela s 5 % masnega dele`a
Fe2O3
Figure 2: Compressive strength of fired samples
Slika 2: Tla~na trdnost `ganih vzorcev
Figure 4: Structure of fly-ash body with 5 % of mass fractions of
Fe2O3
Slika 4: Struktura kosa lete~ega pepela s 5 % masnega dele`a Fe2O3
were examined under a scanning electron microscope
with a sensing element in an environmental form. The
aim was an evaluation of the influence of firing tempe-
rature and additions on the structure of the fly-ash body.
Figures 4 to 7 show the structure of a surface of a
test specimen based on FA3 with a 5 % addition of Fe2O3
and fired at 1200 °C. Figure 4 shows a low proportion of
sintered structure and visible grains of fly ash. Figure 5
shows a part of a sample with melted material at higher
magnification, where non-reacted grains of Fe2O3 are
visible. Figure 6 shows a closer detail of the same place
grains of Fe2O3 and Figure 7 an analysis of the elements
in a larger area where the dominant Fe is evident in the
grains. This fact proves that Fe2O3 did not take part in the
formation of a solid structure, and that it weakened the
fly-ash body.
Figures 8 and 9 show the structure of the test spe-
cimens based on FA3 with a 10 % addition of silica. Fig-
ure 8 shows a sample fired at 1150 °C, where the grains
of fly ash and a minimal proportion of melted material
are clear. Figure 9 shows a sample fired at 1200 °C,
where the proportion of melted material is more con-
V. CERNY, R. DROCHYTKA: ARTIFICIAL AGGREGATE FROM SINTERED COAL ASH
752 Materiali in tehnologije / Materials and technology 50 (2016) 5, 749–753
Figure 8: Structure of fly-ash body with 10 % of mass fractions of
silica
Slika 8: Struktura kosa lete~ega pepela z 10 % masnega dele`a kre-
mena
Figure 6: Non-reacted proportion of Fe2O3 in fly-ash body
Slika 6: Nereagiran dele` Fe2O3 v kosu lete~ega pepela
Figure 9: Detail of structure of fly-ash body with 5 % of mass
fractions of silica
Slika 9: Detajl strukture kosa lete~ega pepela s 5 % masnega dele`a
kremena
Figure 7: Marked composition of elements of sample with non-
reacted Fe2O3
Slika 7: Razporeditev elementov v vzorcu, ki ni reagiral z Fe2O3
siderable and the grains have a stronger and more homo-
geneous structure.
4 CONCLUSIONS
The manufacture of artificial aggregate from fly ash
is one of the options for using the maximum proportion
of this raw material in construction materials. The cha-
racter of the fly ash and good control of the technology
for producing the aggregate by self-firing also brings
savings for traditional raw materials. The paper des-
cribed the experimental testing of possibilities for the
modification of fly ash with the addition of microsilica
or Fe2O3 in order to achieve better quality of the fly-ash
body with a lower energy demand of the process. The
results show how important it is to set the firing tem-
perature at 1200 °C, which makes sure it is possible to
achieve considerably higher strengths in the system. An
evaluation of the influence of the used additions showed
that the addition of microsilica unambiguously improves
the quality of the fly-ash body. Higher strengths and a
lower water-absorbing capacity were achieved compared
to use of pure fly ash. The addition of Fe2O3 did not take
part in the formation of the melt, and stayed almost un-
changed and weakened the fly-ash body structure.
Acknowledgements
This paper has been worked out under the project No.
LO1408 "AdMaS UP – Advanced Materials, Structures
and Technologies", supported by Ministry of Education,
Youth and Sports under the "National Sustainability
Programme I".
5 REFERENCES
1 B. Sandberg, B. Myhre, Microsilica A Versatile Refractory Raw Ma-
terial, Indian Refractories Congress, Jamshedpur, India 1994, 1–7
2 R. Magrla, M. Fridrichova, K. Kulisek, K. Dvorak, O. Hoffmann,
Utilisation of Fluidised Fly Ash for Reduction of CO2 Emissions at
Portland Cement Production, Advanced Materials Research, 1054
(2014), 168–171, doi:10.4028/www.scientific.net/AMR.1054.168
3 J. Britt, All about iron: Iron is everywhere in many different forms,
but that doesn’t mean it has to be boring-or even brown, Cmtech-
nofileiron, 2011
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taining Lightweight Expanded Clay Aggregate on Test Results Ob-
tained by Means of Impact Hammer, Advanced Materials Research,
753–755 (2013), 663–667, doi:10.4028/www.scientific.net/AMR.
753-755.663
5 B. Yun, I. Ratiyah, P. A. M. Basheer, Properties of lightweight con-
crete manufactured with fly ash, furnace bottom ash, and Lytag,
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Technology, Beijing 2004, 77–88
6 T. Melichar, J. Bydzovský, Study of the parameters of lightweight
polymer-cement repair mortars exposed to high temperatures,
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doi:10.4028/www.scientific.net/AMM.395-396.429
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public 2006
V. CERNY, R. DROCHYTKA: ARTIFICIAL AGGREGATE FROM SINTERED COAL ASH
Materiali in tehnologije / Materials and technology 50 (2016) 5, 749–753 753

M. STASZUK et al.: INVESTIGATION STUDIES INVOLVING WEAR-RESISTANT ALD/PVD ...
755–759
INVESTIGATION STUDIES INVOLVING WEAR-RESISTANT
ALD/PVD HYBRID COATINGS ON SINTERED TOOL
SUBSTRATES
PREISKAVE OBRABNE ODPORNOSTI HIBRIDNEGA NANOSA
ALD/PVD NA SINTRANEM ORODJU
Marcin Staszuk, Daniel Paku³a, Tomasz Tañski
Silesian University of Technology, Institute of Engineering Materials and Biomaterials, Konarskiego Street 18A, 44-100 Gliwice, Poland
marcin.staszuk@polsl.pl
Prejem rokopisa – received: 2015-07-22; sprejem za objavo – accepted for publication: 2015-10-13
doi:10.17222/mit.2015.236
This paper is an important research contribution to the development of PVD coatings, in particular, on ceramic substrates,
which, due to their dielectric properties, are difficult materials to coat using this technique. In order to satisfy the desired
expectations relating to the PVD coatings, one of the basic properties must be provided for, the adherence to the substrate. The
main aim of this research is to investigate the structure and mechanical properties of the coatings deposited in a hybrid process,
comprising the atomic-layer deposition (ALD) and cathodic-arc evaporation (CAE-PVD) on sintered carbides and multipoint
ceramic cutting tools. The concept of this research study involves an execution and investigation of ALD + PVD hybrid coatings
on sintered carbides and a sialon-ceramic substrate, and defining the influence of the ALD interlayer on the adherence of the
investigated coatings. The critical load Lc, which is the adhesion measure of coats, was determined with the scratch-test method
and a tribological test made with a pin-on-disk tester. Observations of the surface topography and wear mechanism were
performed using a scanning electron microscope and atomic-force microscopy. The investigation studies showed that an ALD
layer considerably improves the adherence of the PVD layer to the tool-ceramic substrate. The research of the coatings on
sintered-carbide substrates showed that the adherence of the PVD coating to the substrate deteriorates in the case of applying an
ALD interlayer.
Keywords: PVD, ALD, hybrid coatings, tool ceramics, sintered carbides
^lanek je pomemben raziskovalni prispevek k razvoju PVD nanosov, {e posebno na kerami~ni podlagi, ki je zaradi dielektri~nih
lastnosti zahteven material za nana{anje s to tehniko. Da bi zadovoljili pri~akovanja na delu s PVD nanosom, je oprijemljivost
na podlago osnovna lastnost, na katero je potrebno biti pozoren. Glavni namen raziskave je preiskati strukturo in mehanske
lastnosti nanosa, nane{enega s hibridnim postopkom, ki obsega nanos atomskih plasti (ALD) in katodno izparevanje v obloku
(CAE-PVD) na sintrane karbide in ve~to~kovno kerami~no rezilno orodje. Koncept te raziskave vklju~uje izvedbo in preiskavo
hibridnega nanosa ALD + PVD na podlago iz sintranih karbidov in sialon keramike ter dolo~itev vpliva vmesne plasti ALD na
oprijemljivost preiskovanih nanosov. Kriti~na obremenitev Lc, ki je merilo oprijemljivosti nanosov, je bilo dolo~eno s
preizkusom razenja in s tribolo{kim preizkusom trn na plo{~i. S pomo~jo vrsti~nega elektronskega mikroskopa in mikroskopa
na atomsko silo je bilo izvedeno opazovanje topografije povr{ine in mehanizma obrabe. Izvedene preiskave ka`ejo, da ALD
nanos mo~no izbolj{a oprijemljivost PVD nanosa na orodno keramiko. Pri nanosu na podlago iz sintranega karbida je raziskava
pokazala, da je pri uporabi ALD vmesne plasti, oprijemljivost PVD nanosa slab{a.
Klju~ne besede: PVD, ALD, hibridni nanosi, keramika za orodja, sintrani karbidi
1 INTRODUCTION
Sialon-tool ceramics are a group of tool materials
combining the mechanical properties of silicon nitride
Si3N4 and the chemical properties of Al2O3. Ceramic
sinters are fabricated with powder-metallurgy methods,
but in contrast to sintered carbides, they do not contain a
metallic binder. The use of tool ceramics as compared to
sintered carbides is still considerably small but it is
growing, also due to the application of thin wear-resis-
tant layers PVD and CVD.1,2
Surface engineering technologies play a significant
role in material engineering and are viewed as the fun-
damental scope of knowledge in this field of engi-
neering. The research entities working in this field try to
investigate and identify the phenomena taking place on
the surfaces of the machined materials. In order to
improve the operating properties of products, various
surface-machining technologies are applied.2–6 The
advantages arising from the application of PVD coatings
on cutting tools such as a high microhardness and resis-
tance to abrasion, a low friction coefficient of the tool
covered with the coat, the resistance to oxidation or to
the formation of build-up edges on the tool, combined
with the possibility to machine materials without liquid
cooling lubricants are important contributions advocating
the development of this technological field. However, in
order to improve the desired properties of PVD coatings,
one of the basic properties must be taken care of, which
is the adherence to the substrate. The coating of ceramic
materials in PVD processes, including also tool cera-
mics, is difficult due to dielectric properties, since the
inability of the substrate polarization during the coating
Materiali in tehnologije / Materials and technology 50 (2016) 5, 755–759 755
UDK 620.19:620.178.16:621.793.8 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)755(2016)
process makes it difficult to obtain coatings with a good
adherence to ceramic substrates.2,7–11
The fabrication technology of stratified coatings of
the type "adhesive layer/PVD layer" using the multi-
stage surface machining is one of the most modern
methods to modify the properties of the surface layer and
it consists, in a successive application, of two (or more)
technologies of surface engineering. In terms of material
effects, we obtain multilayer systems whereof one part is
made up of an appropriately selected adhesive layer
fabricated on the surface of the substrate, which protects
the "proper" PVD layer against a loss of internal cohe-
sion and an insufficient adhesion to the substrate. And on
the other end of the system, we obtain a PVD coating
which efficiently insulates the substrate, limiting the
impact of the external harmful factors encountered in the
operating process.2,12
The coatings of the (Ti,Al)N type are isomorphic,
with titanium nitride which is still widely applied. The
presence of aluminum in the coatings of this type brings
about the situation where the service-life temperature of
these coatings exceeds 970 K, and in such operating con-
ditions of raised temperature, a layer of Al2O3 is formed
on the surface, which forms a diffusive barrier for the
atmospheric oxygen.13,14
Zinc oxide is a semiconductor characterized by a
high energy gap (3.4 eV), high exciton binding energy
(60 meV) and easy n-type doping.15,16 Zinc oxide, ZnO,
obtained with the use of the ALD technique on a sialon-
ceramic substrate enables the polarization of this sub-
strate during the technological process of PVD and,
hence, the objective of this paper is to investigate the im-
pact of the ZnO layer obtained with the use of the ALD
technique on the adhesion of the hybrid coating
ALD/PVD of the ZnO/(Ti,Al)N type to the sialon sub-
strate. For the sake of comparison, we also investigated
the coatings on sintered carbide substrates.
2 MATERIALS AND METHOD
The research studies were carried out on multipoint
cutting tools made of sintered carbides WC-Co and
sialon-tool ceramics covered in the PVD and ALD/PVD
processes. The tools were covered during the cathodic-
arc evaporation process CAE-PVD with the (Ti,Al)N
coating type and in the ALD/PVD hybrid process using
the ZnO/(Ti,Al)N coating type.
The topography of the surfaces and wear mechanisms
of the coatings was viewed with a Supra 35 scanning
electron microscope from Zeiss. The secondary-electron
(SE) detection was used to obtain the images of the
tested samples, with an acceleration voltage of 5–20 kV.
The topography of the coatings was tested using a
Park Systems XE-100 atomic-force microscope. The
tests were carried out in the non-contact mode.
The adherence of the coatings to the substrate was
assessed with a scratch test on a Revetest device from
CSEM. The method consists of moving a Rockwell C
diamond indenter through the surface of a tested sample
at a constant speed, with the applied force growing
linearly. The ranges of the applied load were 0–100 N
and 0–200 N. The Lc critical load, at which the coating
adhesion is lost, was determined according to the acou-
stic-emission value registered during the measurement
and by observing the scratches formed during the scratch
test.
The abrasive-wear-resistance tests and the wear-fac-
tor tests for the tested coatings were performed with the
pin-on-plate method using a CSEM Tribometer (THT).
A 6 mm Al2O3 ball was used as the counter specimen.
The tests were made at room temperature. The following
test conditions were applied: a normal force of FN = 10 N
and a movement speed of v = 0.1 m/s. The total distance
of 1 km was set for all the tested specimens.
3 DISCUSSION AND TEST RESULTS
The value of critical load Lc, being the adherence
measure for the investigated coatings on sintered-carbide
substrates or sialon-tool ceramic substrates, was deter-
mined using a scratch test (Figure 1). The research
studies show that the ALD/PVD hybrid coat has a much
M. STASZUK et al.: INVESTIGATION STUDIES INVOLVING WEAR-RESISTANT ALD/PVD ...
756 Materiali in tehnologije / Materials and technology 50 (2016) 5, 755–759
Figure 1: a) Acoustic emission (AE) and friction force Ft depending
on load Fn for the ZnO/(Ti,Al)N coating on sialon ceramics, b) scratch
failure at Lc (opt) = 65 N, magnified 200×
Slika 1: a) Akusti~na emisija (AE) in sila trenja Ft v odvisnosti od
obremenitve Fn pri ZnO/(Ti,Al)N nanosu na sialon keramiko, b) po-
{kodba z razo pri Lc (opt) = 65 N, pove~ano 200×
higher adherence to the sialon-ceramic substrate than the
PVD coating without the adhesive layer (Table 1). A re-
verse situation takes place in the case of coatings on
sintered-carbide substrates. Due to high adhesion of the
PVD coating to the sintered-carbide substrate, the ALD
layer causes a drop in the critical load from 109 N to
55 N.
Table 1: Critical loads Lc of the investigated coatings
Tabela 1: Kriti~ne obremenitve Lc preiskovanih nanosov
Type of coating
Critical load Lc, N
Sialon substrate Cemented-carbidesubstrate
(Ti,Al)N 21 109
ZrO/(Ti,Al)N 64 55
The damage done to the coats during the performed
adhesion/scratch tests was identified on the basis of
observations with the scanning electron microscope. The
studies carried out show that it is the delamination,
which is the principle mechanism of the coat damage on
the sialon-ceramic substrate after the critical load Lc has
been exceeded. In the case of the PVD coat, which is
characterised by a low adherence to the ceramic substra-
te, a total delamination with vast chipping was identified.
Furthermore, a low adherence of the PVD coat to the
ceramic substrate indicates a considerable morphological
non-homogeneity and, in particular, the presence of
chippings found on the whole surface of the coat (except
for the scratch area) (Figure 2). A low adherence of the
coat to the substrate results in its unprompted chipping
due to the internal strain of the coat. In the case of the
ZnO/(Ti,Al)N hybrid coat on the sialon substrate, dela-
mination was identified on both sides of the sample
(Figure 3). The chipping on both sides of the scratched
sample is definitely less vast than that observed in the
case of the (Ti,Al)N coat on the same substrate. The only
morphological defects of the hybrid coat are the drop-
shaped microparticles, inseparably connected with the
cathodic-arc evaporation process (Figure 4). In the case
M. STASZUK et al.: INVESTIGATION STUDIES INVOLVING WEAR-RESISTANT ALD/PVD ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 755–759 757
Figure 4: Surface topography of the ZnO/(Ti,Al)N coating deposited
onto the sialon-ceramic substrate: a) SEM, b) AFM
Slika 4: Topografija povr{ine ZnO/(Ti,Al)N nanosa, nane{enega na
podlago iz sialonske keramike: a) SEM, b) AFM
Figure 2: Characteristic failure obtained with the scratch test of the
(TiAl)N coating deposited on the sialon-ceramic substrate
Slika 2: Zna~ilna po{kodba pri preizkusu z razenjem nanosa (Ti,Al)N,
nane{enega na podlagi iz sialonske keramike
Figure 3: Characteristic failure obtained with the scratch test of the
ZnO/(TiAl)N coating deposited on the sialon-ceramic substrate
Slika 3: Zna~ilna po{kodba pri preizkusu razenja ZnO/(Ti,Al)N
nanosa, nane{enega na podlago iz sialonske keramike
of the coats on sintered-carbide substrates, friction is the
dominant mechanism. In the case of the PVD coat on the
sintered-carbide substrate, which is characterised by the
highest adherence among the investigated samples, no
abrasion deep down in the substrate was identified until
the maximum load was reached during the scratch test.
However, numerous cohesive cracks were identified after
surpassing the critical load Lc. In the ZnO/(Ti,Al)N coat
on the same substrate, we identified a uniform abrasion
spreading down into the substrate after the critical load
was exceeded.
With the tests on the abrasive-wear resistance of the
coats deposited on sialon ceramics and on sintered car-
bides using the pin-on-plate method, we found that in
almost all the cases where the coats were applied the
coats were damaged down to the substrate zone (Figures
5 and 6). The dominant wear mechanism of the investi-
gated coats was attrition. We also found that in some
cases the damaged coat sticks onto the material of the
counter specimen, which has a direct impact on the
changing values of the friction coefficient. The research
studies also confirm a positive impact of the ZrO layer
on the attrition resistance as compared to the specimen
without such a layer. At one end of the scratch, we ob-
served a lot of damage to the coat deposited on the sialon
substrate without the ZrO layer. The damage is of
adhesive character – we can observe stratifications and
cracks of the layer (Figure 5). And in the case of the
coat deposited with the ZrO layer, the end of the scratch
is uniformly grated without any adhesive damage (Fig-
ure 6).
4 CONCLUSION
Nowadays, the improvement of the wear resistance of
cutting tools mainly involves surface machining, while
the selection of the technology or coating material aims
at ensuring the proper resistance of a tool to dominant
wear mechanisms. The operating properties of the coats
resistant to the wear are the results of many components,
primarily the microhardness, the grain size and the
adherence to the substrate. Especially the last property is
of key importance, and in view of the performed research
studies2, the grain size, the thickness and the micro-
hardness of the obtained coats have smaller impacts than
the adhesion on the durability of cutting tools since the
changes in these properties have smaller impacts on
service life.
The paper presents the results of research studies in-
volving the application of an ALD layer to obtain a
better adhesion of the PVD coat to the tool sialon-cera-
mic substrate or sintered-carbide substrate. We inve-
stigated the (Ti,Al)N coat deposited on both substrates,
with the ZnO layer and without it. On the basis of the
investigation studies, we can state that zinc oxide con-
siderably improves the adherence of the PVD coat to the
ceramic substrate, which was confirmed with scratch
tests and observations of the damage made during the
tests, using a scanning electron microscope. The critical
load, being the adherence measure of the PVD coat on
the ceramic substrate, increased due to the application of
the adhesive layer by over 200 %. With respect to the
investigated coats on sintered carbides, the adherence of
the PVD coat to the ALD layer decreased as compared to
the coat without ALD.
The improvement in the adherence of the PVD coat
to the ceramic substrate is undaubtly connected with the
possibility to polarize the substrate during the coating
process involving the application of the ZnO layer. The
results of the research studies are important since the
deposition of ceramics in PVD processes is hindered due
to the dielectric properties of ceramics.
M. STASZUK et al.: INVESTIGATION STUDIES INVOLVING WEAR-RESISTANT ALD/PVD ...
758 Materiali in tehnologije / Materials and technology 50 (2016) 5, 755–759
Figure 6: Trace of tribological damage on the surface of the
ZnO/(Ti,Al)N coat deposited on the sialon-ceramic substrate
Slika 6: Sled tribolo{ke po{kodbe na povr{ini ZnO/(Ti,Al)N nanosa,
nane{enega na podlago iz sialonske keramike
Figure 5: Trace of tribological damage on the surface of the (Ti,Al)N
coat deposited on the sialon-ceramic substrate
Slika 5: Sled tribolo{ke po{kodbe na povr{ini (Ti,Al)N nanosa,
nane{enega na podlago iz sialonske keramike
Acknowledgements
The publication was co-financed by the statutory
grant of the Faculty of Mechanical Engineering of the
Silesian University of Technology in 2015.
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Materiali in tehnologije / Materials and technology 50 (2016) 5, 755–759 759

S. P. HOVEIDA MARASHI: DISSIMILAR SPOT WELDING OF DQSK/DP600 STEELS: THE WELD-NUGGET GROWTH
761–765
DISSIMILAR SPOT WELDING OF DQSK/DP600 STEELS:
THE WELD-NUGGET GROWTH
TO^KASTO VARJENJE JEKEL DQSK/DP600: RAST JEDRA ZVARA
Seyed Pirooz Hoveida Marashi
Amirkabir University of Technology, Mining and Metallurgical Engineering Department, Tehran, Iran
pmarashi@aut.ac.ir
Prejem rokopisa – received: 2015-07-30; sprejem za objavo – accepted for publication: 2015-10-12
doi:10.17222/mit.2015.241
The weld-nugget size is the key issue in determining the mechanical properties of resistance spot welds. This paper aims at
investigating the weld-nugget growth of dissimilar-resistance spot welding of ferrite-martensite DP600 and drawing-quality
special-killed (DQSK) low-carbon steel. It was revealed how the weld-nugget size is influenced by the main welding para-
meters: welding current, welding time and electrode force. The weldability lobe was established and proper welding conditions
for the welds with a sufficient size and without an expulsion were determined. Using the experimental data, an empirical
relationship between the weld-nugget size and the welding parameters was developed.
Keywords: resistance spot welding, dual-phase steel, dissimilar welding, weld-nugget growth
Velikost jedra zvara je klju~nega pomena pri dolo~anju mehanskih lastnosti uporovnih to~kastih zvarov. Namen tega ~lanka je
preiskava rasti jedra zvara pri to~kastem uporovnem varjenju feritno-martenzitnega DP600 in pomirjenega maloglji~nega jekla
DQSK za globoki vlek. Ugotovljeno je bilo, kako na velikost jedra zvara vplivajo glavni parametri varjenja: varilni tok, ~as
varjenja in pritisk elektrode. Vzpostavljen je bil varilni kiln in dolo~eni so bili pravilni varilni pogoji za izdelavo dovolj velikih
zvarov, brez izgonov taline. Z uporabo eksperimentalnih podatkov je bila postavljena empiri~na odvisnost med velikostjo jedra
zvara in parametri varjenja.
Klju~ne besede: uporovno to~kasto varjenje, dvofazno jeklo, varjenje razli~nih materialov, rast jedra zvara
1 INTRODUCTION
Resistance spot welding is considered as the domi-
nant process for joining sheet metals in the automotive
industry. Simplicity, low cost, high speed (low process
time) and automation possibility are the advantages of
this process. The quality and mechanical behavior of
spot welds significantly affect the durability and crash-
worthiness of a vehicle.1–3 To ensure and maintain the
structural integrity of a finished component under a wide
range of operating conditions, e.g., a crash situation, a
remotest possibility of producing even one or two defec-
tive welds in a critical component needs to be eliminated.
These requirements, coupled with the uncertainties about
the weld quality due to the difficulty of applying non-
destructive tests to spot welds, are responsible for the
practice of making more spot welds than needed for
maintaining the structural integrity. Typically, there are
about 2000–5000 spot welds in a modern vehicle.
Around 20–30 % of these spot welds are due to the un-
certainty of the quality of spot welds. A significant cost
associated with over-welding provides a considerable
driving force for optimizing this process.4
Resistance spot welding is a process of joining two or
more metal parts using fusion at discrete spots at the
interface of workpieces. The resistance to the current
flow through the metal workpieces and their interface
generates heat; therefore, the temperature rises at the
interface of the workpieces. When the melting point of
the metal is reached, the metal will begin to fuse and a
nugget begins to form. The current is then switched off
and the nugget is cooled down to solidify under
pressure.5,6
It is well established that the geometrical attributes of
spot welds, particularly the weld-nugget size, are the
most important controlling factors determining the me-
chanical strength of RSWs.7–12 In this regard, the
weld-nugget size was included in several empirical rela-
tions. For example, J. Heuschkel13 developed empirical
relations among the tensile-shear strength (P), weld-
nugget diameter (D), base-metal tensile strength (BM),
sheet thickness (t) and base-metal chemical composition
(C, Mn):
[ ]P Dt C= +  ( . )0 05Mn BM (1)
where  and  are material-dependent coefficients.
Similar relations were developed by other researchers.
For example, the following relation was developed by J.
M. Sawhil and J. C. Baker14 for the tensile-shear
strength of spot welds:
P ftD=  BM (2)
where f is a material-dependent coefficient, with a value
between 2.5 and 3.1.
Considering the importance of the weld-nugget size
for the quality of spot welds, there is a need to study the
Materiali in tehnologije / Materials and technology 50 (2016) 5, 761–765 761
UDK 621.791.76/.79:691.714:621.791.05 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)761(2016)
effects of RSW parameters such as welding current,
electrode force and welding time on this key physical
weld attribute. Moreover, an increased use of advanced
high-strength steels (AHSS), particularly ferrite-marten-
site DP steels, led to a wider range of possible material
combinations in the resistance spot welding (RSW) of
body-in-white assemblies. Therefore, there is clearly a
practical need for the study of the weld-nugget growth
during the RSW of dissimilar steel grades. In this paper,
weld-nugget growth characteristics of dissimilar RSW of
ferrite-martensite DP600 and drawing-quality special-
killed (DQSK) low-carbon steel are investigated. The
aim of this paper is to reveal how the weld-nugget size is
influenced by the main welding parameters: welding
current, welding time and electrode force.
2 EXPERIMENTAL PROCEDURE
2-mm-thick drawing-quality special-killed (DQSK)
low-carbon steel and 2-mm-thick DP600 dual-phase
steel sheets were used as the base metals. The chemical
compositions of the base metals are shown in Table 1.
Resistance spot welding was performed using a PLC-
controlled, 120 kVA AC pedestal-type resistance-spot-
welding machine. The welding was conducted using a
45-deg truncated-cone RWMA Class 2 electrode with an
8-mm face diameter.
Table 1: Chemical compositions of steels used in this study
Tabela 1: Kemijska sestava jekel, uporabljenih v {tudiji
Base metal C Mn Si S P
DP600 0.035 1.08 0.388 0.004 0.038
LCS 0.065 0.204 0.095 0.017 0.018
To study the effects of the welding conditions
(welding current, welding time and electrode force) on
the weld failure mode, several welding schedules were
used. Figure 1 shows a schematic of the welding sche-
dules used in this study. A total of 60 combinations of
the welding current, the welding time and the electrode
force were performed.
Samples for the metallographic examination were
prepared using the standard metallographic procedure.
Light microscopy was used to examine the macrostruc-
tures and microstructures and to measure the weld fusion
zone (weld nugget). The samples for the metallographic
examination were prepared using the standard me-
tallographic procedure. A 4 % Nital etching reagent was
used to reveal the macrostructures of the samples. The
FZ size is defined as the width of the weld nugget at the
sheet/sheet interface in the longitudinal direction. The
indentation depth is expressed as the percentage of the
sheet thickness.
3 RESULTS AND DISCUSSION
3.1 Weld macrostructure
Figure 2 shows the macrostructure of a weld joining
DP600 and low-carbon steel indicating that there are
three distinct microstructural zones:
1) The weld nugget (WN) or fusion zone (FZ) which is
melted during the welding process and then resoli-
dified showing a cast structure. The macrostructure
of the weld nugget consists of columnar grains.
2) The heat-affected zone (HAZ) which does not melted
but undergoes microstructural changes.
3) The base metal (BM).
3.2 Effects of the welding parameters on the weld-
nugget growth
The effects of welding parameters on the weld-
nugget size are shown in Table 2. Contour plots for the
weld-nugget size versus the welding current and the
welding time at three levels of the electrode force are
shown in Figure 3. According to these results, the
following points can be drawn:
1) The welding current has a profound effect on the
weld-nugget growth. Increasing the welding current
increases the weld-nugget size.
2) Increasing the welding time increases the weld-
nugget size.
3) Increasing the electrode force decreases the weld-
nugget size. Indeed, when applying electrode force,
S. P. HOVEIDA MARASHI: DISSIMILAR SPOT WELDING OF DQSK/DP600 STEELS: THE WELD-NUGGET GROWTH
762 Materiali in tehnologije / Materials and technology 50 (2016) 5, 761–765
Figure 2: Macrostructure of a weld of DP600 and low-carbon steel:
FZ size is defined as the width of the weld nugget at the sheet/sheet
interface in the longitudinal direction
Slika 2: Makrostruktura zvara DP600 in malooglji~nega jekla: FZ je
{irina pretaljenega jedra na stiku plo~evin v vzdol`ni smeri
Figure 1: Schematic of the welding schedules used in this study
Slika 1: Shema ~asovnega poteka varjenja v tej {tudiji
there is need to use higher welding current and
welding time to obtain a specific weld-nugget size.
The amount of heat generated at the sheet-to-sheet
interface during the spot-welding process is the main
reason for the nugget formation and its strength. The
heat generated during the resistance spot welding can be
expressed as follows:
Q RI t= W W
2 (3)
where Q, R, Iw and tw are the generated heat, the elec-
trical resistance, the welding current and the welding
time, respectively.
Therefore, the three main parameters affecting the
weld-nugget growth are the welding current, the welding
time and the electrical resistance. The heat varies directly
with the interface resistance, the welding time and the
second power of the welding current. Again, this contact
(interface) resistance varies in a complex manner and it
is influenced by the electrode force, the surface condi-
tions of the sheets used and also by the geometry of the
electrode tip.
Increasing the welding current and the welding time
increases the heat generation, which in turn, causes an
enlargement of the weld nugget.
The static electrical resistance (i.e., the contact resis-
tance) is mainly governed by the electrode force, which
in turn controls the weld-nugget formation.15 On a duc-
tile material, where a normal force is applied across the
contact interface, the number of surface asperities
supporting the applied load gradually increases due to
their successive yielding. In other words, the true contact
area will initially be a relatively small fraction of the
macroscopic, or apparent, contact area. Later, the true
contact area will increase with the application of load
and, in the limit, approach the apparent contact area.8
Therefore, an increase in the electrode force decreases
the electric resistance and thus reduces the generated
heat at the sheet/sheet interface.
Since the generated heat is proportional to the
squared current, the current to the duration and the con-
tact resistance is inversely proportional to the electrode
force, another parameter, the so-called heat factor, can be
defined as follows:
Heat factor
W W
e
=
I t
F
2
(4)
S. P. HOVEIDA MARASHI: DISSIMILAR SPOT WELDING OF DQSK/DP600 STEELS: THE WELD-NUGGET GROWTH
Materiali in tehnologije / Materials and technology 50 (2016) 5, 761–765 763
Table 2: Effects of welding current and welding time on the weld-
nugget size at three different electrode forces (S: small, A: acceptable,
E: expulsion)
Tabela 2: Vpliv varilnega toka in ~asa varjenja na velikost pretalje-
nega jedra pri treh razli~nih pritiskih elektrod
Welding
current
(kA)
Welding
time (s)
Weld-nugget size (mm)
F=4.1 kN F=5.1 kN F=5.7 kN
8 0.3 4.35 (S) 2.3 (S) 1.2 (S)
8 0.4 4.67 (S) 3.15 (S) 2.2 (S)
8 0.5 4.9 (S) 3.8 (S) 2.6 (S)
8 0.6 5.12 (S) 4.24 (S) 3.7 (S)
9 0.3 4.8 (S) 4.4 (S) 4 (S)
9 0.4 4.9 (S) 4.8 (S) 4.65 (S)
9 0.5 5.25 (S) 4.9 (S) 4.67 (S)
9 0.6 5.73 (A) 5.4 (S) 5.22 (S)
10 0.3 4.85 (S) 4.55 (S) 4.55 (S)
10 0.4 6.3 (A) 6.35 (A) 5.6 (A)
10 0.5 6.95 (A) 6.65 (A) 6.6 (A)
10 0.6 7.35 (A) 7.25 (A) 7.1 (A)
11 0.3 6.75 (A) 6 (A) 5.9 (A)
11 0.4 7.5 (A) 7.3 (A) 7.1 (A)
11 0.5 7.55 (A) 7.5 (A) 7.2 (A)
11 0.6 8.5 (A) 7.75 (A) 7.95 (A)
12 0.3 7.65 (E) 7.35 (E) 7 (E)
12 0.4 8.2 (E) 7.95 (E) 7.95 (E)
12 0.5 9 (E) 8.95 (E) 8.75 (E)
12 0.6 8.8 (E) 8.9 (E) 8.9 (E)
Figure 3: Weld-nugget size versus welding time and welding current
at electrode forces of: a) 4.1 kN, b) 5.1 kN and c) 5.7 kN
Slika 3: Velikost pretaljenega jedra zvara v odvisnosti od ~asa varjenja
in varilnega toka pri sili elektrode: a) 4,1 kN, b) 5,1 kN in c) 5,7 kN
where Fe is the electrode force. It is expected that the
higher the heat factor, the higher is the weld-nugget
size. As it can be seen in Figure 4, the weld-nugget
growth is proportional to the heat factor, with the ex-
ception of high heat factors. This can be explained with
the fact that increasing the heat factor increases the
probability of expulsion. Expulsion can increase the
heat losses. Therefore, increasing the heat factor beyond
the critical value does not increase the weld-nugget size.
Therefore, it can be concluded that there is not a
proportional relationship between the heat factor and the
weld-nugget size. This is important when selecting the
optimum welding condition to obtain a larger weld-
nugget size.
3.3 Quantitative relation between welding parameters
and the weld-nugget size
To establish a relationship between the weld-nugget
size and the welding parameters, viz., the welding time,
the welding current and the electrode force, the follow-
ing relation was developed using multiple regression:
Weld Nugget Size 4 1.18 W
W
= + +
+ −
06252 2121
569533
.
.
I
t 3 07395 0 2568 0 063452 2. . .F I F Ie W e W+ −
(5)
In Figure 5, the values of the weld-nugget size ob-
tained experimentally are plotted against the weld-
nugget size predicted with Equation (5). As can be seen,
there is very little scatter of the points from the proposed
equation. Deviations are well within the 95 % confidence
limit.
The ability to make a weld, based on the welding
parameters and under production conditions, is best
defined in terms of a šweldability lobe’. The weldability
lobe defines the available tolerances for producing a
weld of a defined quality. In this way, it is possible to
determine the welding parameters that allow an accept-
able weld quality as defined with precise physical limits,
such as the weld size, or non-marking or aesthetic
qualities, associated with the amount of surface inden-
tation.4 According to the AWS D8.1M standard16 for
automotive weld-quality resistance spot welding of
steels, the lower limit of a lobe diagram corresponds to
the welding condition leading to the welds with the
nugget size larger than 4t0.5, where t is the sheet thick-
ness. The upper limits outlining the tolerance box for
acceptable welding are generally defined in terms of the
weld-nugget expulsion.
In Table 2, the range of the welding parameters pro-
ducing reliable spot welds is highlighted in blue. The
welding condition producing small welds is in gray.
Also, the expulsion occurrence is highlighted in red. The
suitable welding-current range is from the current, under
which the minimum nugget diameter (for example, 4t0.5)
is formed to that, under which the expulsion occurs. A
wide suitable welding-current range is desirable because
it is possible to control the nugget diameter within a pre-
scribed range even if the welding current fluctuates. The
welding range for each welding time is about 2 kA,
which is a proper welding-current range indicating good
weldability.
4 CONCLUSIONS
Understanding the influence of RSW parameters on
the weld-nugget growth during resistance spot welding is
a prerequisite for the development of the optimum weld-
ing conditions, ensuring high levels of joint quality in
auto-body manufacture. The results of the present re-
search revealed how the weld-nugget size is influenced
by the main welding parameters, viz., the welding
current, the welding time and the electrode force:
1) Increasing the heat input caused by increasing the
welding current and the welding time led to an en-
largement of the weld nugget due to increasing the
heat generated at the sheet/sheet interface.
S. P. HOVEIDA MARASHI: DISSIMILAR SPOT WELDING OF DQSK/DP600 STEELS: THE WELD-NUGGET GROWTH
764 Materiali in tehnologije / Materials and technology 50 (2016) 5, 761–765
Figure 4: Weld-nugget size versus heat factor (HF)
Slika 4: Velikost pretaljenega jedra zvara v odvisnosti od faktorja
toplote (HF)
Figure 5: Scatter plot for the weld-nugget size
Slika 5: Diagram raztrosa velikosti staljenega jedra zvara
2) Increasing the electrode force can increase the initial
sheet/sheet contact areas and therefore decrease the
sheet/sheet interfacial electrical resistively, which in
turn leads to a reduction in the generated heat at the
sheet/sheet interface. In other words, increasing the
electrode force increases the welding current and
welding time required to melt the sheet/sheet inter-
face.
3) We determined a relation involving the weld-nugget
size and the welding parameters, viz., the welding
current, the welding time and the electrode force.
This helped us to evaluate the combined effect of the
welding parameters on the weld-nugget size. Using
such a quantitative relation, the selection of the opti-
mum welding condition becomes straightforward.
4) Another factor, the heat factor = I2t/F, was defined to
evaluate the combining effect of the welding parame-
ters on the weld-nugget size.
5) The weldability lobe for dissimilar-resistance spot
welding of DP600 and low-carbon steel was deter-
mined using the established criteria of the AWS
standard. A wide welding-current range was estab-
lished indicating good weldability.
5 REFERENCES
1 M. Pouranvari, H. R. Asgari, S. M. Mosavizadeh, P. H. Marashi, M.
Goodarzi, Effect of weld nugget size on overload failure mode of re-
sistance spot welds, Sci. Technol. Weld. Joining, 12 (2007),
217–225, doi:10.1179/174329307x164409
2 M. Pouranvari, Susceptibility to interfacial failure mode in similar
and dissimilar resistance spot welds of DP600 dual phase steel and
low carbon steel during cross-tension and tensile-shear loading
conditions, Mater. Sci. Eng. A, 546 (2012), 129–138, doi:10.1016/
j.msea.2012.03.040
3 M. Pouranvari, E. Ranjbarnoodeh, Dependence of Fracture Mode on
Welding Variables in Resistance Spot Welding of DP980 Advanced
High Strength Steel, Mater. Tehnol., 46 (2012), 665–671
4 N. T. Williams, J. D. Parker, Review of resistance spot welding of
steel sheets, Part 1: Modelling and control of weld nugget formation,
International Materials Reviews, 49 (2004), 45–75, doi:10.1179/
095066004225010523
5 J. C. Feng, Y. R. Wang, Z. D. Zhang, Nugget growth characteristic
for AZ31B magnesium alloy during resistance spot welding, Sci.
Technol. Weld. Joining, 11 (2006), 154–162, doi:10.1179/
174329306x84364
6 M. Pouranvari, S. P. H. Marashi, Critical review of automotive steels
spot welding: process, structure and properties, Sci. Technol. Weld.
Joining, 18 (2013), 361–403, doi:10.1179/1362171813y.0000000120
7 H. Zhang, J. Senkara, Resistance welding: fundamentals and appli-
cations, Taylor & Francis CRC Press, ........... 2005
8 A. De, O. P. Gupta, L. Dorn, An experimental study of resistance
spot welding in 1 mm thick sheet of low carbon steel, Proc. Inst.
Mech. Engrs., Part B: Journal of Engineering Manufacture, 210
(1996), 341–347, doi:10.1243/pime_proc_1996_210_126_02
9 M. Pouranvari, S. P. H. Marashi, On the failure of low carbon steel
resistance spot welds in quasi-static tensile-shear loading, Materials
& Design, 31 (2010), 3647–3652, doi:10.1016/j.matdes.2010.02.044
10 M. Pouranvari, S. P. H. Marashi, S. M. Mousavizadeh, Failure mode
transition and mechanical properties of similar and dissimilar
resistance spot welds of DP600 and low carbon steels, Science and
Technology of Welding & Joining, 15 (2010), 625–631, doi:10.1179/
136217110x12813393169534
11 M. Pouranvari, S. P. H. Marashi, Failure mode transition in AISI 304
resistance spot welds, Welding Journal, 91 (2012), 303–309
12 M. Pouranvari, S. P. H. Marashi, Factors affecting mechanical pro-
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26 (2010), 1137–1144, doi:10.1179/174328409x459301
13 J. Heuschkel, The expression of spot-weld properties, Welding Jour-
nal, 31 (1952), 931–943
14 J. M. Sawhill, J. C. Baker, Spot weldability of high-strength sheet
steels, Welding Journal, 59 (1980), 19–30
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electrical contact resistance in resistance, Welding Journal, 85
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steel, AWS D8:1M, New York, American National Standard, 2007
S. P. HOVEIDA MARASHI: DISSIMILAR SPOT WELDING OF DQSK/DP600 STEELS: THE WELD-NUGGET GROWTH
Materiali in tehnologije / Materials and technology 50 (2016) 5, 761–765 765

T. LAZAR et al.: ARMOUR PLATES FROM KOZLOV ROB – ANALYSES OF TWO UNUSUAL FINDS
767–773
ARMOUR PLATES FROM KOZLOV ROB – ANALYSES OF TWO
UNUSUAL FINDS
OKLEPNI PLO[^I S KOZLOVEGA ROBA – ANALIZE DVEH
NENAVADNIH NAJDB
Toma` Lazar1, Primo` Mrvar2, Martin Lamut3, Peter Fajfar2
1National Museum of Slovenia, Pre{ernova cesta 20, 1000 Ljubljana, Slovenia
2University of Ljubljana, Faculty of Natural Sciences and Engineering, Department of Material Science and Metallurgy, A{ker~eva cesta 12,
1000 Ljubljana, Slovenia
3Slovenian centre of excellence for space sciences and technologies, A{ker~eva cesta 12, 1000 Ljubljana, Slovenia
peter.fajfar@omm.ntf.uni-lj.si
Prejem rokopisa – received: 2015-07-30; sprejem za objavo – accepted for publication: 2015-10-13
doi:10.17222/mit.2015.242
During archaeological excavations of the fortifications on Kozlov rob, two perforated steel plates were discovered, the purposes
of which had never been explained satisfactorily. Detailed examinations and scientific analyses confirmed that in one case at
least we were dealing with fragments of armour from the late Middle Ages or the early Modern Period that had later been
reworked into an entirely unrelated object having, in all possibility, a non-warlike function.
Keywords: armour, weapons, archaeometallurgy, metallography, nano-indentation, Kozlov rob
Med arheolo{kimi izkopavanji utrdbe na Kozlovem robu sta bili odkriti dve preluknjani, jekleni plo{~i, katerih namembnost ni
bila nikoli zadovoljivo pojasnjena. Natan~ne preiskave in znanstvene analize so potrdile, da gre v vsaj enem primeru za
fragment oklepa iz poznega srednjega veka ali iz zgodnjega novega veka. Ta je bil naknadno predelan v predmet, ki po vsej
verjetnosti ni slu`il voja{kemu namenu.
Klju~ne besede: oklep, oro`je, arheometalurgija, metalografija nanovtiskovanje, Kozlov rob
1 INTRODUCTION
Kozlov rob is a lone hill of almost rectangular form
reaching an altitude of 426 m above sea-level overlook-
ing the town of Tolmin. Due to the strategic importance
of the location, a castle was built on Kozlov rob in the
12th Century. The castle remained in use until the
mid-17th Century.1
Archaeological excavations at Kozlov rob were car-
ried out in 1964/5 and 1996/7. The excavations revealed
various finds.2 As elsewhere in Central Europe,3,4,5 the
finds consisted mostly of small, often fragmentary ob-
jects. Among those related to warfare, the largest group
is represented by crossbow bolts.6 In addition to other
fragments of arms and armour, two steel plates were
found, both perforated with rows of holes. One of them
was first reported on in 2008.1,7
The unclear purposes of the plates soon stimulated
lively discussion. The first plate evidently represents the
upper half of a breastplate. The second plate is made
from much thicker steel sheet but lacks any recognisable
features. The two fragments have little in common apart
from the more or less symmetrically distributed holes.
This is an indication that both objects had performed an
identical function during their last period of working life.
At least three hypotheses have been suggested as to their
intended purpose:
1. Fragments of armour were used by the castle garrison
for target practice with crossbows.
2. The perforated breastplate with vent holes is an
example of extremely rare tournament armour for
foot combat.
3. Worn-out fragments of armour were recycled and
modified into something else, perhaps a grate or vent.
In 2010 the plates from Kozlov rob were submitted to
in-depth research for the first time. Four small samples
were removed from each plate and subjected to me-
tallographic analysis. The results indicated that both
plates were made from relatively high-quality steel con-
taining approximately 0.3 % to 0.8 % carbon with a
ferritic-pearlitic structure.8,9
These preliminary results underlined the need for
additional analyses and a more detailed publication
about the plates from Kozlov rob. The following paper
has been compiled to present a comprehensive overview
of the analytical methods used as well as an interdiscip-
linary discussion of the latest findings.
2 EXPERIMENTAL PART
The fragment of breastplate No. 1 (Figure 1) is heav-
ily corroded. It represents the upper half of a steel
cuirass reaching up to the folded rim of the neck and
armpit cut-out. The fragment originally belonged to a
Materiali in tehnologije / Materials and technology 50 (2016) 5, 767–773 767
UDK 620.17:623.445.1:67.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)767(2016)
breastplate of comparatively plain construction, possibly
of North Italian manufacture, in the style typical for the
period around 1500 AD.8,10 It is made from one piece of
steel sheet, approximately 1.5 mm to 3 mm thick. The
length of the plate is 252 mm, the width is 286 mm, and
its weight is 778 g. Due to a progressive state of corro-
sion, the original thickness of the sheet metal is difficult
to measure precisely, but can be estimated at around
2 mm on average. The convex-shaped steel sheet is per-
forated along two thirds of its width on the right-hand
side with 18 holes arranged in more or less symmetrical
rows. The holes are mostly of uniform, square shaped,
pierced from the inner side of the breastplate.1,9
The fragment of steel plate No. 2 (Figure 2) is
shaped as an irregular parallelogram. It is made from a
massive steel sheet with a multi-layered anisotropic mi-
crostructure. The plate is perforated with 16 square and
round holes pierced from both sides. Due to advanced
corrosion the metal sheet has largely disintegrated longi-
tudinally into multiple layers. The surface is covered
with a thick layer of corrosion products. The thickness of
the steel sheet is approximately 3.5 mm to 7 mm, with an
average of around 4.5 mm. The length of the plate is
245 mm, the width 244 mm, and its weight is 1.504 g.8,9
For the microstructural characterization and the
mechanical analysis three small samples were removed
from each steel plate. The locations of the samples of
fragment No. 1 are shown in Figure 1, and their shapes
are shown in Figure 3, and the locations of the samples
of fragment No. 2 are shown in Figure 2, and their
shapes are shown in Figure 4. The obtained samples
were embedded in resin, ground, polished and etched
using a solution of 2 % nital. The microstructural cha-
racterization was performed using light optical micro-
scopy (Olympus Microscope BX61). The tests for me-
chanical analysis were conducted on an Agilent G200
Nanoindenter using the continuous stiffness measure-
ment (CSM) option. The CSM technique allows the
contact stiffness together with the projected area of
hardness impressions to be measured at any point along
the loading curve.11,12 This instrument monitors and re-
cords the dynamic load and displacement of the indenter
during indentation with a force resolution of approxi-
mately 50 nN and a displacement resolution of approxi-
mately 0.01 nm.
3 RESULTS
3.1 Metallographic analysis
Sample 1 originates from the edge of a heavily cor-
roded hole. The examined surface runs along the plane
of the plate. The sample is corroded, hence the micro-
structure of the metal is invisible on the macro photo-
graph (Figure 3a). Sample 2 was removed from the
lower, broken part of the plate. The examined surface
T. LAZAR et al.: ARMOUR PLATES FROM KOZLOV ROB – ANALYSES OF TWO UNUSUAL FINDS
768 Materiali in tehnologije / Materials and technology 50 (2016) 5, 767–773
Figure 1: Fragment of breastplate showing the locations of the
samples (Photo: T. Lazar), Gori{ki muzej, Nova Gorica, s. n.
Slika 1: Odlomek prsne plo{~e z mesti odvzema vzorcev (foto: T.
Lazar), Gori{ki muzej, Nova Gorica, s. n.
Figure 3: Samples from the plate No. 1: a) sample 1, b) sample 2,
c) sample 3
Slika 3: Vzorci plo{~e {t. 1: a) vzorec 1, b) vzorec 2, c) vzorec 3
Figure 2: Fragment of steel plate showing the locations of the samples
(Photo: T. Lazar), Gori{ki muzej, Nova Gorica, s. n.
Slika 2: Odlomek jeklene plo{~e z mesti odvzema vzorcev (Photo: T.
Lazar), Gori{ki muzej, Nova Gorica, s. n.
runs perpendicular to the plane of the plate (Figure 5).
The microstructure contains mostly pearlite (darker
microstructural areas of the grain) with a small propor-
tion of ferrite (white microstructural areas). On the lower
edge of the sample the grain is finely formed within a
directional structure. Also, the proportion of ferrite is
markedly larger. The dark area along the left edge repre-
sents the corroded surface. Sample 3 was taken from the
right edge of the plate. The examined surface runs along
the plane of the plate (Figure 6). The microstructure
contains mostly pearlite with a small proportion of
ferrite. On the upper edge of the sample (the outside
edge of the plate) the grain is smaller and the proportion
of ferrite larger. The dark area around the right edge
represents the corroded surface. The existing corrosion
products are formed mostly on the basis of Fe2O3 and
Fe3O4.
Within the narrow band along the lower right edge,
the microstructure of sample 2 is directed along the flow
of material produced during the perforation process (Fig-
ure 5). The grain is plastically deformed, non-recry-
stallized, as is typical of cold working. Given that the
edge layers contain a significantly higher portion of
recrystallized ferrite grain, it can be estimated that the
steel plate was heated to a temperature between 700 °C
and 900 °C prior to the perforation process. At that
temperature and a suitable time of heating the steel plate
was partly decarburised on the surface. This is confirmed
by the microstructure containing an increased proportion
of ferrite. The narrow band of non-recrystallized grain
may have been formed during a later repair of the hole
done cold.
Sample 4 originates from the edge of a round hole.
The examined surface runs perpendicularly along the
plane of the plate (Figure 7). In the centre of the sample
the microstructure contains mostly pearlite with a small
proportion of ferrite. In the outer layers the grain is
formed more finely and there is an increased content of
ferrite. Sample 5 was removed from the left edge of the
plate. The examined surface runs perpendicular to the
plane of the plate (Figure 8). Along the entire cross-
section pearlite can be observed, indicating that the steel
has a composition containing 0.8 % carbon. Sample 6
was taken from the edge of a square hole. The examined
surface runs perpendicular to the plane of the plate (Fig-
ure 9). The microstructure is fine grained, containing
pearlite and ferrite. The edges of the sample contain
ferrite and the grain is larger, too.
The microstructure of sample 5 is predominantly
pearlitic. The examined surface is directed towards the
interior of the plate. Such a microstructure is typical of
hot working. The micro photographs demonstrate that
T. LAZAR et al.: ARMOUR PLATES FROM KOZLOV ROB – ANALYSES OF TWO UNUSUAL FINDS
Materiali in tehnologije / Materials and technology 50 (2016) 5, 767–773 769
Figure 7: Sample 4, plate No. 2
Slika 7: Vzorec 4, plo{~a {t. 1
Figure 5: Sample 2, plate No. 1
Slika 5: Vzorec 2, plo{~a {t. 1
Figure 6: Sample 3, plate No. 1
Slika 6: Vzorec 3, plo{~a {t. 1
Figure 4: Samples from plate No. 2: a) sample 4, b) sample 5,
c) sample 6
Slika 4: Vzorci plo{~e {t. 2: a) vzorec 4, b) vzorec 5, c) vzorec 6
the plate was formed by hot forging, then slowly air
cooled. The holes on the second plate were punched hot
as well. This is confirmed by the compositions of sam-
ples 4 and 6 that contain fine-grained microstructures
with an increased ferrite content in the outer layers,
which are partly decarburised. The grain is larger than is
the case with the first plate. The proportion of ferrite is
larger, too. Assuming that both plates were heated to the
same temperature during the manufacture, the larger
grain size may be attributed to a longer heating time, as
plate No. 2 is twice as thick as No. 1.
The examined samples also contain non-metallic
inclusions that were oriented perpendicularly to the
direction of deformation during the sequences of hot
forging.13,14
Figure 7 demonstrates that these inclusions are rela-
tively soft and therefore were deformed considerably.
3.2 Mechanical analysis
In order to obtain a deeper insight into the origin of
the archaeological finds in the form of two perforated
steel plates it was decided to obtain some mechanical
data. Since the archaeologists and museum curators are
not in favour of any material being detached from mu-
seum objects, the amount material to work with is
usually fairly small, hence it is difficult to extract its
mechanical properties. Due to such constraints the
nanoindentation technique was used, which can provide
a Young’s modulus and a hardness from very small sam-
ples.
The penetration depth of 1 μm was carefully chosen
in order to avoid the influences of the indentation size
effect, the surface roughness or the surface inclination
and to reach a constant value or plateau region. At each
point at least 6 measurements were taken. In spots where
the results showed extensive dissipation, the number of
measurements was increased considerably.
The first point of interest was plate No. 1, sample 2
(P1-S2), removed from the area next to a hole, where at
the edge finely formed grains are found with an in-
creased portion of ferrite and in the middle part larger
grains with a predominantly pearlitic microstructure. The
results are presented in Table 1. In order to eliminate
point fluctuations in the load-displacement curve the
values are averages at depths between 800 nm and
900 nm from each test. (Note that for the nanoindenta-
tion range the hardness values are greater than for the
micro or macro scale11,12). The hardness measurements of
P1-S2 (Figure 10) show a large dissipation at the edge in
comparison to the area in the middle. It shows a slightly
larger mean hardness (Table 1), but considering the
dissipation it all falls into the margin of error. Since
ferrite has a smaller hardness than pearlite it is assumed
that the influence of cold-worked smaller grains is mini-
mized by a greater ferrite share in the microstructure.
Furthermore, the tests were made separately at the
ferritic and pearlitic grains, but the results overlapped,
most likely due to effects from adjacent grains, so it was
T. LAZAR et al.: ARMOUR PLATES FROM KOZLOV ROB – ANALYSES OF TWO UNUSUAL FINDS
770 Materiali in tehnologije / Materials and technology 50 (2016) 5, 767–773
Figure 9: Sample 6, plate No. 2
Slika 9: Vzorec 6, plo{~a {t. 1
Figure 10: Hardness as a function of contact depth for P1S2: a) the
edge and b) the middle
Slika 10: Trdota v odvisnosti od kontaktne globine za P1S2: a) rob in
b) sredina
Figure 8: Sample 5, plate No. 2
Slika 8: Vzorec 5, plo{~a {t. 1
hard to draw any sound conclusions. Sample 3 (P1-S3),
taken away from the afterwards induced holes, shows a
considerably smaller hardness. The grains remained in
the original state and no plastic deformation occurred,
which agrees well with the above hypothesis.
Table 1: Hardness tests statistics on the first plate
Tabela 1: Statistika preizkusov trdote za prvo plo{~o
P1-S2 middle P1-S2 edge P1-S3
Mean 3.18 3.38 2.29
Std. Dev. 0.24 0.53 0.15
% COV 7.48 15.82 6.36
In Table 2 the hardness results from plate No. 2 are
given. Samples 4 and 6 (P2-S4, P2-S6), taken from the
vicinity of a hole, show higher hardness where the grains
are smaller with a lower share of ferrite, which in these
cases is in the middle of the samples. Where the grains
are larger and a higher share of ferrite is observed, the
hardness is lower, comparable with the samples P1-S3
and P2-S5, taken from the area with no subsequent
reworking of the plates.
Table 2: Hardness tests statistics on the second plate
Tabela 2: Statistika preizkusov trdote za drugo plo{~o
P2-S4
middle
P2-S4
edge P2-S5
P2-S6
middle
P2-S6
edge
Mean 3.05 2.69 2.82 2.97 2.57
Std. Dev. 0.08 0.09 0.3 0.3 0.23
% COV 2.56 3.24 10.53 10.1 9
4 DISCUSSION
Fragments of armour were used by the castle garrison
for target practice with crossbows.
Both steel plates are perforated with holes, mostly of
a square shape, measuring approximately 10 mm × 10 mm
(Figures 11 and 12). Such dimensions correspond to the
typical late-medieval crossbow bolt with a square- or
diamond-shaped head. A sizeable group of bolt heads of
that same type was unearthed on Kozlov rob6,7, seem-
ingly supporting the idea that the plates had been used
for marksmanship practice.
However, several arguments speak against such an
explanation. On both plates the holes are placed quite
neatly in straight rows. Such precise shot placement
would be difficult to achieve even by a skilled marks-
man. At any rate, the arrangement of holes is much more
symmetrical than might be expected from a random
dispersion of hits.
The penetration of late medieval crossbows should
not be overestimated. The spanning force of a "one-foot"
military crossbow was within the range of 1500 N, or
2100 N in the case of the heavier "two-foot" type. In the
15th century, even more powerful crossbows were deve-
loped with draw weights up to 5000 N.16–20 However,
their efficiency was low. At point-blank range, the maxi-
mum penetration of a typical crossbow bolt weighing 70
g might reach up to 60 mm of pinewood. Documented
cases of projectiles stuck within the wooden structures of
medieval castles show a degree of penetration averaging
only 20 mm in seasoned softwood.16
The kinetic energy of a crossbow bolt21 at point-blank
range may be estimated at around 200 J.16 The latter is
clearly superior to a yew longbow with a kinetic energy
of 120 J at a draw weight of 670 N.22 However, it would
be barely sufficient to penetrate even the relatively thin
plate No. 1. The second plate, at least twice as thick,
would prove entirely resistant to crossbow bolts. In fact,
it would be proof against a musket ball possessing at
least ten times greater kinetic energy.23
Therefore, it seems highly unlikely that the two steel
plates from Kozlov rob were pierced by crossbow bolts.
The final answer is provided by the metallographic
analysis. Had the holes been made by the impact of
crossbow bolts the samples would show evidence of
mechanical deformation and cold work hardening of the
microstructure.
The perforated breastplate with vent holes is an
example of extremely rare tournament armour for foot
combat.
In the late Middle Ages, new types of armour were
developed specifically for tournament use. The famous
T. LAZAR et al.: ARMOUR PLATES FROM KOZLOV ROB – ANALYSES OF TWO UNUSUAL FINDS
Materiali in tehnologije / Materials and technology 50 (2016) 5, 767–773 771
Figure 12: Holes on plate No. 1 (Photo: T. Lazar)
Slika 12: Luknje na plo{~i {t. 1 (foto: T. Lazar)
Figure 11: Square-shaped holes of plate No. 2 (Photo: D. Todorovi})
Slika 11: Luknje kvadratne oblike na plo{~i {t. 2 (foto: D. Todorovi})
tournament book of René of Anjou from ca. 1460 depicts
a cuirass with numerous vent holes for fighting on foot
(Figure 13).24,25 Recent archaeological excavations at the
Haus Herbede in the Ruhr district have revealed a breast-
plate of the exact same type.26,27 The discovery of a per-
forated breastplate at Kozlov rob gave rise to speculation
that it might have belonged to a similar suit of tourna-
ment armour designed to prevent overheating.1,7
Nonetheless, a closer examination of the breastplate
disproved such a possibility. The holes are crudely made,
which is clearly inconsistent with the workmanship of
late-medieval tournament armour. At any rate, the breast-
plate would be very uncomfortable to wear due to the
sharp perforations. The second perforated plate is an
even less sophisticated product. The holes were struck
into the steel sheet from both sides, in many cases, form-
ing crude bulges up to a centimetre deep. Such a plate
could never have belonged to armour worn on one’s
person (Figure 14).
Worn-out fragments of armour were recycled and re-
worked into something else, perhaps a grate or vent.
Steel was an expensive commodity in the past, hence
recycling of old or damaged steel products was part of
everyday life. Provided that a damaged fragment of
armour had to be scrapped, it seems plausible that it was
reworked into a new product, even something as mun-
dane as a fireplace grate or a vent. Two very similar
plates used for that purpose were spotted coincidentally
in 2010, built into a modern brick structure on the island
of Rab (Figures 15a and 15b).
The perforations on the plates were clearly made
using a massive chisel or punch. Due to the thickness of
the metal sheet, particularly on plate No. 2., it can be
T. LAZAR et al.: ARMOUR PLATES FROM KOZLOV ROB – ANALYSES OF TWO UNUSUAL FINDS
772 Materiali in tehnologije / Materials and technology 50 (2016) 5, 767–773
Figure 13: Armour for foot combat as depicted in René of Anjou’s
tournament book (Bibliothèque Nationale, Paris, MS Fr 2695, fol.
25v)
Slika 13: Oklep za pehotni dvoboj iz turnirske knjige Renéja An`uj-
skega (Bibliothèque Nationale, Paris, MS Fr 2695, fol. 25v)
Figure 14: Side view of plate No. 2 (Photo: D. Todorovi})
Slika 14: Stranski pogled plo{~e {t. 2 (foto: D. Todorovi})
Figure 15: a), b): Unexpected analogy from the Adriatic coast: two
simple steel vents, designed in the same manner, built into a modern
outbuilding on the island Rab (Photo: T. Lazar)
Slika 15: a), b): Nepri~akovana podobnost z Jadranske obale: dva pre-
prosta jeklena zra~nika, oblikovana na enak na~in in vgrajena v
moderno zgradbo na otoku Rab (foto: T. Lazar)
assumed that the punching was performed while the
plate was heated in a forge. Even so, the process would
have necessitated heavy hammer blows, causing severe
deformation of the material. Further explanation of the
work process may be provided by scientific analyses, in
particular metallography.13,14,28,29
5 CONCLUSION
The fragmentary breastplate No. 1 was a quality pro-
duct made of steel with a relatively high carbon con-
tent.16 The latter also holds true for plate No. 2, even
though, based on its shape alone, it is impossible to
determine what sort of object it had belonged to origi-
nally. One cannot exclude the possibility that it was
merely a steel semi-product.8
The results of the metallographic analyses proved
that both plates had been reworked in a similar manner.
As could be expected given the technological capabilities
of the pre-industrial era, the holes were made by hot
punching. Only later on, as indicated by the non-recry-
stallization, the directional structure of the part of sample
2 (Figure 5), some of the holes may have been repaired
or reworked slightly by cold working. The mechanical
analyses show that in the case of small samples the
nanoindentation technique can provide useful additional
data. The results confirm the metallographic analyses
and further reveal the material properties as well as the
methods of making and processing the plates.
Based on the above observations, it is clear that once
the holes had been punched and formed, at least plate
No. 1 was not exposed to temperatures in excess of
approximately 400 °C that would have caused recry-
stallization of the cold-worked areas. Therefore, it is
possible to conclude that the perforated plates were most
likely used as simple grates, vents or were put to some
other similar uses.
The reworking of the steel plates seems all the more
plausible considering the gradual decline of armour in
Europe during the 17th century. From that standpoint it
might be easier to understand why on Kozlov rob a
breastplate was reworked into a new product, no longer
having anything in common with the armour’s original
purpose.
6 REFERENCES
1 B. @bona Trkman, F. Bressan, Oro`je z gradu Kozlov rob; Vojske,
oro`je in utrdbeni sistemi v Poso~ju, Tolminski muzej, Tolmin 2008,
58
2 B. @bona Trkman, A. Kruh, Stanje raziskav gradov in dvorcev na
obmo~ju histori~ne Gori{ke I. Dobrovo, Kozlov rob, Rihemberk,
[tanjel, Gori{ki letnik, Nova Gorica, 33–34 (2010), 195–217
3 K. Predovnik, Trdnjava Kostanjevica na Starem gradu nad Pod-
bo~jem, Filozofska fakulteta, Ljubljana 2003, 240
4 C. Krauskopf, Plemstvo in predmeti iz njegovega vsakdanjika.
Raziskave materialne kulture 13. in 14. stoletja, Arheo, Ljubljana, 23
(2005), 47–62
5 T. Lazar, Boji{~a visokega in poznega srednjega veka kot izziv slo-
venski arheologiji, Arheo, Ljubljana, 28 (2011), 119–130
6 F. Bressan, Le cuspidi di freccia nel Friuli medievale, Università
degli Studi di Trieste, Trieste 1996, 275
7 T. Lazar, Voja{ka zgodovina slovenskega ozemlja od 13. do 15.
stoletja, Filozofska fakulteta, Ljubljana 2009, cat. no. M 315–346,
442
8 A. Williams: Metalur{ke zna~ilnosti poznosrednjeve{kih oklepov iz
srednje Evrope, Vitez, dama in zmaj; Dedi{~ina srednjeve{kih bojev-
nikov na Slovenskem, 1, Razprave, Ljubljana 2011, 233–247
9 T. Lazar, T. Nabergoj, P. Bitenc, Vitez, dama in zmaj; Dedi{~ina
srednjeve{kih bojevnikov na Slovenskem. 2, Katalog predmetov,
Narodni muzej Slovenije, Ljubljana 2013, cat. no. 256, 257
10 J. Mann, Wallace Collection Catalogues, European Arms and Ar-
mour, Vol I., Armour, The Trustees of the Wallace Collection,
London 1962, cat. no. A 214
11 W. C. Oliver, G. M. Pharr, An Improved Technique for Determining
Hardness and Elastic Modulus Using Load and Displacement
Sensing Indentation Experiments, J. Mater. Res., 7 (1992),
1564–1583, doi:10.1557/JMR.1992.1564
12 R. Rodríguez, I. Gutierrez, Correlation Between Nanoindentation
and Tensile Properties. Influence of the Indentation Size Effect,
Materials Science and Engineering, A361 (2003), 377–384,
doi:10.1016/S0921-5093(03)00563
13 M. Ne~emer, T. Lazar, @. [mit, P. Kump, B. @u`ek, Study of the
Provenance and Technology of Asian Kris Daggers by Application of
X-Ray Analytical Techniques and Hardness Testing, Acta Chim.
Slov., 351 (2013) 60, (2), 351–357
14 Fajfar, J. Medved, G. Klan~nik, T. Lazar, M. Ne~emer, P. Mrvar,
Characterization of a Messer – the late-Medieval single-edged sword
of Central Europe, Materials Characterization, 86 (2013), 232–241,
doi:10.1016/j.matchar. 2013.10.005
15 E. Duka, H. Oettel, T. Dilo, Connection Between Micro and Macro
Hardness. Pearlitic-Ferritic Steel, AIP Conference Proceedings 1476,
2012, 47–5, doi:10.1063/1.4751563
16 E. Harmuth, Die Armbrust, Akademische Druck- und Verlaganstalt,
Graz 1986, p. 232
17 W. F. Paterson, A Guide to the Crossbow, Society of Archer-Anti-
quaries, Burnham 1990, 132
18 J. Alm, European Crossbows, Royal Armouries, Leeds 1996, 25
19 J. Liebel, Springalds and Great Crossbows, Royal Armouries, Leeds
1998, 23
20 H. Richter, Die Hornbogenarmbrust, Hoernig Angelika, Ludwigs-
hafen 2006, 190
21 A. Williams, The Knight and the Blast Furnace, Brill, Leiden,
Boston 2003, 954
22 M. Strickland, R. Hardy, The Great Warbow; From Hastings to the
Mary Rose, Sutton, Stroud 2005, 538
23 P. Krenn, Von alten Handfeuerwaffen, Landeszeughaus, Graz 1989,
149
24 Bibliothèque Nationale, Paris, MS Fr 2695
25 J. Heers, F. Robin, René d’Anjou, Traité des Tournois, Lengenfelder,
München 1993, 30
26 H. W. Peine, 2004: Ein Blick in die Waffenkammer des Hauses Her-
bede an der Ruhr, Das Brigantinen-Symposium auf Schloss Tirol/Il
simposio sulla brigantina a Castel Tirolo, Innsbruck, 2004, 51–53
27 C. Blair, European Armour, B. T. Batsford Ltd., London 1958, 248
28 D. A. Scott, Metallography and Microstructure of Ancient and
Historic Metals, The J. Paul Getty Trust, Los Angeles 1991, 155
29 T. Lazar, N. Neme~ek, Interdisciplinary Research of Museum
Objects. Practical Experience with Various Analytical Methods,
RMZ: Materials and geoenvironment, 60 (2013) 4, 249–261
T. LAZAR et al.: ARMOUR PLATES FROM KOZLOV ROB – ANALYSES OF TWO UNUSUAL FINDS
Materiali in tehnologije / Materials and technology 50 (2016) 5, 767–773 773

D. DANYALI, E. RANJBARNODEH: NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE EFFECT ...
775–782
NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE
EFFECT OF HYDROSTATIC PRESSURE ON THE RESIDUAL
STRESS IN BOILER-TUBE WELDS
NUMERI^NA IN EKSPERIMENTALNA PREISKAVA VPLIVA
HIDROSTATI^NEGA TLAKA NA ZAOSTALE NAPETOSTI V
ZVARU NA KOTLOVSKI CEVI
Daryoush Danyali, Eslam Ranjbarnodeh
Amir Kabir of University of Technology, Department of Mining and Metallurgical Engineering, Tehran, Iran
islam_ranjbar@yahoo.com, islam_ranjbar@aut.ac.ir
Prejem rokopisa – received: 2015-08-13; sprejem za objavo – accepted for publication: 2015-10-09
doi:10.17222/mit.2015.255
A tube to tube-sheet weld joint is a critical section in many boilers. Leakage and failure are two common problems that can
occur with this type of joint. The tensile residual stresses associated with welding can play a major role in these problems. In the
present study, the Finite-Element Method is used to predict the residual stresses in a tube to tube-sheet weld and the effect of
hydrostatic pressure on the residual stresses’ redistribution. Investigations were performed by numerical analyses and
experimental methods. The joint included a circumferential fillet weld in one pass using a gas tungsten arc welding process. The
thermomechanical behavior of the joint is simulated with a two-dimensional axisymmetric model and a subroutine developed in
ANSYS software. The thermography method is used for the thermal verification and the hole-drilling strain gauge under
post-weld heat treatment is used for the mechanical analysis verification. The numerical and experimental results showed that
applied hydrostatic pressure can reduce, by about 58 %, the axial residual stress.
Keywords: hydrostatic test, residual stress, boiler tubes, tube-sheet, finite element method (FEM)
Kriti~ni podro~ji v kotlih sta zvarjen spoj cevi in cevne predelne stene. V tej vrsti spoja sta glavna problema, ki se pojavljata,
pu{~anje in poru{itev. Zaostale natezne napetosti povezane z varjenjem so glavni dejavnik pri teh problemih. V {tudiji je bila
uporabljena metoda kon~nih elementov za napovedovanje zaostalih napetosti v zvarih cevi in cevne predelne stene in vpliv
hidrostati~nega tlaka na razporeditev zaostale napetosti. Raziskave so bile izvedene z numeri~no analizo in z eksperimentalnimi
metodami. Spoj je vklju~eval obodni zvar izdelan v obloku z volframovo elektrodo in za{~ito s plinom. Termomehansko
obna{anje spoja je bilo simulirano z uporabo dvodimenzijskega osnosimetri~nega modela in programa razvitega z ANSYS
programsko opremo. Termografska metoda je uporabljena za preverjanje toplote in pri toplotni obdelavi zvarov je bila
uporabljena metoda z merilnimi listi~i na izvrtini za preverjanje mehanske analize. Numeri~ni in eksperimentalni rezultati so
pokazali, da uporabljen hidrostati~ni tlak, lahko za 58 % zmanj{a osne zaostale napetosti.
Klju~ne besede: hidrostati~ni preizkus, zaostala napetost, kotlovske cevi, cevna predelna stena, metoda kon~nih elementov
(FEM)
1 INTRODUCTION
Boilers are a very important part of many industries,
such as power generation, food, hospitals, and hotels.
Boilers are the most troublesome components of elec-
trical power generation plants. It costs the US utility
industry over $5 billion per year in unscheduled shut-
downs, repairs and power replacements.1 In general, the
safety and structural integrity of a boiler are of para-
mount importance. Welding is common method in the
boiler-manufacturing process and the process of welding
has a direct influence on the integrity of the components
and their thermal and mechanical behaviors during
service. Micro-cracking in tube to tube-sheet welds, in
both the radial and circumferential directions, is
commonplace.2
The nature of failures in the vicinity of the tube to
tube-sheet interfaces is particularly difficult to determine
because of the large temperature gradients as well as the
relatively high pressures. Due to the high temperatures
introduced during welding and the subsequent cooling of
the welded metal, welding can produce undesirable
residual stresses and deformations. Radial cracking is
more likely to occur, but circumferential cracking also
occurs and results from the thermal cycling, which is
exacerbated by a high residual stress.3 The thermal
stress, mechanical stress and welding residual stress of
tube to tube-sheet has received a lot of attention in recent
years.
B. Yildrim and H. F. Nied4 investigated the residual
stress and distortion in boiler-tube panels with welded
overlay cladding. X. Qingren5 reported that residual
stress in the circumferential weld in pipes is reduced when
hydrostatic pressure is increased. To the best of the
knowledge of the author, a very limited number of FEM
models have been proposed for predicting the residual
stresses in a tube to tube-sheet weld and the effect of
hydrostatic pressure test on the residual stresses in this
type of joint. The present study uses a two dimensional
Materiali in tehnologije / Materials and technology 50 (2016) 5, 775–782 775
UDK 519.6:620.179:67.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)775(2016)
axisymmetric FEM model for the calculation and
distribution of the residual stresses in a tube to
tube-sheet joint. The Thermography Method is used for
the thermal verification and the hole-drilling strain gauge
under post-weld heat treatment is used for the
mechanical analysis verification and then the effect of
hydrostatic pressure on the redistribution residual
stresses is investigated by numerical analyses and experi-
mental methods.
2 FEM AND VERIFICATION PROCEDURES
2.1 Welding verification procedures
To verify the results of the FEM method, an
experimental test with the same geometry, material and
welding parameters was produced, as shown in Figure 1.
2.2 Material properties
The base metal of the tube was steel St35.8 and
tube-sheet was steel 17Mn4. The chemical compositions
are listed in Table 1.
Table 1: Chemical composition in mass fractions (w/%)
Tabela 1: Kemijska sestava v masnih odstotkih (w/%)
Composition C Si Mn P S Cr
St35.8 0.17 0.35 0.60 0.03 0.03 –
17Mn4 0.20 0.20 0.120 0.03 0.025 0.30
For the thermal and mechanical analyses, the rela-
tionship between the thermo-physical and the thermo-
mechanical properties of the materials with temperature
is incorporated, as shown in Figure 2.
The welding parameters used were exactly as same as
the boiler-repair parameters. The dimensions of the
welded parts are presented in Table 2.
Table 2: The dimensions of the pieces used
Tabela 2: Dimenzije uporabljenih kosov
Material Diameter(mm)
Width
(mm)
Width
(mm)
Thickness
(mm)
Tube 51.8 – 200 3.2
Tube-sheet – 80 80 12
2.3 Thermal verification procedure
To verify the thermal model, thermography was used.
In this study, the welding temperature was monitored
with a model T8 thermography camera. It helps to record
the temperature during the welding and after the weld
cooling.
2.4 Mechanical model verification procedure
To verify the weld profile obtained using the FEM
method, a section was prepared for macroscopic exami-
nation of the welded samples. The section was prepared,
polished and etched in a solution with 50 mL of 37 %
hydrochloric acid and 50 ml of distilled water for 30
minutes. After the etching period it was examined
macroscopically.
2.5 Residual stress verification procedure
The hole-drilling strain gauge under post-weld heat
treatment is based on the principle that if a hole in a
piece is created, stresses around the hole and the hole
will be released. When this amount increases, the mobi-
lity of the hole will increase. The strain and thus the
amount of residual stresses can be calculated by deter-
mining the displacement of the hole. This procedure is
based on Tait’s research.3
The amounts of axial residual stresses were produced
by the drilling under post-weld heat treatment technique.
D. DANYALI, E. RANJBARNODEH: NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE EFFECT ...
776 Materiali in tehnologije / Materials and technology 50 (2016) 5, 775–782
Figure 1: Schematic of geometrical model for verification
Slika 1: Shematski prikaz geometrijskega modela za preverjanje
Figure 2: Thermo-physical and mechanical properties
Slika 2: Termofizikalne in mehanske lastnosti
The samples were drilled using a TOSKURIM drilling
machine.
The high-speed steel drill was 2 mm in diameter,
known as the drill chuck. The rotation speed was
2743 min–1 and the forward speed of the drill was 0.05
mm/min. After drilling, the distances between the holes
were measured and recorded with a MITUTOYO
micrometer type with an accuracy of up to 0.02 mm.
After measuring the distance between the holes, the
samples were heat treated according to the guidelines for
heat treatment.5
The applied heating rate was 220 °C/h. The welded
parts were held at 620 °C for an hour and then they were
allowed to cool with a cooling rate of about 275 °C/h,
cooled in furnace to 300 °C and after that in air. Figure 3
shows the applied heat-treatment cycle for the welded
parts in this study.6
After that the samples were cooled, the distances
between the holes were measured and recorded again in
the same manner using the same micrometer. The resi-
dual stresses in the specimens are released by drilling the
specimens, and the axial strain, ex, and transverse re-
leased strain, ey, are measured. Using the measured
strains, the axial residual stress, x, can be obtained from
Equation (1):7
x = [–E/(1 – v
2)](ex + vey) (1)
where E is Young’s modulus and v is Poisson’s ratio.
2.6 Hydrostatic test verification procedure
Before applying hydrostatic pressure, the distances
between holes were measured and recorded again. The
drilling apparatus and procedure were the same as used
in the calculation of the residual stresses. Figure 4
shows a hydrostatic test setup and Figure 5 shows the
sample drilled to investigate the behavior of residual
stresses after the hydrostatic pressure.
To create the real conditions, like hydrostatic test
conditions, the welded part was restrained between the
fixed jaws of a clamp. The fittings for the pressure-test-
ing machine were installed and water with a temperature
of about 15 °C was injected into the welded part. When
the air bubbles exit completely, a pressure equal to 1.5
times the boiler’s design pressure was applied to the
welded part.6 Due to the design the pressure of the
investigated boiler in this study was 0.4 MPa, while the
D. DANYALI, E. RANJBARNODEH: NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE EFFECT ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 775–782 777
Figure 3: Heat-treatment cycle
Slika 3: Potek toplotne obdelave
Figure 5: The sample drilled for axial residual stress measurement
and hydrostatic tests
Slika 5: Zvrtan vzorec za merjenje aksialnih zaostalih napetosti in
hidrostati~ni preizkus
Figure 4: Hydrostatic test setup
Slika 4: Sestav za hidrostati~ni preizkus
Figure 6: Applied hydrostatic pressure as a function of time
Slika 6: Uporabljen hidrostati~ni tlak v odvisnosti od ~asa
applied hydrostatic pressure to the welded part was about
0.6 MPa.
Figure 6 shows the applied hydrostatic pressure as a
function of time. After 1800 s, the hydrostatic pressure
was removed. After the completion of the hydrostatic
testing procedure, the distances between the drilling
locations were measured and recorded with the same
micrometer.
2.7 FEM meshed model
The FEM model of the welding process is performed
on the St35.8 steel tube with an outer diameter of 51 mm
and a wall thickness of 3 mm, and 17Mn4 steel tube-
sheet with a wall thickness of 12 mm using a single pass.
There was no gap between the tube and tube-sheet. The
welding process used was gas tungsten arc welding. For
a more convenient calculation, a part of the tube-sheet
was selected for the FEM analysis. Although the real
welding is a three-dimensional procedure, it is often
considered sufficient to represent a circumferential weld
with a two-dimensional axisymmetric FEM model.7 The
two-dimensional axisymmetric FEM model is much
faster and easier to perform.8,9 Therefore, the methodo-
logy described here is based on a two-dimensional
axisymmetric model.
The meshed model is shown in Figure 7. For the
meshing, plane55 and Surf151 elements were used for
thermal analysis and the plane82 element was used for
the mechanical analysis. In total, 346 nodes and 303
elements were generated. The welding of the tube to
tube-sheet was performed with a single pass weld. In the
weld adjacent, the meshing was very fine due to the high
thermal gradient of this region during the welding
process.
2.8 Thermal analysis
The heat equation is a parabolic partial differential
equation that describes the distribution of heat (or vari-
ations in temperature) in a given region over the time and
is generally describes by Fourier’s relation, Equation
(2):6
c
t x
K
T
x y
K
T
y z
Kx y zp
d
d
d
d
∂
∂
∂
∂
∂
∂
∂
∂
⎛
⎝
⎜ ⎞
⎠
⎟ = ⎡⎣⎢
⎤
⎦⎥+
⎡
⎣⎢
⎤
⎦⎥
+
d
d
T
z
Q
⎡
⎣⎢
⎤
⎦⎥+ (2)
where  is the density (kg/m3), Cp is the specific heat
capacity (J/kg K), K (W/m K) is the thermal conduc-
tivity, Q is the heat input (W), and T is the temperature
(K).
The Gaussian function is calculated by following
Equation (3):6,10
q r
Q
r
r
r
( ) exp=
⎛
⎝
⎜
⎞
⎠
⎟ −
⎛
⎝
⎜
⎞
⎠
⎟
⎡
⎣⎢
⎤
⎦⎥
3
3
0
2
0
2π
(3)
where q(r) is the heat flux (W/m2), r0 is the radius of the
welding heat source that is estimated to be about one
half of the Gaussian curve width, r is the distance from
the middle point of the heat source to the center, and Q
is the heat input (W) that is calculated using Equation
(4):6
Q = V·I· (4)
where V is the welding voltage, I is the welding current,
and ç is the welding efficiency. In this paper V, I,  and
r0 were 10 V, 120 A, 0.85 and 10 mm, respectively. The
filler metal, welding speed and gas flow rate were
ER70S-6, 5cm/min and 15 L/min, respectively.
The fluid flow in the welding pool increases the
heat-transfer rate, which has an important effect on the
temperature field in the welding process. The effect of
fluid flow on the weld pool temperature for the whole
temperature field is considered by the heat-transfer
coefficient, h. In this study, as in a previous study, h = 15
D. DANYALI, E. RANJBARNODEH: NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE EFFECT ...
778 Materiali in tehnologije / Materials and technology 50 (2016) 5, 775–782
Figure 8: The applied thermal boundary conditions
Slika 8: Uporabljeni toplotni robni pogoji
Figure 7: FEM meshed model
Slika 7: Model mre`e za FEM
W/K m2.11 Due to the small thickness of the parts,
preheating before welding is not done and the initial
temperature of the welded parts was assumed to be
27 °C. Figure 8 shows the applied thermal boundary
conditions.
2.9 Mechanical analysis
The residual stress was calculated by using the tem-
perature distribution obtained from the thermal analysis
as input data. The material properties relevant to the re-
sidual stress are the Young’s modulus, the yield strength,
the Poisson’s ratio, and the coefficient of thermal
expansion. The total strain can be decomposed into three
components, as in Equation (5):11
thotal = e + p + th (5)
where e, p, and t are the elastic strain, plastic strain
and thermal strain, respectively. The elastic strain was
modeled using the isotropic Hooke’s law with a tempe-
rature-dependent Young’s modulus and Poisson’s ratio.
During the mechanical analysis, boundary conditions
were applied to prevent the rigid-body motion. Figure 9
shows the applied mechanical boundary conditions.
Because of the tube-sheet rigidity, BC-X is constrained
in the X-direction and BC-Y is constrained in the
Y-direction. The S.L lines are axisymmetric lines.
3 RESULTS AND DISCUSSION
3.1 Thermal results
The thermal contour recorded by the experimental
method is shown in Figure 10 and the thermal results
recorded by experiment and the FEM are shown in
Figure 11. The weld-pool peak temperatures of the FEM
model and the experiments are 1900 °C and 1716 °C,
respectively. This 10 % difference shows the good agree-
ment between the experimental results and the FEM
model. This difference could be due to changes in
thermo-physical properties of the materials as a function
of the temperature, the difference in the thermal con-
ductivity. In the same way, Majzobi and colleagues con-
ducted a study in which a thermography technique was
used to record the thermal history.12 The thermal history
D. DANYALI, E. RANJBARNODEH: NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE EFFECT ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 775–782 779
Figure 10: Contours of the thermal history
Slika 10: Plastnice termi~ne zgodovine
Figure 9: The mechanical boundary conditions
Slika 9: Mehanski robni pogoji
Figure 12: Temperature contours for the FEM model during welding
Slika 12: Razporeditev temperature pri FEM modelu med varjenjem
Figure 11: Curves of the thermal history
Slika 11: Krivulji toplotne zgodovine
of the contours recorded by the other researches also
shows a difference of about 11 %, which confirms the
results of the present study.
The weld profile that is created with the FEM is
shown in Figure 12. The peak of the weld’s penetration
is 1.06 mm into the tube. The weld profile that is created
using the experimental method is shown in Figure 13.
The peak of the weld penetration is 1.2 mm into the tube.
Comparing these two sets of results shows that great-
est difference in the amount of weld penetration into the
tube-sheet is about 12 %. The results are proved to be in
a logical agreement for the FE method and experimental
results.
3.2 Mechanical behavior under hydrostatic effect
The residual stresses contour in the axial direction is
shown in Figure 14. The axial residual stress calculated
by the FEM and experimental methods is shown in Fig-
ure 15. The peak axial residual stress of 105 MPa is
located in the weld root zone. The peak axial residual
stress calculated by the experimental method was 85.94
MPa. The axial residual stresses first decrease from the
peak value, and then start to increase. All the axial resi-
dual stresses decrease along the weld and the outer
surface of the tube. The axial residual stress values are
D. DANYALI, E. RANJBARNODEH: NUMERICAL AND EXPERIMENTAL INVESTIGATION OF THE EFFECT ...
780 Materiali in tehnologije / Materials and technology 50 (2016) 5, 775–782
Figure 17: Residual stresses in axial direction using the FEM and
experimental methods, after the hydrostatic test
Slika 17: Zaostale napetosti v osni smeri po FEM in po eksperimen-
talni metodi, po hidrostati~nem preizkusu
Figure 14: Residual stresses results in axial direction using the FEM
method, before the hydrostatic test
Slika 14: Zaostale napetosti v osni smeri po FEM metodi, pred hidro-
stati~nim preizkusom
Figure 13: The weld profile using the experimental method
Slika 13: Profil zvara pri preizkusu
Figure 15: Residual stresses results in axial direction using the FEM
method, before the hydrostatic test
Slika 15: Zaostale napetosti v smeri osi, po FEM metodi, pred hidro-
stati~nim preizkusom
Figure 16: Residual stresses in axial direction using the FEM and
experimental methods, after the hydrostatic test
Slika 16: Zaostale napetosti v smeri osi po FEM in eksperimentalni
metodi, po hidrostati~nem preizkusu
reasonable, compared to Sattari-Far’s research re-
sults.13,14 The maximum difference between the two
methods is 18 % and the lowest is 2 %. As compared to
Tait’s results2, the findings of the present study are
acceptable. The peak axial tensile residual stress is 105
MPa in the root of the joint, which caused the initiation
and growth of fatigue cracks under cyclic loading in this
boiler joint.
In order to investigate the effect of the hydrostatic
test on the residual stress in the axial direction in a
boiler-tube weld, hydrostatic pressure is applied using
the FEM and experimental methods. Figure 16 shows
the curve of the residual stress results in the axial direc-
tion using the FEM and experimental methods, before
the hydrostatic test. It is clear that the hydrostatic pres-
sure affects the residual stress in the axial direction, as
shown in Figure 17. The maximum difference between
the two methods is about 9 %. The results of the FE
method show a peak in the residual stress in the axial
direction after the application of the hydrostatic pressure
is 52.4 MPa. The results of the experimental method
show that the peak residual stresses in the axial direction
after the application of hydrostatic pressure is 47.7 MPa.
The residual stresses in the axial direction calculated
by the FEM method in the weld root are shown in Fig-
ure 18. The residual stresses in the axial direction cal-
culated by the FEM method in weld root show that the
application of a hydrostatic test pressure after the
welding process can reduce the tensile residual stresses
in the axial direction from 105 MPa to of 52.4 MPa, i.e.,
by approximately 51 %. The results of the residual
stresses in the axial direction calculated using the FEM
method in the weld toe shows that hydrostatic pressure
can reduce the tensile residual stresses in the axial
direction from 40.8 MPa to 20.4 MPa, i.e., a reduction of
about 50 %. The residual stresses in the axial direction
calculated using the experimental method in the weld
shows that applying a hydrostatic test pressure after the
welding process can reduce the tensile residual stresses
in the axial direction from 20.74 MPa to 8.71 MPa. This
means a reduction of about 58 %. The reason for this
reduction can be sought in the accumulation of the
hydrostatic pressure and the previous residual stresses
that were caused by plastic deformation. In the plasti-
cally deformed region, the total stresses exceeded the
yield strength and partial relaxation of the residual
stresses can be occurred.
4 CONCLUSIONS
This study predicted the axial residual stresses and
their behavior under hydrostatic pressure using the FEM
and experimental methods. The results of this study can
be summarized as follows:
1. The root of tube to tube-sheet joint and the weld toe
were two critical points with the maximum axial
tensile residual stress.
2. Applying a hydrostatic pressure after the welding
process could reduce the axial tensile residual
stresses in the weld by about 58 %.
3. Applying a hydrostatic pressure after the welding
process could reduce the axial tensile residual
stresses in the root of tube to tube-sheet joint by
about 50 %.
4. Applying hydrostatic pressure after the welding
process could reduce the axial tensile residual
stresses in the weld toe by about 51 %.
5. It was easier and cheaper to calculate the residual
stresses in the design of a complex geometry in
connection with a corner joint with a very low
thickness using the technique of strain measurement
accuracy of the hole under the heat treatment.
6. The results of the present study could be generalized
for other critical boiler joints, such as the shell to
furnace joint and the furnaces to tube-sheet joint.
7. The results of the present study could be generalized
to other high-pressure equipment including heat
exchangers, hot-water boilers, hot-oil boilers and any
other equipment that has a tube to tube-sheet joint.
5 REFERENCES
1 X. Shugen, W. Weiqiang, Numerical investigation on weld residual
stresses in tube to tube sheet joint of a heat exchanger, Int. J. Pres.
Ves. Pip., 101 (2013), 37–44, doi:10.1016/j.ijpvp.2012.10.004
2 R. Tait, J. Press, Investigation An experimental study of the residual
stresses, and their alleviation, in tube to tube-sheet welds of
industrial boilers, Engineering Failure Analysis, 8 (2001), 15–27,
doi:10.1016/j.engfailanal.2013.04.01
3 K. Abouswa, F. Elshawesh, A. Abuargoub, Stress corrosion cracking
(caustic embrittlement) of super heater tubes, Desalination, 222
(2008), 682-688, doi:10.1016/j.desal.2007.02.073
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Materiali in tehnologije / Materials and technology 50 (2016) 5, 775–782 781
Figure 18: Residual stresses in the axial direction using the FEM and
experimental methods, before and after the hydrostatic test
Slika 18: Zaostale napetosti v osni smeri po FEM in po eksperimen-
talni metodi, pred in po hidrostati~nem preizkusu
4 B. Yildrim, H. F. Nied, Residual Stress and Distortion in Boiler Tube
Panels with Welded Overlay Cladding, Journal of Pressure Vessel
Technology – Transactions of ASME, 126 (2004), 426–431,
doi:10.2320/matertrans.126.426
5 X. Qingren, F. Yaorong, H. Chun-Yong, The measurement and
control of residual Stress in spiral submerged arc welded pipe,
Proceedings of 4th International pipeline conference, Calgary, 2002,
615–622
6 British Standards Institute, BS2790, Design and manufacturing of
shell boilers of welded construction, 1992, 80–100
7 C. Lee, K. Chang, Numerical investigation of residual stresses in
Strength mismatched dissimilar steel butt welds, J. Strain Analysis,
43 (2008), 55–66, doi:10.1243/03093247JSA313
8 T. Soanes, W. Bell, A. Vibert, Optimizing residual stresses at a repair
in a steam header to tube plate weld, Int. J. Pres. Ves. Pip., 82 (2005),
311–318, doi:10.1016/j.ijpvp.2004.08.009
9 D. Deng, H. Murakawa, Numerical simulation of temperature field
and residual stress in multi-pass welds in stainless steel pipe and
comparison with experimental measurements, Comp. Mater. Sci., 37
(2006), 269–277, doi:10.1016/j.commatsci.2005.07.007
10 X. Shugen, Z. Yanling, Using FEM to determine the thermo
mechanical stress in tube to tube–sheet joint for the SCC failure
analysis, Engineering Failure Analysis, 34 (2013), 24–34,
doi:10.1016/j.engfailanal.2013.07.01
11 T. Tso-Liang, F. Chin-Ping, C. Peng-Hsiang, Y. Wei-Chun, Analysis
of residual stresses and distortion in T-joint fillet welds, Int. J.
Pres.Ves. Pip., 78 (2001), 523–538, doi:10.1016/j.ijpvp.2000.07.008
12 J. Otegui, P. Fazzini, Failure analysis of tube–tubesheet welds in
cracked gas heat exchangers, Engineering Failure Analysis, 11
(2004), 903–913, doi:10.1016/j.engfailanal.2004.01.003
13 I. Sattari-Far, Y. Javadi, Influences of welding sequence on welding
distortion in pipes, Int. J. Pres. Ves. Pip., 85 (2008), 265–274,
doi:10.1016/j.engfailanal.2007.02.004
14 I. Sattari-Far, M. Farahani, Effect of the weld groove shape and pass
number on residual stresses in butt-welded pipes, Int. J. Pres. Ves.
Pip., 86 (2009), 769–777, doi:10.1016/j.ijpvp.2009.07.007
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782 Materiali in tehnologije / Materials and technology 50 (2016) 5, 775–782
P. SKUBISZ et al.: EFFECT OF DIRECT COOLING CONDITIONS ON THE MICROSTRUCTURE ...
783–789
EFFECT OF DIRECT COOLING CONDITIONS ON THE
MICROSTRUCTURE AND PROPERTIES OF HOT-FORGED HSLA
STEELS FOR MINING APPLICATIONS
VPLIV POGOJEV OHLAJANJA NA MIKROSTRUKTURO IN
LASTNOSTI VRO^E KOVANIH HSLA JEKEL ZA UPORABO V
RUDARSTVU
Piotr Skubisz, £ukasz Lisiecki, Tadeusz Skowronek, Artur ¯ak, W³adys³aw Zalecki
1AGH University of Science and Technology, Department of Metals Engineering and Industrial Computers Science, 30A Mickiewicz Ave,
Kracow, Poland
2Ferrous Metals Institute, 12-14 K. Miarki, Gliwice, Poland
pskubisz@metal.agh.edu.pl
Prejem rokopisa – received: 2015-09-18; sprejem za objavo – accepted for publication: 2015-10-26
doi:10.17222/mit.2015.298
The article presents hot deformation and controlled direct cooling of hardened medium-carbon HSLA steel with micro-additions
of alloying elements Ti and/or V. It also contains a study of the effect of thermomechanical-processing conditions on grain
refinement and precipitation kinetics in the replacement of the conventional reheating-requiring heat treatment with a
cost-effective technology. Controlled cooling with accelerated air and mist adjusted to lower carbon and hardenability-related
alloying elements by means of employing experimentally designed heats, varying in the Mo content, was designed to meet the
mining-industry requirements for mechanical properties. Besides microstructural-property relations, the vital problem of
non-uniformity of the obtained properties is addressed with regard to within-part variances in the cooling rate of the as-forged,
undeformed and dynamically recrystallized material. The strength and plasticity versus the microstructure of the produced fine
pearlite/bainite structure with grain-boundary ferrite were evaluated. Microstructure-property relations allowed the formulation
of the conclusions on the effect of direct-cooling conditions on the microstructure and grain substructure, their mutual synergic
effect in controlling the microstructure, as well as guidelines for the transfer of the locally established process parameters into
technological conditions.
Keywords: thermomechanical processing, drop forging, accelerated cooling, grain refinement, tempered martensite
^lanek predstavlja vro~o deformacijo in kontrolirano ohlajanje srednjeoglji~nega HSLA jekla z mikrododatkom legirnih
elementov Ti in/ali V. Vsebuje tudi {tudijo vpliva pogojev termomehanske obdelave na udrobnjenje zrn in kinetiko izlo~anja pri
nadome{~anju obi~ajnega re`ima ogrevanja za toplotno obdelavo s cenej{o tehnologijo. Da bi dosegli zahteve rudarske
industrije in mikrostrukturne lastnosti, je bilo vzpostavljeno kontrolirano ohlajanje s tokom zraka in me{anice zraka ter vode,
prilagojeno ni`jemu ogljiku in drugim elementom. Ti vplivajo na prekaljivost z uporabo eksperimentalnih talin, v katerih se je
spreminjala vsebnost Mo. Poleg odvisnosti lastnosti od mikrostrukture, je glavni problem neenakost dobljenih lastnosti, kar je
delno povezano z razlikami pri ohlajanju v kosu kovanega, nedeformiranega ali dinami~no rekristaliziranega materiala.
Ocenjeni sta bili trdnost in plasti~nost v odvisnosti od nastale drobnozrnate perlitno/bainitne mikrostrukture s feritom po mejah
zrn. Odvisnost mikrostrukture od lastnosti omogo~a postavitev zaklju~kov o vplivu pogojev pri neposrednem ohlajanju na
mikrostrukturo in zrna substrukture ter njihov medsebojni vpliv pri kontroli mikrostrukture, kot tudi navodila za prenos lokalno
ugotovljenih procesnih parametrov v tehnolo{ke pogoje.
Klju~ne besede: termomehanska obdelava, kovanje s padalnim kladivom, pospe{eno ohlajanje, zmanj{anje zrn, popu{~en
martenzit
1 INTRODUCTION
Forging has long been more than a mere shaping
technique. In addition to the accuracy requirements, such
as allowances and yield, mechanical properties must be
fulfilled. As an alternative to the traditional quenching
and tempering (Q&T) heat treatment, thermomechanical
processing (TMP) in controlled conditions of forging and
direct cooling is increasingly used on an industrial scale,
offering a good combination of strength and ductility at a
lower cost.1–3
The strength levels obtained with TMP are often
lower as compared to a traditional Q&T material4, due to
an excessive surplus of indices in the case of the latter.
Furthermore, some mechanical properties of TMP mi-
croalloyed grades, such as the fatigue strength, are
second to those treated with Q&T.5 The key issue is to
control the chemistry and processing cycle so as to
achieve the sustainable strength, plasticity, cracking
and/or fatigue resistance required by the end user of a
finished product.
Thus, numerous microalloyed grades for specific
applications have been designed over the decades. So
have the technologies of forging and subsequent cool-
ing.6,7 This work is devoted to the design or modification
of a microalloyed steel grade exhibiting as-forged air-
cooled mechanical properties suitable for a mining appli-
cation. The target application of this combination is a
miners’ linking hook for coal transport, which, after a
Materiali in tehnologije / Materials and technology 50 (2016) 5, 783–789 783
UDK 621.78.08:67.017:691.714 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)783(2016)
controlled hot-forging and direct heat treatment, involv-
ing quenching and self-tempering, is expected to provide
a yield strength (YS) of 800 MPa, the ultimate tensile
strength (UTS) of 1050 MPa, an elongation to fracture of
A 15 % and/or an impact strength at room temperature of
KCV 60 J/cm2 to substitute C45 or SJ355 heat-treatable
structural steel grades.
2 EXPERIMENTAL METHODS
The research plan included testing the effect of forg-
ing temperature on the microstructure and mechanical
properties obtained after direct cooling in forced air and
atomizers, applied to designed microalloyed grades, hot
deformed in an inconvenient high-speed hammer-forging
process. The goal of the study was both the investigation
of the effect of material-process conditions on the final
(as-forged) properties, and its applicability to drop-forg-
ing industrial conditions with respect to the uniformity of
the properties.
Eight different sets of alloying elements were com-
posed on the basis of the computation in Thermocalc
under equilibrium thermodynamics conditions. Diagrams
were established on the basis of a volumetric-change
analysis (dilatometer DIL 805 with probe LVDT) CTT,
followed by an examination of the microstructures ob-
tained with different cooling rates of 0.17 °C/s,
0.42 °C/s, 1 °C/s, 2 °C/s, 4 °C/s, 10 °C/s, 20 °C/s,
40 °C/s, 50 °C/s and 100 °C/s, taking into consideration
the cooling rates predicted in the modeling of physical
cooling of the target geometry. Having had selected two
grades, further described as A and B, laboratory tests of
rolling and direct cooling with water and air were con-
ducted, as well as forging on a hydraulic press of 5 MN
with a ram velocity of 50 mm/s.
To estimate the amount of deformation and the actual
forge-end temperature in the bulk of the part, a nume-
rical calculation of an equivalent strain and temperature
progression in hot forging was carried out. Since the
forging temperature should be high enough to dissolve
carbides and carbonitrides so as to enable their precipi-
tation during the post-forging direct cooling, a soaking
temperature of 1190 °C was established and forging
temperatures ranging between (1100, 1000 and 900) °C
were assumed. The cooling rate – fast enough to provide
harder microstructural components during the sub-
sequent cooling8 – was realized with atomizers.
Numerical modeling was conducted with the finite-
element method (FEM) in code QForm3D. Besides
supplying the data for the prediction of transformation
products or microstructural development, and carbide
and carbonitride precipitation kinetics, the results of the
calculation formed the basis for experimental controlled
cooling tests on a laboratory continuous cooling line,
QuenchTube Duo (Figure 1).
The laboratory cooling line simulated an industrial
cooling line, providing a forced stream of mist with
intervals resulting from the transfer of a part between
consecutive cooling zones. All runs involved the same
cycle, irrespective of the forging temperature and the
alloy. Cooling was conducted at a cooling rate adapted to
the temperature changes, calculated for a drop-forged
miners’ link hook (Figure 2), which formed a case study
for an evaluation of the effectiveness of the selected pro-
cessing conditions with respect to the forging tempera-
ture and chemical composition.
3 MATERIALS
The study involved experimentally designed steel
grades, which were to provide directly cooled forgings
with the mechanical properties comparable to typical
mining-industry grades, such as C45, 41Cr, 36CrNiMo4
and 23MnNiCrMo5.
Eight grades were made, based on the assumption of
a reduction of carbon and the major alloying elements,
and micro-additions of Ti, V and Nb, with or without
Mo. The heating was performed in an induction-heating
vacuum furnace, VSG-100, with a 100 kg nominal capa-
city of the melting pot. Thermodynamic stability and
volume fractions of phase components and precipitates
were calculated in Thermocalc and confirmed with a
dilatometric analysis. Upon the investigation of the pri-
mary and transformed austenite grain size, the predicted
or calculated yield and impact strength, two grades were
P. SKUBISZ et al.: EFFECT OF DIRECT COOLING CONDITIONS ON THE MICROSTRUCTURE ...
784 Materiali in tehnologije / Materials and technology 50 (2016) 5, 783–789
Figure 2: Geometry of the target application; T1, T2 – thermocouple
locations
Slika 2: Geometrija preizku{anca; T1, T2 – lokacija termoelementovFigure 1: Laboratory simulation of the cooling line
Slika 1: Laboratorijski simulator linije ohlajanja
selected (Table 1), indicating the smallest grain size and
a potential combination of strength and ductility.
Utilizing the deformed material from plastometric tests
with formulas for characteristic times,9 continuous-cool-
ing-transformation diagrams (CCT) of non-recrystallized
materials were made, as shown in Figure 3.
4 RESULTS AND EVALUATION
4.1 Analysis of the target forging technology
Forging-machine kinematics has an enormous impact
on thermomechanical parameters for a hot-forged piece.
As could be expected, hammer forging produces an
increased amount of generated heat and temperature
gradients, which influence the flow stress and metal dis-
placement. Numerical modeling with FEM allowed for
an estimate of the level, variance and gradients of the
temperature and strain development during the forging.
This information made it possible to evaluate the ob-
tained results of direct cooling where, unlike for the
traditionally heat-treated material, the as-forged steel
inherits the aftermath of dynamic processes. Thus, to fit
the laboratory-test results to the target application, we
carried out an insightful analysis of the development of
the thermomechanical indicators of the changes influenc-
ing the final microstructure and properties.
The results of the FEM analysis are summarized in
Table 2.
Table 2: Evolution of thermo-mechanical conditions in the forged part
(in locations shown in Figure 2)
Tabela 2: Razvoj termo-mehanskih pogojev v kovanem kosu (na me-
stih prikazanih na Sliki 2)
Operation 1 2 3 4 5
Effective strain
T1 0.62 0.62 0.62 0.84 0.88
T2 0.62 1.25 2.51 2.75 2.81
Temp., °C
T1 1115 1107 1089 1105 1090
T2 1115 1130 1136 1144 1140
4.2 Effect of the cooling rate
To assess the dependence of the structural compo-
nents and the grain size on the cooling rate, the micro-
structure derived from a dilatometric investigation of
normalized material was examined. Micrographs ob-
tained for both alloys are shown in Figures 4 and 5.
4.3 Physical modeling of forging and cooling
The effect of the run-out table temperature on the ki-
netics of austenite restoration and precipitation of car-
bides and carbonitrides, and the resulting strengthening
P. SKUBISZ et al.: EFFECT OF DIRECT COOLING CONDITIONS ON THE MICROSTRUCTURE ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 783–789 785
Figure 4: Microstructure of as-received (hot-rolled) alloy A after anisothermal cooling at rates: a) 1 °C/s, b) 10 °C/s, c) 50 °C/s
Slika 4: Mikrostruktura izhodne (vro~e valjane) zlitine A, po anizotermnem ohlajanju s hitrostjo: a) 1 °C/s, b) 10 °C/s, c) 50 °C/s
Figure 3: Continuous-cooling diagrams calculated with TTSteel code,
modified with theoretical equations9
Slika 3: Diagram kontinuiranega ohlajanja, izra~unan s TTSteel kodo
in spremenjen s teoreti~nimi ena~bami9
Table 1: Chemical compositions of experimental heats of microalloyed steels used in the study
Tabela 1: Kemijska sestava eksperimentalnih talin mikrolegiranih jekel, uporabljenih v {tudiji
Steel % C % Mn % Cr % Si % Mo % Ti % V % Nb % N Ac1,°C Ac3,°C
A 0.30 1.50 0.42 0.26 0 0.011 0.09 0.039 0.011 809 785
B 0.28 1.24 0.42 0.27 0.2 0.019 0.067 0.047 0.010 728 727
efficiency was investigated by varying the forging tem-
perature while employing the same cooling conditions
after the deformation.
According to the plots of temperature measured
during the experiments, the assumed temperatures
changed during the deformation and therefore the forge
end temperature was not exactly the same as expected.
Due to a non-uniform distribution of strain in the bulk,
the amount of generated deformation heat varied with
the location (Table 2). However, the assumed tempera-
tures of forging, (1100, 1000 and 900) °C, allowed an
analysis of the material response in the representative
ranges in relation to the temperatures of the dynamic-
recrystallization stop and the solution of carbonitrides/
carbides. The highest temperature, assumed for the total
content of carbon and microalloying elements, was
found in the solution, available for precipitation. The
lowest forging temperature of 900 °C dropped to 880 °C
due to the material transfer and die cooling; it was meant
to provide the alloys with a reasonable amount of accu-
mulated strain, based on the calculated recrystallization
stop temperature. The forging at 1000 °C was to show
the effect of the precipitates formed prior to forging on
the evolution of dynamically recrystallized grains. As
alloy B did not contain Mo, this test was conducted for
alloy A only.
The obtained cooling curves are shown in Figure 6
for steels A and B, respectively. For clarity, the plots
were shifted and the moment of deformation was
indicated by plotting the forging load (in grey). This was
also an opportunity to indicate unexpected dependence
of the load extreme on the temperature for steel B, which
shows a higher value at 1000 °C than at 900 °C. The
experiment conducted showed that the deformation
temperature is high enough to avoid ferrite nucleation,
which adversely affects the strength properties.10,11 The
load observed at 900 °C implies there is no occurrence of
ferrite dynamic precipitation, while incomplete softening
can be found with the aid of a metallographic analysis
(Figures 7 and 8).
5 MICROSTRUCTURE AND MECHANICAL
PROPERTIES
The microstructure obtained using accelerated cool-
ing with mist is composed of martensite and bainite,
depending on forging temperature portions of fine and
narrow-spaced pearlite, differing in the fractions of the
constituents. The cooling rate of 20 °C/s is sufficient to
omit the high-temperature onset of a diffusion-driven
transformation into recrystallized austenite. Industrial-
like conditions of direct cooling (based on the assump-
tion of quenching directly after deformation) caused an
increased amount of ferrite, with a simultaneous increase
in the fraction of martensite instead of bainite or
Widmanstätten ferrite. As indicated in the physical
simulation included in the study, 80° s–1 in a range of
P. SKUBISZ et al.: EFFECT OF DIRECT COOLING CONDITIONS ON THE MICROSTRUCTURE ...
786 Materiali in tehnologije / Materials and technology 50 (2016) 5, 783–789
Figure 5: Microstructure of as-received (hot-rolled) alloy B after anisothermal cooling at rates: a) 1 °C/s, b) 10 °C/s, c) 50 °C/s
Slika 5: Mikrostruktura izhodne (vro~e valjane) zlitine B, po anizotermnem ohlajanju s hitrostjo: a) 1 °C/s, b) 10 °C/s, c) 50 °C/s
Figure 6: Experimental cooling-curve plots for: a) steel A, and
b) steel B; load peaks (grey) indicate forge-end points
Slika 6: Diagrami eksperimentalnih krivulj ohlajanja: a) jeklo A in
b) jeklo B; konice obremenitve (siva barva) ka`ejo konec kovanja
800–500 °C suffices to produce a 50 % martensite
transformation after the post-forging operations like
trimming the flash and transfer on air. As shown in
Figure 7, by forging at 900 °C, we obtained non-
recrystallized microstructures with a high density of
crystallographic defects, such as shearing bands and
sub-cells with very fine grains at the grain boundaries,
illustrating the increment of the total strength enhance-
ment due to the reduction of the forging regime.
Based on comparable cooling conditions, the micro-
structure analysis indicated a strong effect of forging
conditions on the grain size and morphology, which is
reflected by mechanical properties. Steel A exhibits a
higher strength, with UTSs of 1688 MPa after the forging
at 1100 °C and 1486 MPa for 900 °C (Table 3). Steel B
shows a better ductility associated with the Mo addition.
Furthermore, directly cooled steel A showed a 5 %
elongation while steel B indicted an 8 % elongation,
which is a satisfactory result, taking into account that no
tempering was carried out, while auto-tempering, enabl-
ing a ductility enhancement, could be utilized in indu-
strial processes.
6 CONCLUSIONS
Deformation at a varied temperature, in connection
with direct cooling, indicates a strong influence of
hammer-forging conditions on the resultant mechanical
properties. This gave us an opportunity for investigating
P. SKUBISZ et al.: EFFECT OF DIRECT COOLING CONDITIONS ON THE MICROSTRUCTURE ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 783–789 787
Figure 8: Microstructures of alloy B in as-forged direct-cooling con-
ditions, after forging at: a) 900 °C, b) 1000 °C, c) 1100 °C
Slika 8: Mikrostruktura zlitine B, ohlajene takoj po kovanju na:
a) 900 °C, b) 1000 °C, c) 1100 °C
Figure 7: Microstructures of alloy A in as-forged direct-cooling con-
ditions, after forging at: a) 1100 °C, b) 900 °C
Slika 7: Mikrostruktura zlitine A, ohlajene takoj po kovanju:
a) 1100 °C, b) 900 °C
Table 3: Mechanical properties of the microalloyed steels after accele-
rated cooling
Tabela 3: Mehanske lastnosti mikrolegiranih jekel po pospe{enem
ohlajanju
Alloy
Forging
tempera-
ture, °C
YS,
MPa
UTS,
MPa
Elongation at
fracture, A10
%
Area
reduction,
%
Steel A
1100 °C 1104 1688 3.2 6
900 °C 915 1486 5.8 20
Steel B
1100 °C 706 1208 9.4 44
1000 °C 1043 1514 7.5 38
900 °C 757 1012 11.0 59
the possibilities of controlling microstructural properties
and comparing them.
According to the prevailing studies of the TMP utili-
zation in combination with microalloyed steels, detri-
mental hammer-forging conditions and the required final
properties make big differences. Nevertheless, hammer-
forging plants, as well as those employing high-speed-
press forging, are quite numerous and the control and
design of the TMP process in such conditions are still
challenging.
Both the chemical compositions of the designed
alloys and the technology based on the concept of direct
cooling from the forge-end condition assure simplicity
and attractiveness to the forging industry, offering
reduced costs and making it possible for steelworks to
keep the imposed limits of alloying elements. Moreover,
the predefined cooling conditions are easily transferable
to forge-plant conditions. The contents of the micro-
alloying elements in alloys are kept within reasonable
limits required from the standpoint of control of the
grain-structure evolution due to the interactions of the
precipitates with microstructural defects during forging
and the subsequent cooling.12
The incremental character of forging, the thin end,
optionally allows continuous-cooling deformation
modeling. However, it may be noticed that the hook
undergoes a double reduction: during the flattening and
finishing of the impression. Hence, a simple compression
of a 25 mm flat bar ensures appropriate geometrical and,
more importantly, phenomenological similarity condi-
tions. Thus, the forging tests reflected the industrial pro-
cesses of both forging and cooling. Having passed the
800–500 °C range, the cooling rate was slowed down
producing a type of the equalizing hold (Figures 6 and 7),
reducing thermal stresses on the one hand and the time
for the Nb carbide precipitation13 on the other hand.
Precipitates greatly contribute to the hardening of an
alloy; however, the efficiency of their strengthening de-
pends on the volume fraction and size of the precipitates.
Depending on the particle relation to the dislocation
parameters, it can pile up and loop the dislocations or be
sheared by it.14 In the case of the analyzed alloys, both
instances were witnessed. The produced carbides and
carbonitrides TiC and Ti(C,N) are of the largest size
among the occurring precipitates. They are reported to
take on the form of cube-like shapes, reaching a micro-
meter in diameter, which enabled an observation with the
immersion technique using an optical microscope. On
the other hand, Nb carbides have smaller particle dia-
meters. They are found to have a spherical shape and the
size of dozens of nanometers. Their contribution to the
total strength enhancement reaches 90 MPa13,15, which
causes high effectiveness in pinning grain boundaries
during deformation and grain-restoration processes.
The volume of these precipitates allow for a signifi-
cant microstructural controllability brought about by
Nb(C,V), provided the whole Nb content dissolves at
temperatures of 1170–1190 °C. The applied Nb content
is supposed to retard recrystallization, which calls for a
lower forging temperature.
The overall level of the strength properties results
from the grain refinement, which contributes up to
250 MPa.14 It is enhanced by the presence of N and Al,
lowering the tendency for coagulation of the V(C,N)
precipitates. In addition to reducing the amount of vana-
dium dissolved in the austenite16 by decreasing the self-
diffusion of iron, N reduces the grain-growth tendency17,
as long as it is below the content necessary for the VN
formation.18
Besides the precipitates, the grain is refined due to
the grain-structure restoration during the forging, which
greatly influences the concentration of nucleation sites.
Decreasing the forging temperature from 1180 °C to
1000 °C resulted in lowering the forge-end point from
about 1210 °C to 1034 °C, producing a fine-grained
structure, which formed a base for fine colonies of pear-
lite or bainite during the cooling. The produced micro-
structure exhibits a grain size decreasing with the lower-
ing forging temperature. However, contrary to higher
temperature trials, a microstructure abundant in cry-
stallographic defects was obtained during the forging at
about 900 °C, such as shearing bands and non-recry-
stallized sub-cells with particularly fine grains at the
grain boundaries, which can be attributed to disconti-
nuous dynamic recrystallization, resulting in a selective
renovation of the grains, typical of a low-temperature
deformation.19,20 The presence of such grains confirms
that the deformation was completed below the tempera-
ture, at which the recrystallization and/or the strain-
induced precipitation inhibiting static recrystallization
can occur. 21
The mixture of fine recrystallized and strain-hard-
ened grains produced a significant strengthening, re-
sulting in an UTS of 1700 MPa, similar to the related
studies.22 On the other hand, a low plasticity, reaching at
most 11 % of A10 elongation to fracture was obtained.
However, in an industrial practice, higher plasticity indi-
ces should be expected as, in contrast to small laboratory
samples, massive parts allow the auto-tempering effect to
occur, enabling a ductility improvement. It must be noted
that the flat specimens did not have a privileged
grain-flow direction. On the contrary, the grain-flow
orientation in the gauge area of the tensile specimens
was perpendicular to the tension direction. In the
considered part, the situation was different – the metal
flow pattern produced an evident longitudinal grain flow,
allowing the ductility enhancement.
Acknowledgements
Financial assistance of NCBiR within project PBS2/
B5/29/2013 agreement 19.19.110.86730, is acknowl-
edged.
P. SKUBISZ et al.: EFFECT OF DIRECT COOLING CONDITIONS ON THE MICROSTRUCTURE ...
788 Materiali in tehnologije / Materials and technology 50 (2016) 5, 783–789
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P. SKUBISZ et al.: EFFECT OF DIRECT COOLING CONDITIONS ON THE MICROSTRUCTURE ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 783–789 789

I. DINAHARAN, E. T. AKINLABI: INFLUENCE OF THE TOOL ROTATIONAL SPEED ON THE MICROSTRUCTURE ...
791–796
INFLUENCE OF THE TOOL ROTATIONAL SPEED ON THE
MICROSTRUCTURE AND JOINT STRENGTH OF FRICTION-STIR
SPOT-WELDED PURE COPPER
VPLIV HITROSTI VRTENJA ORODJA NA MIKROSTRUKTURO IN
TRDNOST TORNO VRTILNO TO^KASTO ZVARJENEGA SPOJA
^ISTEGA BAKRA
Isaac Dinaharan, Esther T. Akinlabi
University of Johannesburg, Department of Mechanical Engineering Science, Auckland Park, Kingsway Campus,
Johannesburg 2006, South Africa
dinaweld2009@gmail.com
Prejem rokopisa – received: 2015-09-21; sprejem za objavo – accepted for publication: 2015-10-05
doi:10.17222/mit.2015.301
Copper is very difficult to be spot welded with conventional fusion-welding techniques due to a high thermal diffusivity. Fric-
tion-stir spot welding (FSSW) is a novel solid-state welding process, suitable and effective for spot welding copper.
Commercially pure copper sheets of 3 mm thickness were spot welded using an industrial friction-stir welding machine. The
spot welds were made by varying the tool rotational speed at three levels. The spot welds were characterized using light
microscopy. The shear-fracture load was evaluated using a computerized tensile-testing machine. The results revealed that the
tool rotational speed remarkably influenced the microstructure, the shear-fracture load and the mode of fracture.
Keywords: copper, friction-stir spot welding, microstructure, shear load
Zaradi velike toplotne prevodnosti se baker zelo te`ko to~kasto vari pri obi~ajnih postopkih varjenja z zlivanjem. Torno vrtilno
to~kasto varjenje (FSSW) je nov na~in varjenja v trdnem stanju, ki je primerno in primerljivo s torno vrtilnim to~kastim
varjenjem bakra. Bakrene plo~evine, debeline 3 mm, iz ~istega komercialno dostopnega bakra, so bile to~kasto zvarjene s torno
vrtilnim to~kastim varjenjem, z uporabo industrijske naprave za tovrstno varjenje. To~kasti zvari so bili izdelani pri treh
hitrostih vrtenja orodja. Karakterizirani so bili z uporabo svetlobne mikroskopije. Stri`na trdnost je bila ocenjena z uporabo
ra~unalni{ko vodenega nateznega stroja. Rezultati so pokazali, da hitrost vrtenja orodja mo~no vpliva na mikrostrukturo, stri`no
trdnost in na~in preloma.
Klju~ne besede: baker, torno vrtilno to~kasto varjenje, mikrostruktura, stri`na obremenitev
1 INTRODUCTION
Pure copper is extensively used in the optical and
electronic industries owing to its excellent properties
such as good ductility, high electrical, thermal conduc-
tivity and good corrosion resistance. The welding of
copper is often encountered in the electrical, nuclear and
automobile industries.1,2 Spot welding of pure copper is
generally hard to do with conventional fusion welding
because of the high thermal diffusivity, which is about 10
to 100 times higher than in many steels and nickel alloys.
The heat input required is much higher than in almost
any other material. Further, pure copper is susceptible to
solidification cracking and blowhole formation.3,4 Fric-
tion-stir spot welding (FSSW) is a novel solid-state
welding technique, promising for spot welding copper
without the problems associated with fusion-welding
techniques.5
FSSW is a derivative process based on friction-stir
welding (FSW) which was developed at The Welding
Institute in 1991. There are several differences between
FSSW and FSW. One distinct difference is that there is
no translation of tool during FSSW. FSW is commonly
employed to join metallic plates in a butt configuration
along the line of contact. On the other hand, FSSW is
performed on thinner plates kept in a lap configuration.
A rotating, non-consumable, cylindrical-shouldered tool
with a pin is plunged, at a predetermined feed rate, into
the overlapping plates to a depth slightly shorter than the
total thickness of both plates. Frictional heat is generated
between the plate material and the rotating tool, which
plasticizes the material. The rotating action of the pin
induces the material flow in both the circumferential and
axial directions. The axial force applied along the tool
axis forges the plasticized material and forms an annular,
solid-state bond around the pin. At this moment, the
rotating tool is retracted, leaving the exit hole behind.
The major process parameters, which influence the joint
strength, are the tool geometry, the tool rotational speed,
the tool penetration depth and the dwell time. FSSW
exhibits key advantages such as excellent mechanical
properties, a low distortion, ease of handling, low cost,
and clean working environment.6–8
The FSSW technique has been successfully used to
spot weld aluminum9, magnesium10, steel11 and plastics.12
Both similar and dissimilar spot welds were reported in
Materiali in tehnologije / Materials and technology 50 (2016) 5, 791–796 791
UDK 621.663:67.017:621.791.05:669.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)791(2016)
the literature13,14. FSSW of copper and its alloys is least
explored. Z. Barlas15 spot welded pure copper and brass
sheets using FSSW and analyzed the influences of the
tool rotational speed, dwell period and material location
on the microstructure and the tensile shear-fracture load
(TSFL). He found that the welding parameters had an
important influence on the TSFL and the failure mode.
The TSFL increased with an increase in the tool rota-
tional speed and dwell time. R. Heideman16 spot-welded
AA6061 and pure-copper sheets and studied the effects
of the tool pin length, the shoulder plunge depth, the
welding time and the tool rotational speed on the TSFL.
Literature on FSSW of pure copper is scantily. There-
fore, the present work focuses on spot welding copper
sheets of 3 mm using FSSW and analyzes the influence
of the tool rotational speed on the TSFL.
2 MATERIALS AND METHODS
Commercially available pure copper sheets of 3 mm
were used in this study. The optical photomicrograph of
the as-received copper sheet is shown in Figure 1. Lap
joint configuration was used to fabricate the spot welds
where the rolling direction of the material was kept
parallel to the loading directions and the joint was
initially obtained by securing the sheets in position using
mechanical clamps. A non-consumable tool made of
high-carbon steel was used to fabricate the joints. The
tool had a shoulder of 18 mm and a conical profile of a
diameter varying from 5 mm to 3 mm along the length of
pin. The pin length was 5.7 mm. An industrial-purpose
FSW machine (I-STIR), depicted in Figure 2, was used
for FSSW. The FSSW procedure and the welding cycle
are depicted in Figure 3. The tool rotational speed was
varied at three levels of 1200 min–1, 1600 min–1 and 2000
min–1. Other parameters were kept constant. The dwell
period and axial force were 3 s and 10 kN, respectively.
Spot welds were made on two sheets clamped in the lap
configuration. Sufficient cooling was allowed between
successive spot welds. The spot-welded copper sheet is
shown in Figure 4. Six spot welds were made for each
set of the process parameters. Specimens were machined
from the welded sheets and polished semi-automatically
in a polishing machine (Struers LaboPol 25). The
mirror-polished specimens were etched with a color
etchant containing 20 g of chromic acid, 2 g of sodium
I. DINAHARAN, E. T. AKINLABI: INFLUENCE OF THE TOOL ROTATIONAL SPEED ON THE MICROSTRUCTURE ...
792 Materiali in tehnologije / Materials and technology 50 (2016) 5, 791–796
Figure 3: Schematic illustration of the friction-stir spot-welding
process
Slika 3: Shematski prikaz postopka torno vrtilnega to~kastega varjenja
Figure 1: Light micrograph of pure copper
Slika 1: Svetlobni posnetek mikrostrukture ~istega bakra
Figure 2: I-STIR friction-stir welding machine
Slika 2: I-STIR naprava za torno vrtilno varjenje
Figure 4: Friction-stir spot-welded sheet
Slika 4: Torno vrtilno to~kasto zvarjena plo~evina
sulfate and 1.7 mL of HCl (35 %) in 100 mL distilled
water. The macrostructure was recorded using a stereo
microscope (OLYMPUS SZX16). The microstructure
was observed using a metallurgical microscope (OLYM-
PUS BX51M). The TSFL was evaluated using a compu-
terized tensile tester (INSTRON 1195) at a crosshead
speed of 2 mm/min.
3 RESULTS AND DISCUSSION
Macrographs of the copper friction-stir spot-welded
as a function of the tool rotational speed are presented in
Figure 5. It is possible to spot weld successfully using
the chosen parameters. The weld zones are almost sym-
metrical with respect to the axis of the keyhole. The tool
rotation generates frictional heat, which plasticizes the
copper. The force applied along the axis of the tool
promotes the vertical motion of the plasticized copper.
The axial force further consolidates the plasticized
copper and a spot joint is formed. The joint width is
clearly visible on all the joints. Considerable areas of
both the top and bottom sheets are bonded together. The
hook does not extend to the keyhole area, which indi-
cates bonding between the sheets. The bond width was
measured and was found to be (7.5, 8.5 and 10) mm, res-
pectively, at the selected tool rotational speeds. An
increase in the tool speed improves the bond width
because the frictional-heat generation depends upon the
tool rotational speed.17 The higher the tool rotational
speed, the higher is the heat generation. Hence, more
copper is plasticized and the bond width increases.
Regions with different microstructures were observed
on these macrographs. They are the stir zone (SZ), the
thermomechanically affected zone (TMAZ), the heat-
affected zone (HAZ) and the base copper. These regions
are presented in the micrographs in Figure 6 as a func-
tion of the tool rotational speed. The stir zone is adjacent
to the keyhole area and displays very fine grains. FSSW
was derived from FSW. The plasticized material in FSW
undergoes dynamic recrystallization, which results in the
formation of fine grains. The width of the stir zone
reduces as the tool rotational speed is increased due to a
higher heat input. The TMAZ region presents slightly
elongated grains with an array of oxide particles. The
shearing of the plasticized material from the advancing
side to the retreading side causes the grains in this region
to elongate. The width of the TMAZ is found to reduce
with an increase in the tool rotational speed. The third
I. DINAHARAN, E. T. AKINLABI: INFLUENCE OF THE TOOL ROTATIONAL SPEED ON THE MICROSTRUCTURE ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 791–796 793
Figure 6: Light micrographs of the transition zone of friction-stir
spot-welded copper at: a) 1200 min–1, b) 1600 min–1 and c) 2000 min–1
Slika 6: Svetlobni posnetki mikrostrukture prehodne cone torno
vrtilno to~kasto zvarjenega bakra pri: a) 1200 min–1, b) 1600 min–1 in
c) 2000 min–1
Figure 5: Macrographs of friction-stir spot-welded copper at:
a) 1200 min–1, b) 1600 min–1 and c) 2000 min–1
Slika 5: Makro posnetek torno vrtilno to~kasto zvarjenega bakra pri:
a) 1200 min–1, b) 1600 min–1 in c) 2000 min–1
region, the HAZ, contains coarse grains. The coarseness
of the grains increases with an increase in the tool
rotational speed as presented in Figure 7.
A higher frictional heat causes the grains to become
coarse. The grain size in the HAZ region is higher com-
pared to the grain size of the as-received pure copper in
Figure 1. This region is clearly influenced by the
frictional heat generated during FSSW, leading to the
grain growth. The stir zone is very thin at the higher tool
rotational speed of 2000 min–1. The higher heat causes
the recrystallized grains to grow further and diminishes
the width of the stir zone. The microstructure of the un-
bonded zone is presented in Figure 8. The grains of the
upper and lower sheets are not uniform. The grains in the
upper sheet are coarser than those in the lower sheet.
This can be attributed to the proximity of each sheet to
the tool shoulder. The upper sheet receives more heat
input compared to the lower sheet, causing the grains to
become coarser.
The influence of the tool rotational speed on the
TSFL is depicted in Figure 9. It is evident from the
figure that the increase in the tool rotational speed
increased the TSFL. Although the increase in the tool
rotational speed caused grain growth in different regions
across the joint, the TSL improved upon the increased
tool rotational speed because the TSFL depends upon the
joint width and the hook position from the stir-zone area.
A hook is formed in a FSSW joint as the material from
the bottom sheet flows upwards during the process. The
distance of the hook from the stir-zone area determines
the load-carrying capacity of the joint. The hook is not
I. DINAHARAN, E. T. AKINLABI: INFLUENCE OF THE TOOL ROTATIONAL SPEED ON THE MICROSTRUCTURE ...
794 Materiali in tehnologije / Materials and technology 50 (2016) 5, 791–796
Figure 8: Light micrographs of the unbonded zone of friction-stir
spot-welded copper at: a) 1200 min–1, b) 1600 min–1 and c) 2000 min–1
Slika 8: Svetlobni posnetki nezvarjenega podro~ja torno vrtilno
to~kastega zvara bakra pri: a) 1200 min–1, b) 1600 min–1 in c) 2000 min–1
Figure 7: Light micrographs of the heat-affected zone of friction-stir
spot-welded copper at: a) 1200 min–1, b) 1600 min–1 and c) 2000 min–1
Slika 7: Svetlobni posnetki mikrostrukture toplotno vplivane cone
torno vrtilno to~kasto zvarjenega bakra pri: a) 1200 min–1, b) 1600 min–1
in c) 2000 min–1
clear in the macrograph in Figure 1 due to the annealing
effect of copper during FSSW.
The other factor is the bond width. The bond width
increased with an increase in the tool rotational speed.
This provided more tangential area to resist the external
load. Hence, the TSFL increased as the tool rotational
speed was increased. The photographs of failed speci-
mens are shown in Figure 10. The failed specimens de-
monstrate different modes of failure. The nugget pull-out
failure is observed at the tool rotational speeds of
1200 min–1 and 1600 min–1. It can be inferred from the
I. DINAHARAN, E. T. AKINLABI: INFLUENCE OF THE TOOL ROTATIONAL SPEED ON THE MICROSTRUCTURE ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 791–796 795
Figure 12: Images of tensile shear specimens during loading as
marked in Figure 11
Slika 12: Posnetki nateznih stri`nih vzorcev med obremenjevanjem,
kot je prikazano na Sliki 11
Figure 9: Effect of tool rotational speed on tensile shear load of
friction-stir spot-welded copper
Slika 9: Vpliv hitrosti vrtenja orodja na natezno stri`no obremenje-
nega torno vrtilno zvarjenega bakra
Figure 11: Load-versus-displacement graph for a tensile shear-tested
specimen at the tool rotational speed of 1600 min–1
Slika 11: Odvisnost obremenitve in raztezka pri nateznem stri`nem
preizkusu vzorcev pri hitrosti vrtenja orodja 1600 min–1
Figure 10: Images of failed sheets at tool rotational speeds of: a) 1200
min–1, b) 1600 min–1 and c) 2000 min–1
Slika 10: Posnetki poru{enih plo~evin pri hitrosti vrtenja orodja:
a) 1200 min–1, b) 1600 min–1 in c) 2000 min–1
photographs (Figures 10a and 10b) that the nugget
pull-out originated from the bottom of the keyhole area.
The diameter of the hole in the upper failed sheet is
higher at 1500 min–1 than 1200 min–1. This is due to an
increase in bond width, which provided a wider me-
tallurgical bonding. The mode of failure at the tool rota-
tional speed of 2000 min–1 is tearing. The tearing started
circumferentially around the unbonded region and could
not pull the nugget out because the higher bond width at
2000 min–1 caused the sheet to tear like a tensile speci-
men.
The load-versus-displacement curve for the tensile
shear-tested specimen at the tool rotational speed of
1600 min–1 is presented in Figure 11. The load increases
with a constant slope until point šc’ is reached. Corres-
ponding photographs in Figure 12 show details of the
points marked in Figure 11. As the load is increased to
point šb’, the specimen bends. The axis of the normal
load does not coincide with the centre of the tested spe-
cimen, which creates a small moment causing the bend-
ing. The specimen starts to yield, i.e., the nugget pull-out
commences at point šc’. The load does not drop sharply
to zero, unlike in the cases of tensile failures, because the
nugget pull-out is not instant. The nugget pulls out
gradually as the separation increases until point šg’ is
reached. The load drops to zero, which indicates a com-
plete separation of both sheets.
4 CONCLUSIONS
Commercially pure copper sheets were successfully
spot welded using the novel FSSW technique. The
influence of the key processing parameter, the tool
rotational speed, on the microstructure and TSFL was
analyzed. The increase in the tool rotational speed
increased the bond width and coarsened the grains in the
stir zone and HAZ. The TSFL improved as the tool rota-
tion increased. The nugget pull-out failure took place at
1200 min–1 and 1600 min–1 and the tear-out failure took
place at 2000 min–1.
5 REFERENCES
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I. DINAHARAN, E. T. AKINLABI: INFLUENCE OF THE TOOL ROTATIONAL SPEED ON THE MICROSTRUCTURE ...
796 Materiali in tehnologije / Materials and technology 50 (2016) 5, 791–796
J. ROZMAN et al.: MEASUREMENT OF BIO-IMPEDANCE ON AN ISOLATED RAT SCIATIC NERVE ...
797–804
MEASUREMENT OF BIO-IMPEDANCE ON AN ISOLATED RAT
SCIATIC NERVE OBTAINED WITH SPECIFIC CURRENT
STIMULATING PULSES
MERITEV BIOIMPEDANCE NA IZOLIRANEM @IVCU
ISCHIADICUS PRI PODGANI, VZBUJENEM S POSEBNIMI
TOKOVNIMI STIMULACIJSKIMI IMPULZI
Janez Rozman1,3, Monika C. @u`ek2, Robert Frange`2, Samo Ribari~3
1Center for Implantable Technology and Sensors, ITIS d. o. o., Lepi pot 11, 1000 Ljubljana, Slovenia
2Institute of Physiology, Pharmacology and Toxicology, Veterinary Faculty, University of Ljubljana, Gerbi~eva 60, 1000 Ljubljana, Slovenia
3Institute of Pathophysiology, University of Ljubljana, ZaMedical Faculty, lo{ka 4, 1000 Ljubljana, Slovenia
samo.ribaric@mf.uni-lj.si
Prejem rokopisa – received: 2015-09-30; sprejem za objavo – accepted for publication: 2015-10-26
doi:10.17222/mit.2015.307
In this study we designed and tested a four-probe, bio-impedance measurement set-up for peripheral nerves, based on the Red
Pitaya open-source measurement and control tool. The set-up was tested on an isolated rat sciatic nerve (RSN) while it was
stimulated with specific current stimulating pulses. The measurements tested the hypothesis that the specific waveform of a
stimulating pulse elicits current differences at the double layer (DL) interface between the platinum (Pt) stimulating electrode
and the nerve tissue. Impedance spectroscopy was used to electrically characterize the interface between the Pt electrode and the
nerve tissue and measure the interface electrical impedance (Z). An analysis of the frequency response and the impedance, with
specific current stimulating pulses, characterised the structure and the composition-related electrical properties of the RSN. An
analysis of the voltage responses (VRs), measured at the same time, showed that the maximum negative polarization across the
electrode-electrolyte interface (Emc) and the maximum positive polarization across the electrode-electrolyte interface (Ema) did
not exceed the safety limits for water electrolysis. We conclude that the voltage and current changes, elicited at the DL of the in-
terface between the Pt stimulating electrode and the nerve tissue, do not lead to tissue damage. Based on the obtained
electrophysiological results we conclude that the developed stimulating electrodes and the stimulus pattern could act as a useful
tool for developing nerve-stimulating electrodes.
Keywords: electrical impedance, electrical impedance spectroscopy, bio-impedance, voltage response, current source, platinum,
rat sciatic nerve, electrode-nerve tissue interface
Izdelali in preizkusili smo napravo za {tiri-to~kovno merjenje elektri~ne impedance (Z) perifernih `ivcev, katere osnova je
odprtokodna merilna in kontrolna konzola (Red Pitaya). Naprava je bila preizku{ana na izoliranem `ivcu ischiadicus pri podgani
(RSN) tako, da smo `ivec lahko dra`ili z izbranimi tokovnimi stimulacijskimi impulzi. Preverili smo hipotezo, da izbrani
stimulacijski impulzi izzovejo razlike na dvojni plasti (DL) prehoda med platinovo (Pt) stimulacijsko elektrodo in `iv~nim
tkivom. Z metodo impedan~ne spektroskopije smo izmerili Z na prehodu med Pt elektrodo in `iv~nim tkivom. Analiza
frekven~nega odgovora RSN, medtem ko je bil stimuliran s posebnimi tokovnimi stimulacijskimi impulzi, je podala od sestave
in zgradbe odvisne elektri~ne lastnosti `ivca. Nadalje je analiza so~asno merjenega napetostnega odgovora (VR) pokazala, da
nobena od obeh, tako najve~ja negativna polarizacija Emc kot tudi najve~ja pozitivna polarizacija Ema, na prehodu med elektrodo
in elektrolitom, ni presegla varne meje pri kateri lahko pride do elektrolize vode. Zaklju~ujemo, da napetostne in tokovne
razlike, ki so nastale na DL prehodu med Pt stimulacijsko elektrodo in `iv~nim tkivom, nimajo {kodljivega u~inka na `iv~no
tkivo. Kon~ni sklep raziskave je, da je mogo~e pridobljene elektrofiziolo{ke rezultate koristno uporabiti pri nadaljnem
na~rtovanju in razvoju stimulacijskih elektrod.
Klju~ne besede: elektri~na impedanca, elektri~na impedan~na spektroskopija, bioimpedanca, napetostni odgovor, tokovni vir,
platina, kol~ni `ivec podgane, prehod med elektrodo in `iv~nim tkivom
1 INTRODUCTION
Tissue-characterizing techniques based on electrical
impedance (Z) are being used to study the Z variations of
biological tissues over a range of frequencies. Some of
these analysing techniques provide Z values of the tissue
sample as a lumped estimation at a suitable frequency
range and the information on several complex bioelec-
trical phenomena occurring in cells and tissues under an
alternating (AC) electric current signal.1,2 In this regard,
a measurement of the complex Z in biological systems is
one of the developing technologies for monitoring and
determining the pathological and physiological status of
biological tissues.3 Namely, with AC electrical
stimulation, biological tissues produce a complex Z that
depends on the tissue composition, structures, health
status, and applied stimulus frequency.4
In biological tissues, plasma membranes of the cells
are composed of electrically non-conducting lipid
bilayers and an ion-conducting proteins channels.5,6 This
structure provides a capacitance to the applied AC signal
and contributes to the overall response of the biological
tissues by producing a complex Z, which is a function of
the tissue composition as well as the frequency of the
applied AC signal. Z is a measurement of the overall
Materiali in tehnologije / Materials and technology 50 (2016) 5, 797–804 797
UDK 621.317.33:616.833.58:599.323.452:543.42 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)797(2016)
ability of a medium to conduct electrical current, defined
as the ratio of the voltage in an object to the current in a
conductive medium. Z is a complex quantity, which
consists of a real part (resistance) and an imaginary part
(reactance).7 Its units are ohms (
). In this regard,
Wtorek et al.8 presented the construction of a probe for
immittance spectroscopy based on the four-electrode
technique. They tested the probe in "in-vivo"
measurements on swine gluteal tissue.
The simplest set-up for measuring Z is the so-called
two-probe set-up, where a current source is used in order
to feed a known current (I) into the unknown Z. A volt-
age measurement using a voltage-measurement device
determines the voltage drop, which is assumed to be pro-
portional to the unknown Z. However, the resistivity of
cables and the contact resistance also fully contribute to
the voltage drop and to the measured voltage. As a re-
sult, the measurement result Z is actually higher than the
true Z.
In the diagnosis of neuromuscular diseases, for in-
stance, the Z of muscles can be measured to monitor pa-
tients with amyotrophic lateral sclerosis, muscular dys-
trophy and inflammatory myopathies. In clinical
practice, a surface-electrode Z measurement is used for Z
measurements of the skin tissue, which can be used for
monitoring various conditions relating to the physical or
medical state of a patient. The accuracy of a Z measure-
ment with surface electrodes can be degraded by un-
known contact impedances.
In order to avoid the disadvantages of the two-probe,
Z measurement set-up, the four-probe, Z measurement
set-up employs an additional pair of electrodes. In the
four-probe, Z measurement set-up, the current is fed
through two feeding electrodes1,4 (considered as drive
terminals) within the line of four equidistantly placed
electrodes, while the voltage drop is measured between
two measurement electrodes2,3 (considered as
measurement terminals). The voltage is measured with
an instrument that has a very high input Z, so that almost
no current flows through the measurement electrodes,
i.e., their cables and contact resistance play almost no
role in the measurement. Therefore, the measured
voltage is almost identical to the voltage drop at the
unknown Z.9 For many applications, the four-probe, Z
measurement set-up thus provides sufficient accuracy.
The measurement of Z in biological systems, such as
nerve tissue, with the four-probe measurement set-up,
eliminates the influence of the Z of the electrode/nerve
tissue interface on the nerve tissue Z. Namely, electrode
polarization is avoided and the contact Z is eliminated
from the measurement.10,11 However, the Z measurement
error cannot be fully eliminated due to current flow
through the measurement electrodes. As the use of Pt
electrodes in selective nerve-stimulation applications
(SNS) became indispensable, significant research was
dedicated to understanding the electrode-electrolyte
interface in “in vivo” and the electrode-nerve tissue
interface in "in vivo". Of particular importance to SNS is
Schwan’s limit of linearity: the voltage at which the
electrode system’s Z becomes nonlinear and is often
exceeded during SNS.12,13
The main goal of this study was to use the electrical
impedance spectroscopy (EIS) and voltage response
(VR) experimental results to enhance our understanding
of the effect of physical processes at the interface Z, with
the expressed purpose of improving the interface design
of future multi-electrode stimulating systems for SNS.
The measurements were performed to test the hypothesis
that a specific waveform of current stimulating pulse
elicits controlled and tissue-safe voltage differences at
the double layer (DL) of the interface between the Pt
stimulating electrode and the nerve tissue. These repeti-
tive voltage differences are potentially harmful since
over time they can prevent effective excitation of the
nerve tissue at the stimulating electrode due to electroly-
sis-induced nerve-tissue damage.
2 MATERIAL AND METHODS
2.1 Experimental design and set-up
Among the several custom-designed pieces of equip-
ment, the most important was a temperature-controlled
measuring chamber and proprietary stimulation/record-
ing cables (Figure 1d). The chamber was developed for
the two-probe and the four-probe electrophysiological
measurements. The bulk body of the chamber was ma-
chined out of Plexiglas, using a computerized numerical
control milling machine. The main part of the chamber
was a ladder with seven Pt hook-shape electrodes
mounted horizontally over the chamber aperture and
filled with a Krebs-Ringer solution (154-mM NaCl,
5-mM KCl, 2-mM CaCl2, 1-mM MgCl2, 5-mM HEPES,
11-mM D-glucose, pH 7.4). The distance between the
electrodes was 4.8 mm.
The hook-shape electrodes were manufactured from a
cold-drawn Pt wire (99.99 % purity) with a thickness of
1 mm. Pure Pt is commonly used as a stimulating elec-
trode material because it can effectively supply high-
density electrical charge to activate the neural tissue. Pt
is normally used in the form of a pure metal because im-
purities and alloying elements may adversely affect its
mechanical characteristics, working characteristics and
its stability against corrosion in physiological media. Pt
also has many physical properties that are of great value
for their use in the technology of implantable stimulating
electrodes. The contact surface of the Pt electrodes was
increased, relative to the electrodes’ geometric area, by
sanding with a metallurgical grade sandpaper number
1000.
Consequently, in the developed measuring system,
the two Pt hooks (electrode 1 and 7) of the ladder,
representing the first two points of the four-probe
measuring set-up, are used as shown in Figure 1d. For
measurements of the voltage drop elicited by the current
J. ROZMAN et al.: MEASUREMENT OF BIO-IMPEDANCE ON AN ISOLATED RAT SCIATIC NERVE ...
798 Materiali in tehnologije / Materials and technology 50 (2016) 5, 797–804
pulse, the two hooks (electrode 3 and 5) representing the
second two probes of the four-probe measuring set-up,
were used. The hook electrode 1 is connected to the high
current source output (IH) of the bio-impedance front-end
device (Figure 1c), while the hook electrode 7 is
connected to the low current source output (IL) of the
bio-impedance front-end device. Electrode 3 was con-
nected to the high-voltage input (VH) of the bio impe-
dance front-end device and electrode 5 was connected to
the low-voltage input (VL) of the bio impedance
front-end device. Before the rat’s sciatic nerve (RSN)
was placed on the ladder of electrodes, the chamber
canal was filled with saline to prevent the drying out of
the RSN. To reduce the polarization and voltage shifts,
the electrodes were soaked for 15 min to 30 min.
2.2 Nerve dissection and preparation
Z measurements were performed on two isolated
RSNs harvested from adult male Wistar rats, weighing
300 g to 350 g, purchased from Charles River. Animal
handling complied with the Prevention of Cruelty to
Animals law, which is consistent with the European
Community Directive 2010/63/EU and was approved by
the Institutional Review Board. The rats are euthanized
by CO2 and exsanguination. Dissection began with an
incision through the skin on the lateral side of the hind
limb from the knee to the hip region. To expose the RSN,
the superficial layer of the thigh muscle is carefully
dissected between the m. gluteus superficialis and m.
biceps femoris. The length of RSN available for
measurement depends on the point at which the RSN
divides into the peroneal, tibial and sural nerves. To
prevent nerve injury during handling, a thread was tied
around the distal end of the RSN. After the attached
tissue was removed from the RSN, the RSN was cut as
close to the knee joint as possible. Similarly, at the
proximal end, the RSN was cut as close as possible to
the spinal column. By doing so, the RSN obtained was
long enough (between 30 mm and 35 mm) to be in
contact with all the electrodes in the ladder.
Afterwards, the RSN was placed in a glass sili-
con-lined Petri dish filled with Krebs-Ringer solution.
As the buffer was used for cell physiology experiments,
it was gassed with 100 % oxygen for 10–20 min, so that
the RSN could recover from the dissection. The RSN
was transferred from the Petri dish into the chamber and
placed on the ladder so that it touched all the electrodes.
The RSN was kept moist with saline to prevent it from
drying out and rendering it useless for the experiment.
Finally, to prevent evaporation and drying of the tissue
the chamber was covered with a lid. All of the recordings
were made at an ambient temperature of 22–23 °C.
2.3 Measuring system
The Red Pitaya (Red Pitaya, Solkan, Slovenia) devel-
opment platform was used for the stimulating/measuring
set-up shown schematically in Figure 1. This platform
enables signal generation, signal acquisition and signal
processing.
The set-up used two built in D/A and A/D converter
inputs enabling measurements with a 14-bit resolution at
a 125 MHz/s sampling rate. This sampling rate could be
adjusted with a decimal/dividing factor to 1, 8, 64, 1024,
8192 and 65536, respectively. Red Pitaya was attached to
the Z front-end device as shown in Figure 1c. This de-
vice was composed of a voltage-controlled current
source (VCCS) based on the Howland method. The
VCCS was designed to provide a constant current signal
(ic), up to 4 mA, independent of the value of the attached
load within a frequency range from 1 Hz to 1 MHz. The
Z front-end device enabled simultaneous, four- and
two-probe Z measurements.3,7
2.4 Bio-impedance spectroscopy measurements
The EIS-based frequency response studies of Z, of
any material, can provide the material’s structure- and
J. ROZMAN et al.: MEASUREMENT OF BIO-IMPEDANCE ON AN ISOLATED RAT SCIATIC NERVE ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 797–804 799
Figure 1: Schematic diagram of the stimulating/measuring set-up: a) programming part of the Z measurement set-up, b) Red Pitaya,
c) Z front-end device, d) 3D view of the measuring chamber with a ladder of electrodes
Slika 1: Shematski prikaz stimulacijsko/merilne naprave: a) programska enota naprave za merjenje Z, b) Red Pitaya, c) vhodno izhodna naprava
(Z), d) tridimenzionalna slika merilne celice z lestvico elektrod
composition-dependent electrical properties as well as its
frequency response.14 In this study, EIS was used to esti-
mate the complex (Z(f)) and phase angle ((f)) of the
RSN at different frequencies fi (fi: f1, f2, f3,..., fn) by mea-
suring the surface voltage (V(f)) developed for a constant
low-amplitude, low-frequency alternating sinusoidal cur-
rent (I(f)) injected into a RSN using both the two- and
four-probe methods.15–18 Afterwards, the frequency-de-
pendent electrical bio-impedance Z(fi) was found as the
transfer function of the RSN, and thus the Z(fi) was cal-
culated by dividing the voltage data (V(fi)) measurement
by the applied current (I(fi)), as shown in the following
Equation (1):
Z(fi) = V(fi)/I(fi) (1)
The programming part of the Z measurement set-up,
shown in Figure 1a, was performed with a Lenovo T420
portable computer (Lenovo, Singapore) and MATLAB
R2007a software (The Mathworks Inc., USA). In this
way, communication and concomitant functional control
of the Red Pitaya was made and monophasic and
biphasic voltage pulses to drive (driving pulse) a VCCS
within the front-end device, having a specific waveform
shown in Figures 2a and 2b, were generated. An afore-
mentioned waveform of the stimulating pulse was
adopted from a recent study where this waveform was
tested by selective nerve stimulation of the isolated por-
cine left cervical vagus nerve segment.19 Afterwards,
driving pulses were delivered to the VCCS, where they
were converted into constant current pulses by means of
the high-output-impedance current generator. To avoid a
measurement error in developing the presented Z mea-
surement set-up, it was essential that a high-output Z was
maintained over the operating frequency range so that
the injected current was constant and that the current
source did not load the sample.20 The stimulating current
pulses, having precisely the same waveform as the driv-
ing pulses that were delivered to the VCCS, were con-
stant current, biphasic, partially charge-balanced and
asymmetric, composed of a precisely determined
quasi-trapezoidal cathodic phase with a square leading
edge of intensity ic, a plateau of the cathodic phase with
the width of tc, followed by an exponentially decaying
phase with the width of texp and the time constant ôexp,
and ended by a wide, rectangular, anodic phase with the
width ta and intensity of ia.
In the four-probe measurement, the generated current
stimulating pulse waveform ic was applied to an isolated
RSN via the two outer hook electrodes within the ladder,
electrode number 1 and electrode number 7 that served
as current feeding probes IH and IL. They were separated
from the electrodes number 3 and 5 that served as volt-
age probes VH and VL, respectively. In the two-probe
measurement, the generated current stimulating pulse
was applied on an isolated RSN via the electrode number
1, serving as a current feeding probe IH, and connected to
the electrode number 3 that served as a voltage probe VH
and via the electrode number 7, that served as a current
feeding probe IL and was connected to electrode number
5 that served as a voltage probe VL.
These measurements were used to obtain the Bode
and Nyquist diagrams with which the equivalent circuit
model elements could be calculated for each of the sam-
ples. The Nyquist diagram consists of a Zimaginary vs.
Zreal plot, representing the imaginary and real values for
the Z. This paper however, only reports an “in-vitro”
study of RSN samples conducted using EIS to find the
frequency-dependent Z.
2.5 Voltage-response (VR) measurements
The presented set-up also enabled measurements of
VRs that were elicited in the RSN using generated
quasi-trapezoidal monophasic and biphasic pulses. The
generated stimulating pulse was measured indirectly as a
voltage drop at the reference resistor with a trains imped-
ance amplifier to assess the value of the phase angle .
Later on, this value was used in the processing of data
for the Z calculation. The aforementioned stimulating
pulse was used for excitation of the RSN that represented
a floating load; therefore, the elicited voltage drop was
measured differentially using an instrumental amplifier.
The RSN was stimulated with the generated current
pulses while information on the VR and stimulating cur-
rent was recorded simultaneously. In this regard, ic was
1.8 mA while tc was 60 μs. In each measurement, where
50 pulses were generated, the last pulse was saved and a
time to gap of 250 μs was introduced. The EIS was per-
J. ROZMAN et al.: MEASUREMENT OF BIO-IMPEDANCE ON AN ISOLATED RAT SCIATIC NERVE ...
800 Materiali in tehnologije / Materials and technology 50 (2016) 5, 797–804
Figure 2: Voltage driving pulses to control the VCCS current source:
a) a monophasic pulse and b) a biphasic pulse
Slika 2: Napetostna krmilna pulza za kontrolo VCCS tokovnega vira:
a) monopolarni pulz in b) bipolarni pulz
formed using current sinusoidal signals with an ampli-
tude of 400 μA that were generated at frequencies evenly
distributed within the frequency range between 1 Hz and
500 kHz and delivered to the RSN. While the RSN was
stimulated with the aforementioned current sinusoidal
signals, the stimulating current and VR were saved again.
All of the data were analysed off-line and Z was calcu-
lated with the lock-in method.
3 RESULTS
3.1 Bio-impedance spectroscopy
Figure 3a shows the Bode diagram of the absolute
value of Z, measured on the RSN, using the two- and
four-probe set-ups while the RSN is stimulated with two
specific monophasic and biphasic excitation pulses, re-
spectively. Figure 3b shows the  obtained with the two-
and four-probe set-ups expressed in the Bode diagram
during RSN stimulation with monophasic and biphasic
stimuli. Figure 3c shows the Z spectrum measured with
the two- and four-probe set-ups, presented with the
Nyquist diagram (Cole-Cole diagram), while the RSN is
stimulated with monophasic and biphasic stimuli.
The absolute Z value, expressed in the Bode diagram
while the RSN is stimulated with monophasic and
biphasic stimuli and measured using the two-probe
set-up, decreased almost exponentially from approxi-
mately 60 k
 measured at 1 Hz to slightly below 800 

measured at 20 kHz (Figure 3a). At frequencies between
20 kHz and 100 kHz, Z remained almost unchanged; at
frequencies above 100 kHz, however, Z began to
decrease rapidly. It was noted that traces of Z belonging
to stimulation with monophasic and biphasic stimuli
were almost identical.
Figure 3b, representing the  expressed in the Bode
diagram, shows that for the two-probe measurement, the
capacitive nature of an interface between RSN and plati-
num electrodes prevailed at frequencies below 500 Hz.
Namely, a corresponding  measured at 1 Hz was –60 °C
and  measured at 500 Hz was –10 °C, respectively. At
frequencies between 500 Hz and 30 kHz, the nature of
the interface remained still slightly capacitive, express-
ing the same value of the phase angle , i.e., –10 °C.
Within this frequency range, the ohmic nature of the in-
terface between the RSN and platinum electrodes pre-
vailed. At frequencies above 100 kHz, however, the na-
ture of the interface between the RSN and the platinum
electrodes started to be progressively capacitive again,
expressing the sharp decrease of  towards negative val-
ues. The traces of , belonging to the stimulation with
the monophasic and biphasic stimuli, were almost identi-
cal.
In Figure 3c, the Z spectra measured with two- and
four-probe set-ups expressed in a Nyquist diagram
(Cole-Cole diagram) while the RSN is stimulated with
monophasic and biphasic stimuli differ considerably.
This diagram represents the relationship between the
imaginary and the real component of Z, as measured at
the interface between the RSN and the particular plati-
num electrodes using the two- versus four-probe set-up
and monophasic versus biphasic stimuli. In other words,
this diagram represents the relationship between capaci-
tive and ohmic natures of the interface between the RSN
and the platinum electrodes. The difference was not sig-
nificant for monophasic versus biphasic stimuli but was
for the two- versus four-probe measurements. Traces of
Z spectra, belonging to stimulation with monophasic and
biphasic stimuli, were almost identical for the two- as
well as for the four-probe measurements. For the two-
J. ROZMAN et al.: MEASUREMENT OF BIO-IMPEDANCE ON AN ISOLATED RAT SCIATIC NERVE ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 797–804 801
Figure 3: Results of Z spectroscopy measurements while the RSN is
stimulated with monophasic and biphasic stimuli: a) Absolute Z value
measured with a two- and four-probe set-ups expressed in a Bode dia-
gram, b)  obtained with a two- and four-probe set-ups expressed in a
Bode diagram, c) Z spectrum measured with two- and four-probe
set-ups expressed in a Nyquist diagram (Cole-Cole diagram)
Slika 3: Rezultati spektroskopskih meritev Z, medtem, ko je bil RSN
stimuliran z monofaznimi in bifaznimi pulzi: a) absolutna vrednost Z
merjena z dvo in {tirito~kovno metodo, prikazana z Bodejevim
diagramom, b)  merjen z dvo in {tirito~kovno metodo, prikazan z
Bodejevim diagramom, c) spekter Z merjen z dvo in {tirito~kovno
metodo prikazana z Nyquistovim diagramom (Cole-Cole diagram)
and four-probe measurement, the Z spectrum below 10
k
 of the real/ohmic component, showed an almost
identical relationship between the imaginary/capacitive
and the real/ohmic nature of the interface between the
RSN and the Pt electrodes. Above 10 k
 of real/ohmic
component, however, the two-probe measurement
showed an almost linear decrease of an imaginary/capac-
itive nature and, at the same time, an almost linear in-
crease of the real/ohmic nature, while the four-probe
measurement showed no further increase in the
real/ohmic nature.
3.2 Voltage-response (VR) measurements
VRs, during the excitation of the RSN with the
quasi-trapezoidal monophasic and biphasic pulses using
the two- and four-point measurement set-ups, are shown
in the Figures 4a to 4c. Figure 4a shows the VRs elic-
ited with a monophasic pulse and measured with the
two- and four-point set-ups. VRs measured with the
four-point set-up follow the quasi-trapezoidal pulse more
accurately than the VRs measured using the two-point
set-up. This was in accordance with Figure 3 where the
real/ohmic nature of the interface between the RSN and
the Pt electrode in the four-point measurement prevails
over the imaginary/capacitive nature of the interface be-
tween the RSN and the Pt electrode.
Figure 4b shows VRs measured with the two- and
four-point set-ups elicited with the biphasic pulse. VRs
measured with the two-point set-up follow the
quasi-trapezoidal pulse more accurately than the VRs
measured with the four-point set-up. This is in accor-
dance with Figure 3 where the real/ohmic nature of the
interface between the RSN and the Pt electrode in the
four-point measurement prevails over the imaginary/ca-
pacitive nature of the interface between the RSN and the
Pt electrode.
As seen in all three panels of Figure 3, the leading
edge of the cathodic phase ic and the plateau of the cath-
odic phase with the width tc of the generated stimulating
current pulses were not of exactly the same shape as the
waveform of pulses that were delivered to the VCCS.
Other parts of the generated current pulses, such as expo-
nentially decaying phase with the width of texp and the
time constant exp in both monophasic and biphasic
pulses as well as the rectangular, anodic phase with the
width ta and intensity of ia in the biphasic pulse, are al-
most identical. It was presumed that in calculations of Z,
after a monophasic and biphasic pulse, the aforemen-
tioned dissimilarity did not influence Z significantly.
Figure 4c shows VRs measured using the four-point
set-up, elicited with a train of generated biphasic pulses
with a cathodic intensity (ic) of 1.8 mA and the elicited
VRs with the entire voltage drop of V = 1.7 V. As
shown in Figure 4c, an onset of ic elicited the near-in-
stantaneous voltage increase with the same time course.
However, when ic was terminated, the course of the volt-
age in the exponential decay region where ic exponen-
tially approached the lowest value, the voltage also expo-
nentially approached the value of approximately -0.15 V.
Almost at the same time as the cathodic pulse ic was ter-
minated, the onset of an anodic intensity (ia) of 0.6 mA
elicited the near-instantaneous positive voltage of ap-
proximately 0.65 V, having a slightly different course.
As in the VR of the tested stimulation pulse, the max-
imum negative polarization across the electrode-nerve
tissue interface (Emc) and the maximum positive polariza-
tion across the electrode-nerve tissue interface (Ema)
reached -0.15 V and 0.65 V, respectively. None of them
exceed the safety limit range for water electrolysis, from
-0.60 V to +0.85 V.21
J. ROZMAN et al.: MEASUREMENT OF BIO-IMPEDANCE ON AN ISOLATED RAT SCIATIC NERVE ...
802 Materiali in tehnologije / Materials and technology 50 (2016) 5, 797–804
Figure 4: Stimulation pulses and corresponding VRs: a) VRs measured
using a two- and four-point measurements elicited with a monophasic
pulse, b) VRs measured using a two- and four-point measurements
elicited with a biphasic pulse, c) VRs measured using a four-point
measurement elicited with a train of biphasic pulses
Slika 4: Stimulacijski pulzi in pripadajo~i VR-ji: a) VR-ji merjeni z
dvo in {tirito~kovno metodo, izzvani z monofaznim pulzom, b) VR-ji
merjeni z dvo in {tirito~kovno metodo, izzvani z bifaznim pulzom, c)
VR-ji merjeni z dvo- in {tirito~kovno metodo, izzvani z vlakom
bifaznih pulzov
4 DISCUSSION
The presented work demonstrates the feasibility of a
Z measurement on an isolated RSN with quasi-trape-
zoidal current biphasic stimulating pulses. The feasibility
of the stimulating paradigm, where the specific stimulus
waveform was used to selectively stimulate different
types of nerve fibres within the cervical segment of por-
cine vagus nerve, has already been published.19
In the last two decades, particular attention has been
paid to vagus nerve stimulation techniques that are used
to treat, among others, a number of nervous-system dis-
orders, neuropsychiatric disorders, eating disorders,
sleep disorders, cardiac disorders, endocrine disorders,
and pain.22–25 Considerable scientific and technological
efforts have been devoted to developing systems of elec-
trodes that interface the human vagus nerve with elec-
tronic implantable devices.26,27
The Z characterization of the electrode-electrolyte in-
terface is of paramount importance in the field of
neuroprotheses, where small electrodes are required for
SNS and recording of superficial regions of peripheral
nerves via multi-electrode systems. The electrode-elec-
trolyte interface is determined from the known electrode
Z, which should be as low as possible to avoid nerve tis-
sue damage. A high Z would result in a large applied
electrode voltage leading to undesirable electrochemical
reactions that damage the nerve tissue.28 As shown in the
presented study, neither the maximum negative polariza-
tion across the electrode-electrolyte interface (Emc), nor
the maximum positive polarization across the elec-
trode-electrolyte interface (Ema), exceeded the safety lim-
its for water electrolysis.21
Nevertheless, large surface area stimulating elec-
trodes (due to a roughened surface), with a relatively
much smaller geometric surface area and thus a much
lower Z than current designs, could be beneficial, espe-
cially where a larger number of stimulating electrodes
within the multi-electrode system29 is needed. The au-
thors demonstrated “in vitro” that by increasing the real
surface area of stimulating electrodes, a significantly
lower polarization impedance and electrode impedance,
as well as a low residual direct current, could be ob-
tained. Also, all three electrical parameters, electrode im-
pedance, access resistance and polarization, could be cal-
culated from the current and voltage measurements.
Thus, a theoretical equivalent circuit model of the inter-
face between the RSN and Pt electrodes could be mod-
elled by fitting the equivalent circuit model elements to
the EIS-based frequency-response spectroscopy experi-
mental results.28,30,31
5 CONCLUSIONS
The first conclusion is that at frequency range
between 0 kHz and 2.5 kHz, which confines the power
spectrum density of the proposed quasi-trapezoidal
stimulating pulse (power spectrum density not shown in
this paper), the ohmic nature of the interface between
RSN and Pt electrodes prevailed over the imaginary/
capacitive nature, yielding the relative low Z of approxi-
mately 1 k
.
The second conclusion is that the differences elicited
at the DL of the interface between the Pt stimulating
electrode and the nerve tissue, while the RSN was stimu-
lated using biphasic quasi-trapezoidal stimuli, have not
exceeded the voltage range considered safe for nerve tis-
sue.
The third conclusion is that the designed and tested
four-probe bio-impedance measurement set-up, based on
the Red Pitaya open-source measurement and control
tool, provided data on the structure- and composition-re-
lated electrical properties of the RSN tissue that contrib-
ute to the development of multi-electrode nerve stimulat-
ing systems.
Definitions of field-specific terms
SNS – selective nerve stimulation
RSN – segment of a rat sciatic nerve
AC – alternating current
Z – electrical impedance
VCCS – voltage controlled current source
driving pulse – voltage pulses to drive a VCCS
EIS – Electrical impedance spectroscopy
Pt – platinum
chamber – temperature-controlled measuring chamber
CNC – computerized numerical control
ic – intensity of the cathodic phase
tc – width of the cathodic phase
texp – width of the cathodic exponential decay
exp – time constant of the exponential decay
ia – intensity of the anodic phase
ta – width of the anodic phase
IH and IL – current feeding probes
VH and VL – voltage probes
VR – voltage response
R – reference resistor
 – phase angle
DL – double layer
Ema – maximum positive polarization across the elec-
trode-electrolyte interface
Emc – maximum negative polarization across the elec-
trode-electrolyte interface
Acknowledgement
This work was financed by the research grant
P3-0171 and P4-0053 from the Slovenian Research
Agency (ARRS), Ministry of Education, Science and
Sport, Ljubljana, Republic of Slovenia. The authors ac-
knowledge Professor Dejan Kri`aj from the Faculty of
Electrical Engineering, University of Ljubljana, for the
J. ROZMAN et al.: MEASUREMENT OF BIO-IMPEDANCE ON AN ISOLATED RAT SCIATIC NERVE ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 797–804 803
useful discussion and enabling us to use a measuring set
up developed in his laboratory.
Ethical approval
The neural tissue was treated in accordance with the
approval guidelines of the ethics committee of the Veteri-
nary Administration, Ministry of Agriculture, Forestry
and Food, Republic of Slovenia.
Conflict of interest statement
The authors declare that there are no conflicts of in-
terest regarding the publication of this paper. None of the
authors received any financial or other compensation for
undertaking this study or during its execution. No per-
sonal relationships with other people or organizations in-
appropriately influenced the work.
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804 Materiali in tehnologije / Materials and technology 50 (2016) 5, 797–804
A. KOCIJAN et al.: INFLUENCE OF DIFFERENT PRODUCTION PROCESSES ON THE BIODEGRADABILITY ...
805–811
INFLUENCE OF DIFFERENT PRODUCTION PROCESSES ON THE
BIODEGRADABILITY OF AN FeMn17 ALLOY
VPLIV RAZLI^NIH PROCESOV IZDELAVE NA
BIORAZGRADLJIVOST ZLITINE FeMn17
Aleksandra Kocijan, Irena Paulin, ^rtomir Donik, Matej Ho~evar, Klemen Zeli~,
Matja` Godec
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
aleksandra.kocijan@imt.si
Prejem rokopisa – received: 2016-03-31; sprejem za objavo – accepted for publication: 2016-04-06
doi:10.17222/mit.2016.055
The purpose of this research was to evaluate the biodegradability of a cast FeMn17 alloy that was processed by hot rolling and
annealing, which influenced both the mechanical and corrosion properties of the FeMn17 alloy. The corrosion behaviour and
in-vitro biodegradability were investigated by light microscopy, scanning electron microscopy, X-ray diffraction and immersion
tests in Hank’s solution. Compared to pure Fe, the cast FeMn17 alloy has a better biodegradability (higher corrosion rate). We
showed that hot rolling additionally improves the biodegradability, while the annealing process lowers the biodegradability of
the FeMn17 alloy.
Keywords: biodegradability, FeMn alloy, corrosion, XRD, SEM
Namen raziskave je bil oceniti biorazgradljivost zlitine FeMn17 s predelavo z vro~im valjanjem in `arjenjem. Vro~e valjanje in
`arjenje vplivata na mehanske in korozijske lastnosti zlitine FeMn17. Korozijske lastnosti in biorazgradljivost smo preverjali s
svetlobno mikroskopijo, vrsti~no elektronsko mikroskopijo, rentgensko difrakcijo in s testi potapljanja v Hankovo raztopino. V
primerjavi s ~istim Fe, ima zlitina FeMn17 veliko bolj{o biorazgradljivost (ve~jo korozijsko hitrost). Pokazali smo, da vro~e
valjanje dodatno izbolj{a biorazgradljivost, z `arjenjem pa se biorazgradljivost zlitine FeMn17 poslab{a.
Klju~ne besede: biorazgradljivost, FeMn zlitina, korozija, XRD, SEM
1 INTRODUCTION
Biodegradable metallic materials represent a novel
class of bioactive biomaterials that can temporarily sup-
port tissue healing and should progressively degrade
completely without a negative effect on the healing pro-
cess.1–3 Potential applications of these biomaterials are
paediatric, orthopaedic (fixation screws and pins) and
cardiovascular implants (coronary stents).1,4 Biodegrad-
able polymers were first investigated as bioactive bio-
materials; however, in recent years biodegradable metal-
lic materials, especially Fe and Mg alloys, have received
more attention due to their superior mechanical proper-
ties and their cytocompatibility.1–3,5 Compared with
Mg-based materials, Fe-based materials possess similar
mechanical properties to stainless steel and are more at-
tractive for applications that require high strength and
ductility.1 Despite the immense potential of Fe and Mg
alloys, experiments and clinical trials also exposed their
weaknesses: too rapid degradation rates, poor mechani-
cal properties and significant hydrogen evolution during
the corrosion process of Mg-based alloys and a too slow
degradation of Fe-based alloys.5–8
Fe-based alloys may also present problems with cer-
tain imaging devices (magnetic resonance imaging, for
example) due to the Fe’s ferromagnetic nature. However,
alloying and heat treatment can modify the mechanical,
corrosion, and ferromagnetic properties of pure Fe.1,2
The choice and the amount of alloying element is impor-
tant with respect to the toxicity and degradation behav-
iour of the Fe alloy.2 Mn represents a suitable alloying
element based on microstructural, magnetic, corrosion,
and toxicological considerations.2 Mn (austenite-forming
element) transforms Fe into a nonmagnetic material,
lowers the standard electrode potential of Fe and thus en-
hances the degradation of the material, which represents
an essential trace element necessary in many enzymatic
reactions. Newly developed Fe-Mn-based alloys contain-
ing up to 35 % of mass fractions of Mn have comparable
mechanical properties to stainless steel, faster degrada-
tion and improved MRI compatibility.1,2 Despite this, the
low corrosion rate still represents the major problem
fornewly developed Fe–Mn-based alloys.
Non-conventional processing techniques such as
powder metallurgy, electrodeposition and inkjet 3D-
printing can achieve the faster degradation of Fe–Mn al-
loys.1,2 However, these techniques are rather complex
and expensive, therefore it is necessary to further investi-
gate conventional methods, such as casting with addi-
tional steel-processing techniques in order to find an eco-
nomically favourable solution. Research has rarely been
made on the influence of subsequent processing and heat
treatment on the properties of any conventional, cast,
Materiali in tehnologije / Materials and technology 50 (2016) 5, 805–811 805
UDK 620.1/.2:669.056:67.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(5)805(2016)
biodegradable, metallic materials, including Fe–Mn-
based alloys. The aim of the present study was to investi-
gate the influence of three basic steel-processing meth-
ods (casting, hot rolling and annealing) on the corrosion
behaviour of the biodegradable Fe–Mn-based alloy due
to the formation of less-corrosion-resistant deformational
martensite.9
2 EXPERIMENTAL PART
Material preparation – The investigated Fe-based al-
loy with 17 % of mass fractions of Mn (FeMn17) was
produced from relatively pure Fe and Mn. Both materials
were melted in an induction furnace under air atmo-
sphere at approximately 1700 °C and cast into two iron
moulds. One mould was left to cool down in air
atmosphere and the other one was after casting in mould,
hot rolled at approximately 1000 °C for a 33 % reduc-
tion. Parts of both samples were annealed at 1050 °C for
1 h and then furnace cooled. The materials were ana-
lysed and the results were compared with pure Fe. The
chemical compositions of the alloy and the pure Fe were
determined using an X-ray fluorescence spectrometer
XRF (Thermo Scientific Niton XL3t GOLDD+) and
arepresented in Table 1.
Metallographic investigation – Samples were cut
with a water-cooled saw and cross-sections of the sam-
ples were prepared by the standard metallographic tech-
niques of grinding and polishing. The samples were
etched with 3 % Nital, an ethanol solution of HCl(aq) and
an ethanol solution of HNO3(aq). The characterization of
the material was performed using a light microscope
(LM, Microphot FXA Nikon with Olympus DP73) and a
scanning electron microscope coupled with an en-
ergy-dispersive spectrometer (SEM, JEOL JSM-6500F,
EDS INCA ENERGY 400) for analyses of the inclusions
and phases.
X-ray Powder Diffraction – The samples were mea-
sured using a Panalyitical XPERT Pro PW 3040/60
goniometer 2 between 15–90° with a step size of 0.002°
and a scan step time of 100 s on each step. Cu with (K =
0.154 nm) anode was used with a current of 40 mA and a
potential of 45 kV.
Mechanical properties – The microhardness was
measured by Vickers HV5 (5 kg load, 11s -Instron
Tukon 2100B).
Electrochemical measurements – Were performed on
prepared specimens, ground with SiC emery paper down
to 1200 grit. The experiments were carried out in a simu-
lated physiological Hank’s solution, containing 8 g/L
NaCl, 0.40 g/L KCl, 0.35 g/L NaHCO3, 0.25 g/L
NaH2PO42H2O, 0.06 g/L Na2HPO42H2O, 0.19 g/L
CaCl22H2O, 0.41 g/L MgCl26H2O, 0.06 g/L
MgSO47H2O and 1 g/L glucose, at pH = 7.8. All chem-
icals were from Merck, Darmstadt, Germany. The mea-
surements were performed using a three-electrode, flat
BioLogic corrosion cell (volume 0.25 L). The test speci-
men was employed as the working electrode (WE). The
reference electrode (RE) was a saturated calomel elec-
trode (SCE, 0.242 V vs. SHE) and the counter electrode
(CE) was a platinum net. Electrochemical measurements
were recorded by using a BioLogic Modular Research
Grade Potentiostat/Galvanostat/FRA Model SP-300 with
an EC-Lab®software V10.44. The specimens were im-
mersed in the solution 1 h prior to the measurement in
order to stabilize the surface at the open-circuit potential
(OCP). The potentiodynamic curves were recorded after
1 h of sample stabilisation at the open-circuit potential
(OCP), starting the measurement at 250 mV vs. SCE
more negative than the OCP. The potential was then in-
creased, using a scan rate of 1 mV s–1, until the
transpassive region was reached. The linear polarisation
measurements were performed at ±25 mV according to
the OCP, using a scan rate of 0.01 mV s–1.
3 RESULTS AND DISSCUSION
3.1 Microstructure characterization
The microstructures of all four samples, i.e., cast, hot
rolled and both samples after annealing at 1050 °C and
furnace cooled down, are shown in Figure 1. There is no
difference in the chemical composition, where as there
are some differences in the microstructure. In the cast
sample, the cast structure occurs upon cooling, though
the cooling rate is too high for equilibrium phase trans-
formation, as predicted from the Fe–Mn phase diagram.
We observed some boundary segregations rich in Mn and
inclusions of MnS and TiN, which were also confirmed
by the EDS analyses. The observed average grain size
for these samples was approximately 350 μm. There was
also some smaller porosity present in the microstructure.
The cast + annealed sample (Figure 1b) has a similar
microstructure with the same precipitations and inclu-
sions that are ubiquitous at grain boundary and in the
grains. The difference is in the much larger grain sizes
(average of approx. 1200 μm) and more segregation of
Mn at the grain boundaries according to the prolongation
of the time available for the cooling after annealing.
In the hot-rolled sample was observed, beside austen-
ite, also traces of strain-induced martensite and deforma-
tion twins. Due partly to the recrystallization, the micro-
structure consists of large and small grains. The
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806 Materiali in tehnologije / Materials and technology 50 (2016) 5, 805–811
Table 1: Chemical composition in mass fractions, w/%
Tabela 1: Kemijska sestava v masnih dele`ih, w/%
Material Mo Ni Mn Cu Ti Si C Fe
Pure Fe <LOD 0.14 0.262 0.043 <LOD 0.051 0.19 balance
FeMn17 alloy 0.021 0.182 17.08 0.446 0.058 0.167 0.08 balance
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Materiali in tehnologije / Materials and technology 50 (2016) 5, 805–811 807
Figure 1: LM images of microstructure: a) cast sample, b) cast + annealed sample, at 1050 °C c) hot-rolled sample, and d) hot-rolled + annealed
sample
Slika 1: LM-posnetki mikrostruktur: a) ulit vzorec, b) ulit + `arjen na 1050 °C, c) vro~e valjan vzorec in d) vro~e valjan + `arjen vzorec
Figure 2: SEM images of microstructure: a) cast sample, b) cast + annealed sample, at 1050 °C, furnace cooled down, c) hot-rolled sample and
d) hot-rolled + annealed at 1050 °C
Slika 2: SEM-posnetki mikrostruktur: a) ulit vzorec, b) ulit + `arjen na 1050 °C, c) vro~e valjan vzorec in d) vro~e valjan + `arjen vzorec
increased concentration of Mn was observed on the grain
boundaries well as the MnS and TiN inclusions in the
matrix as determined by the EDS analyses.
After annealing the sample has a recrystallized
microstructure with smaller grains, an average of ap-
proximately 40 μm and the strain-induced martensite is
still present in the microstructure. This means the an-
nealing temperature was too low or the annealing time
was too short for complete recrystallization. The SEM
imaging was performed at higher magnification for de-
tails of the microstructures (Figure 2). In the microstruc-
tures of the hot rolled and the hotrolled + annealed
FeMn17 samples the similar mixture of martensitic and
austenitic microstructure was observed, similar to those
previously described by Y. K. Lee at al.10, J. Martinez at
al.11 and M. Schinhammer at al.12 However, there was no
literature presenting only cast FeMn17 material. The
EDS analyses were performed for a determination of the
precipitations and inclusions in the microstructure (Fig-
ure 3).
3.2 Microhardness
Microhardness by Vickers (HV) was measured with a
5-kg load for 11 s. The results are in Table 2. The pure
Fe has, as expected, the lowest microhardness. The
higher microhardness is achieved in the hot-rolled sam-
ples, whereas the HV decreases with the annealing of the
samples. In cast sample, the HV is not different accord-
ing to the sample cast + annealed and cooled down with
the furnace. On the other hand, all the HVs are lower
thanthe one of the hot-rolled samples.
Table 2: Microhardness HV 5
Tabela 2: Mikrotrdota po HV5
Samples Pure Fe Cast Hotrolled
Cast +
annealed
Hot rolled +
annealed
HV 5 110 181 240 180 210
3.3 X-ray powder diffraction
Diffraction peaks of the FeMn17 specimens corre-
spond to the phases previously observed in FeMn17.
These phases are similar to those observed by Y. K. Lee
et al.10, J. Martinez et al.11 and H. Hermawan et al.13 in
various Fe–Mn binary alloys at room temperature. Fig-
ure 4 shows observed peaks for the FeMn17 alloy and
two starting materials, pure Fe and Mn. The highest peak
at 2 = 43.6° corresponds to (111), the next peak at 50.5°
to (002) and the last peak, at 74.4°, corresponds to (022).
The results correspond to a cubic crystal system, with the
space group F m –3 m and space group number 225, the
cubic cell unit is face centred (fcc). The XRD spectrum
shows peaks at different 2 for pure Fe and Mn, which
proves that all the starting materials were melted in a
new alloy with a different cell unit and different chemi-
cal composition and mechanical properties was prepared.
The highest intensity peak forthe pure Fe, as a start-
ing material, is observed at 44.8° and corresponds to
(011), the next peak is observed at 82.4° corresponding
to (112) and the peak with lowest intensity was at 65.1°,
which corresponds to (002). This is also a cubic crystal
system with a different space group, I m –3 m, and the
space number is 229. The cell unit for used pure Fe, as
the main alloying material is body centred cubic (bcc).
The other starting material used for the alloying was
Mn. The XRD spectrum for the Mn shows main peaks at
43.0°, which corresponds to (330), next peak at 47.8°
corresponds to (332) and two peaks with similar intensi-
ties at 52.3° and 78.9° that correspond to (510) and
(721), respectively. This is also a cubic crystal system
with the space group I –4 3 m and the corresponding
space group number 217, the cubic cell unit is body
centred cubic (bcc).
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808 Materiali in tehnologije / Materials and technology 50 (2016) 5, 805–811
Figure 3: SEM image of microstructure and areas of EDS measure-
ments on the hot-rolled sample
Slika 3: SEM-posnetek mikrostrukture z mesti EDS-analiz vro~e
valjanega vzorca
Figure 4: The XRD patterns for specimens – the two original materi-
als,pure Fe and Mn, and the specimen of cast FeMn17 alloy
Slika 4: XRD-spektri dveh osnovnih materialov, ~istega Fe in Mn, ter
ulite zlitine FeMn17
From the observed results it can be concluded that
the alloy consists of just one phase, FeMn and some
non-metallic inclusions. Any other potential phases in
the alloy are below 2 %, which is also the limit of detec-
tion of the XRD technique. This corroborates SEM/EDS
and LM results where onlysome additional non-metallic
inclusions and Mn-rich grain-boundaries are present in
the matrix of the investigated FeMn17 alloy.
Figure 5 shows the comparison of XRD spectra be-
tween differently processed FeMn17 alloys. The only
difference between the spectra is in the intensities of the
peaks and not in the peak positions, which corroborates
the fact that the only difference between these specimens
was just the thermal and mechanical treatments. The
most intense peak was for the first peak at 43.6° of
FeMn17 that was cast in a metal mould, and hot rolled,
while the preferential orientation was observed due to
the hot rolling, which explains the decrease of the inten-
sity of the peak at 74.4° in comparison to the other speci-
mens. The annealing of the specimens reduces the pref-
erential orientation so that the intensities in the XRD
spectrum are shifted to lower values and become very
similar to the samples just cast in a metal mould.
3.4 Electrochemical measurements
Figure 6 represents the potentiodynamic behaviour
of the investigated materials in a simulated physiological
Hank’s solution at pH = 7.8. The influence of different
production parameters on the corrosion behaviour of the
tested materials compared to pure Fe was investigated.
After 1 h of stabilization at the OCP, the corrosion poten-
tial (Ecorr) for pure Fe in Hank’s solution was approxi-
mately –0.67 V vs. SCE. Hot rolled + annealed FeMn17
as well as cast + annealed FeMn17 alloys exhibited simi-
lar potentiodynamic behaviour as pure Fe, with Ecorr at
approximately –0.68 V vs. SCE. In the case of cast +
hot-rolled FeMn17 alloys the corrosion stability de-
creased compared to other three specimens and the Tafel
region was shifted to higher corrosion-current densities
and Ecorr was –0.74 V vs. SCE. None of the investigated
samples in Hank’s solution exhibited typical
potentiodynamic behaviour with a passive region fol-
lowed by a transpassive region.
Figure 7 represents linear polarisation curves for the
investigated samples in a simulated physiological Hank’s
solutionat pH = 7.8. The calculations were performed
from linear polarization measurements using Equation
(1):
Rp = ac/(2.3 Icorr (a + c)) (1)
The polarization resistance, Rp, is evaluated from the
linear polarization curves by applying a linear least-
squares fit of the data around ±10 mV Ecorr. The corro-
sion current, Icorr, is calculated from Rp, the least-squares
slope, and the Tafel constants, a and c, of 120 mV de-
cade–1. The value of E(I = 0) is calculated from the
least-squares intercept. The corrosion rate was calculated
by using the following conversion equation (2):
vcorr= C(EW/d)(Icorr/A) (2)
where EW is the equivalent weight of the sample in g, A
is the sample area in cm2, d is its density in g/mL, and C
is a conversion constant, which is dependent upon the
desired units. The results demonstrated steeper linear
polarisation curves and therefore a lower corrosion
resistance of cast + hot-rolled FeMn17 alloys compared
to pure Fe, cast + annealed FeMn17 and hot rolled +
annealed FeMn17 alloys. The corrosion parameters cal-
culated from the linear polarisation and potentiodyna-
mic measurements in a simulated physiological Hank’s
solution indicated that the corrosion stability of the cast
+ hot rolled FeMn17 alloys was lower compared to pure
Fe, cast + annealed and hot rolled + annealed FeMn17
alloys. This was proven by the higher corrosion currents
(Icorr), the higher corrosion rates (vcorr) and the lower
polarization resistances (Rp), compared to pure Fe, cast
+ annealed and hot rolled + annealed FeMn17 alloys
(Table 3).
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Materiali in tehnologije / Materials and technology 50 (2016) 5, 805–811 809
Figure 5: XRD patterns for specimens of differently produced
FeMn17 alloy
Slika 5: XRD-spekter razli~no obdelane zlitine FeMn17
Figure 6: Potentiodynamic curves for pure Fe, cast, hot rolled, cast +
annealed and hot rolled + annealed FeMn17 alloy in a simulated
physiological Hank’s solution
Slika 6: Potenciodinamske krivulje za ~isto Fe, ulito in vro~e valjano,
ulito + `arjeno in vro~e valjano + `arjeno FeMn17 zlitino
A wide variety of biodegradability/corrosion studies
have been made using the electrochemical testing of the
desired designed materials due to aconvenient and easy
method to evaluate the corrosion property by testing
theOCP (open-circuit potential), polarization curves and
EIS(electrochemical impedance spectroscopy) via a
three-electrode system.6–8,12,13,15–17 Very similar results of
corrosion testing with potentiodynamic testing were
observed by H. Hermawan13 with one distinctive diffe-
rence, our pure Fe behaved biodegradability much more
similar to the FeMn17 alloy than theirs.
In the present biodegradable material design study,
the three design criteria as listed by H. Hermawan1 have
been considering fordeveloping a new alloy for bio-
degradability: a) mechanical properties that approach to
those of AISI 316L; b) degradation rate that matches
with the remodelling period (6–12 months) A. Schomig
et al.14 and complete disappearance of the implant mate-
rial within a reasonable period (corrosion rate in the
range of 0.1 mm/year); and c) no toxic substances are re-
leased during the degradation process (just pure Fe and
Mn with minor non-metallic inclusions).
Table 3: Corrosion parameters calculated from the electrochemical
measurements
Tabela 3: Korozijski parametri izra~unani iz elektrokemijskih meritev
Material E(I=0)(mV)
Icorr
(μA)
Rp
(k
)
vcorr
(nm/year)
Pure Fe –675 8.7 4.8 102
Cast –737 40.6 2.4 486
Cast + hot rolled –742 44.7 2.1 537
Cast + annealed –676 18.0 3.7 216
Cast + hot rolled +
annealed –682 14.3 4.3 172
4 CONCLUSIONS
In the present study, we focused on a comparison of
the biodegradability of a produced FeMn17 alloy after
different preparation processing, i.e., casting, hot rolling
and annealing. All the prepared specimens were com-
pared to pure Fe.
With additions of Mn the mechanical properties in-
crease and the corrosion resistance decreases. The pro-
cess parameters influenced the biodegradability as well
as the mechanical properties. The produced material,
cast and hot rolled, has interior stress and that increases
the biodegradability, though the annealing process in-
creases the stability of the material and the corrosion re-
sistance. The microstructure differences of the material
are notable in the grain growth of the annealed sample
and the recrystallization in the annealed hot-rolled sam-
ple.
The electrochemical investigation confirmed the
enhanced biodegradability of the developed FeMn17
alloy compared to pure Fe. We confirmed that cast as
well as additionally hot-rolled specimens exhibited a
pronounced corrosion instability compared to the
annealed material.
Acknowledgement
This work was carried out within the framework of
the Slovene programme P2-0132, "Fizika in kemija
povr{in kovinskih materialov" of the Slovenian Research
Agency, whose support is gratefully acknowledged by
the authors.
5 REFERENCES
1 H. Hermawan, Biodegradable Metals, From Concept to Applications,
1st ed., Springer, Berlin 2012, 69, doi:10.1007/978-3-642-31170-3
2 Y. F. Zheng, X. N. Gu, F. Witte, Biodegradable metals, Materials Sci-
ence and Engineering: R, Reports, 77 (2014) 1–34, doi:10.1016/
j.mser.2014.01.001
3 H. Hermawan, D. Dubé, D. Mantovani, Developments in metallic
biodegradable stents, Acta Biomaterialia, 6 (2010) 5, 1693–1697,
doi:10.1016/j.actbio.2009.10.006
4 H. Hermawan, D. Mantovani, Degradable metallic biomaterials: the
concept, current developments and future directions, Minerva
Biotecnologica, 21 (2009) 4, 207–216
5 M. P. Staiger, A. M. Pietak, J. Huadmai, G. Dias, Magnesium and its
alloys as orthopedic biomaterials: A review, Biomaterials, 27 (2006)
9, 1728–1734, doi:10.1016/j.biomaterials.2005.10.003
6 Z. Zhen, T. Xi, Y. Zheng, A review on in vitro corrosion performance
test of biodegradable metallic materials, Transactions of Nonferrous
Metals Society of China, 23 (2013) 8, 2283–2293, doi:10.1016/
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7 B. Liu, Y. F. Zheng, Effects of alloying elements (Mn, Co, Al, W, Sn,
B, C and S) on biodegradability and in vitro biocompatibility of pure
iron, Acta Biomaterialia, 7 (2011) 3, 1407–1420, doi:10.1016/
j.actbio.2010.11.001
8 D. Vojtìch, J. Kubásek, J. ^apek, I. Pospí{ilová, Comparative me-
chanical and corrosion studies on magnesium, zinc and iron alloys as
biodegradable metals, Mater. Tehnol., 49 (2015) 6, 877–882,
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810 Materiali in tehnologije / Materials and technology 50 (2016) 5, 805–811
Figure 7: Linear polarisation curves for pure Fe, cast, hot rolled, cast
+ annealed and hot rolled + annealed FeMn17 alloy in a simulated
physiological Hank’s solution.
Slika7: Linearne polarizacijske krivulje za ~isto Fe, ulito in vro~e
valjano, ulito + `arjeno in vro~evaljano + `arjeno zlitino FeMn17
9 L. L. Shreir, R. A. Jarman, G.T. Burstein, Corrosion Volume 1,
Metal/Environment Reactions, 3rd ed., Oxford 1994, 3184,
doi:10.1016/B978-0-08-052351-4.50001-7
10 Y. K. Lee, J. H. Jun, C. S. Choi, Damping capacity in Fe-Mn bi- nary
alloys, ISIJ International, 37 (1997) 10, 1023–1030, doi:10.2355/
isijinternational.37.1023
11 J. Martinez, S. M. Cotes, A. F. Cabrera, J. Desimoni, A. Fernández
Guillermet, On the relative fraction of  martensite in -Fe-Mn al-
loys, Materials Science and Engineering A, 408 (2005), 26–32,
doi:10.1016/j.msea.2005.06.019
12 M. Schinhammer, A. C. Hänzi, J. F. Löffler, P. J. Uggowitzer, Design
strategy for biodegradable Fe-based alloys for medical applications,
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13 H. Hermawan, D. Dubé, D. Mantovani, Degradable metallic bio-
materials, Design and development of Fe–Mn alloys for stents, Jour-
nal of Biomedical Materials Research Part A, 93 (2010) 1, 1–11, doi:
10.1002/jbm.a.32224
14 A. Schömig, A. Kastrati, H. Mudra, R. Blasini, H. Schühlen, V.
Klauss, G. Richardt, F. J. Neumann, Four-year experience with
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witzer, Degradation performance of biodegradable Fe-Mn-C(-Pd) al-
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doi:10.1016/j.msec.2012.10.013
16 F. Moszner, Fe–Mn–Pd maraging steels for biodegradable implant
applications, PhD thesis, ETH-Zürich 2014, doi.org/10.3929/
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shape memory alloy as potential biodegradable metallic material,
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Materiali in tehnologije / Materials and technology 50 (2016) 5, 805–811 811

V. BÍLEK et al. EFFECT OF A COMBINATION OF FLY ASH AND SHRINKAGE-REDUCING ADDITIVES ...
813–817
EFFECT OF A COMBINATION OF FLY ASH AND
SHRINKAGE-REDUCING ADDITIVES ON THE PROPERTIES OF
ALKALI-ACTIVATED SLAG-BASED MORTARS
VPLIV KOMBINACIJE LETE^EGA PEPELA IN DODATKA ZA
ZMANJ[ANJE KR^ENJA NA LASTNOSTI MALTE IZ Z
ALKALIJAMI AKTIVIRANE @LINDRE
Vlastimil Bílek, Luká{ Kalina, Jan Koplík, Miroslava Mon~eková,
Radoslav Novotný
Brno University of Technology, Faculty of Chemistry, Materials Research Centre, Purkyòova 118, 612 00 Brno, Czech Republic
xcbilekv@fch.vutbr.cz, bilek@fch.vut.cz
Prejem rokopisa – received: 2015-06-29; sprejem za objavo – accepted for publication: 2015-10-26
doi:10.17222/mit.2015.133
This study is aimed at reducing the drying shrinkage of alkali-activated slag (AAS) through the use of low-calcium fly ash (FA)
and commercially available shrinkage-reducing additives (SRAs) originally developed for OPC-based binders, since there are no
such admixtures tailored for AAS systems. Generally, all the SRAs tested are based on modified alcohols. All the mortars were
based on slag activated by waterglass, with the water-to-slag ratio equal to 0.40 and the sand-to-binder ratio 2:1. In the first step,
the effect of the partial replacement of slag by FA (25, 50 and 75) % of mass fractions on the drying shrinkage and compressive
strength was investigated. On the basis of the obtained results a mortar with 50 % FA in the binder was chosen for subsequent
experiments, where the influence of three types of SRAs on the drying-shrinkage behaviour was examined. It was observed that
while 25 % of the FA did not affect the drying shrinkage significantly, 50 % and 75 % of FA in the binder decreased the drying
shrinkage by 57 % and 78 %, respectively. However, with an increasing content of FA, the compressive strength markedly
decreased. All the tested SRAs had a similar effect on the drying shrinkage of the slag/fly ash (50/50) mortar: at a dose of
0.50 % (by mass of slag) the shrinkage was reduced only slightly, whereas 1.0–3.0 % of SRA resulted in a decrease by 49–66 %.
Also, the drying shrinkage rate during the first days of drying was modified. However, all the SRAs reduced the compressive
strength as compared to the neat slag-FA mortar, especially when the doses were higher than 0.50 %.
Keywords: alkali-activated slag, fly ash, shrinkage reducing additives, shrinkage, strength
Namen {tudije je zmanj{anje kr~enja pri su{enju z alkalijami aktivirane `lindre (AAS) z uporabo lete~ega pepela (FA), z majhno
vsebnostjo kalcija in komercialno dostopnega dodatka za zmanj{anje kr~enja pri su{enju (SRA), originalno razvitega za veziva
na osnovi OPC, ker do sedaj ni bilo takega dodatka primernega za AAS sisteme. Na splo{no so vsi preizku{eni SRA temeljili na
modificiranih alkoholih. Vse malte so bile iz `lindre, aktivirane z vodnim steklom, z razmerjem voda:`lindra (0,40) in
pesek:vezivo v razmerju 2:1. Na prvi stopnji je bil preiskovan vpliv delne nadomestitive `lindre s FA (25, 50 in 75) % masnih
odstotkov, na kr~enje pri su{enju in na tla~no trdnost. Na osnovi dobljenih rezultatov je bila za nadaljevanje preizkusov izbrana
malta s 50 % FA v vezivu, kjer se je preiskoval vpliv treh vrst SRA na skr~ek pri su{enju. Opa`eno je bilo, da medtem ko 25 %
FA ni vplivalo pomembno na skr~ek pri su{enju, je 50 % in 75 % FA v vezivu, zmanj{alo skr~ek pri su{enju za 57 % oziroma
78 %. Vendar pa se je z nara{~ajo~im dele`em FA tla~na trdnost opazno zmanj{ala. Vse preizku{ene SRA so imele podoben
vpliv na kr~enje pri su{enju malte, z razmerjem `lindra:lete~i pepel (50:50), pri koli~ini 0,50 % (masa:`lindra) se je skr~ek
zmanj{al samo delno, medtem ko se je pri 1,0–3,0 % SRA skr~ek zmanj{al za 49–66 %. V prvih dneh se je spremenila tudi
hitrost kr~enja pri su{enju. Vendar pa je SRA zmanj{ala tla~no trdnost, v primerjavi z maltami iz `lindre in FA, {e posebno, ~e
sta bila njuna dele`a ve~ja od 0,5 %.
Klju~ne besede: z alkalijami aktivirana `lindra, lete~i pepel, dodatki za zmanj{anje kr~enja, skr~ek, trdnost
1 INTRODUCTION
Concretes based on Portland cement (PC) provide a
very good mechanical performance for arelatively low
cost. Besides some durability problems that may occur,
PC production is connected with the significant emis-
sions of greenhouse gases such as CO2 and NOx and a
high energy consumption. It is estimated that the PC
industry contributes 7 % to the global CO2 emissions.
The utilization of some industrial wastes or by-products
such as blast-furnace slag (BFS) and fly ash (FA) may be
a possible way to partially solving the above-mentioned
problems. BFS and/or FA can be activated by an alkaline
activator to formulate so-called alkaline cements with simi-
lar or even better properties than PC-based materials.1
Alkali-activated slag (AAS), especially when the
activator is waterglass, shows similar mechanical proper-
ties to PC-based materials.2 However, a high shrinkage,
leading to cracking, usually occurs,3 which is considered
to be the most serious complication for the use of AAS
in practice.4 The fact that the drying shrinkage of AAS
concrete is several times higher thana that of PC-based
concrete is often associated with a higher mesopore
volume in the AAS matrix, which results in substantial
capillary tensile forces at the water-air meniscii and the
material shrinks.5
Materiali in tehnologije / Materials and technology 50 (2016) 5, 813–817 813
UDK 67.017:621.763:621.745.4 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 50(5)813(2016)
One possible way to mitigate shrinkage is the use of
shrinkage-reducing admixtures (SRAs). Generally, SRAs
belong to the group of chemical substances called sur-
factants. The molecules of surfactants are composed of a
polar head (ionic or non-ionic) and a non-polar tail.
Considering the water-vapour interface, the polar head is
attracted to the water phase, while the non-polar tail is
oriented towards the gaseous phase, which means that
surface tension is reduced and thus the capillary stresses
decreas.6 On the other hand, there are several doubts
about the importance of surface tension and capillary
stresses on shrinkage. For example, F. Wittmann7 stated
that even a significant reduction in the surface tension of
the pore solution did not noticeably affect the shrinkage
of a cementitious system and found the disjoining pres-
sure to be more important. Also, M. J. Setzer8 empha-
sized the role of the disjoining pressure and changes in
the surface energy. Nevertheless, SRAs are successfully
used in concrete production and a reduction in shrinkage
of up to 50 % was reported.6
A limited number of studies concerning the use of
SRAs in AAS are available. C. Bilim et al.9 and also M.
Palacios and F. Puertas10 reported a reduction in shrink-
age through the use of an SRA based on polypropylene
glycol, especially in the case of moist curing. T. Bakha-
rev et al.11 achieved a lower shrinkage thanks to some
non-standard SRA and also when using some air-entrain-
ing agent.
2 EXPERIMENTAL WORK
2.1 Materials and sample preparation
Ground granulated BFS with specific surface area of
400 m2/kg was used as a reference binder and siliceous
sand as a fine aggregate. The sand-to-binder ratio was
2:1. The BFS was activated by liquid sodium silicate
with a SiO2/Na2O ratio of 1.85 at the doses of 4.2 % by
binder mass. The water-to-binder ratio was 0.40.
In the first series the BFS was partially replaced by
low-calcium FA. The weight ratios of the S/FA were
100/0 (marked as S), 75/25 (FA25), 50/50 (FA50) and
25/75 (FA75). The mortar FA50 was additionally modi-
fied using three commercially available SRAs in doses of
0.50, 1.0, 2.0 and 3.0 % by mass of BFS + FA. Based on
the safety data sheets, these SRAs were generally a
mixture of alcohols and glycols. The first one (A) was a
mixture of 5-ethyl-1,3-dioxane-5-methanol, 2-ethylpro-
pane-1,3-diol and 2-ethyl-2-(hydroxymethyl)propane-
1,3-diol, the second one (B) was based on
2,2-dimethylpropane-1,3-diol and 2-butylaminoethanol
and the third one (C) consisted mainly of 2-methoxy-
methylenoxy propanol and propane-1,2-diol.
At the beginning of the mixing procedure, all the
liquid components were combined and the BFS + FA
were added. After 30 s of mixing the sand was added.
The total mixing procedure took 4 min. Then the fresh
mortar was cast into the moulds. After 24 h of moist
curing, the specimens were de-moulded and the strength
or shrinkage testing followed.
2.2 Compressive strength testing
In the case of the first series (BFS partially replaced
by FA), the compressive strength was tested on broken
parts of the specimens (from the three-point bending
test) with the dimensions 20 mm × 20 mm × 100 mm,
while the influence of the different dosages of SRAs was
studied on broken parts of the specimens with the dimen-
sions 40 mm × 40 mm × 160 mm. All these tests were
performed for 24 h, 7 d and 28 d after the mixing. After
their de-moulding, all the specimens were submerged in
water until the start of the test.
2.3 Shrinkage tests
The drying shrinkage was measured as length
changes on the basis of the procedure described in
ASTM C596. After the de-moulding the specimens were
immersed in water, in which they were stored for 3 d.
Then they were taken out to the laboratory conditions
(approximately 50 % relative humidity and 23 °C), their
surfaces were dried with a wet towel and measurements
of their length changes began. Also, their mass changes
during drying were measured and evaluated.
3 RESULTS
The results obtained for the first series, i.e., the
influence of the partial replacement of BFS by FA on the
compressive strength development, the drying shrinkage
and the mass changes during drying, are presented in
Figures 1 to 3, respectively. On the basis of these data,
the mortar FA50 with a relatively low drying shrinkage
and a concurrently satisfactory compressive strength was
V. BÍLEK et al. EFFECT OF A COMBINATION OF FLY ASH AND SHRINKAGE-REDUCING ADDITIVES ...
814 Materiali in tehnologije / Materials and technology 50 (2016) 5, 813–817
Figure 1: Effect of fly ash on the compressive strength of alkali-acti-
vated slag mortar
Slika 1: Vpliv lete~ega pepela na tla~no trdnost malte iz z alkalijami
aktivirane `lindre
chosen for further experiments, where the effect of
different dosages of different SRAs on the compressive
strength (Table 1), the drying shrinkage (Figure 3) and
the mass loss during drying (Figure 4) was investigated.
In order to compare the compressive strengths of the
mortars containing SRAs with the reference mortar FA50
by using the specimens of the same dimensions, the
FA50 mortar was prepared once more. Thus, the desig-
nation R for the FA50 mortar was used in the second part
of this study.
4 DISCUSSION
As can be seen from Figures 1 and 2, the partial
replacement of the BFS in the binder by FA led to a
gradual decrease in the compressive strength, but also to
a significant shrinkage reduction in the case of mortars
with 50 % and 75 % of FA in the binder. These results
correspond with those reported by M. Marjanovi} et al.12,
where the compressive strength of mortars with a pre-
dominant content of BFS in the binder were significantly
higher than those with a predominant content of FA.
Also, the slightly higher shrinkage of the mortar with a
BFS/FA ratio of 75/25 than that of the 100/0 mortar was
recorded in the same study. The increasing content of FA
led to a substantial increase in the mass loss during the
drying process under laboratory conditions, as can be
seen from Figure 3. This could imply that a smaller
quantity of hydration products was formed in the speci-
mens with higher doses of FA, because the alkali activa-
tion of FA is generally favoured by an elevated tempe-
rature and a larger amount of Na2O in the mixture.12
Moreover, the main hydration product of FA activation,
the N-A-S-(H) phase, does not chemically bind water,13
which may also affect the mass changes during drying.
Table 1 shows the impact of SRAs on the compres-
sive strength on the reference mortar R, whose compo-
sition was the same as the mortar FA50. The difference
was only in the size of the specimens for compressive
strength testing, as stated in Section 2.2. However, simi-
lar results for the compressive strength were obtained,
when compared Table 1 and Figure 2. Regardless of the
type of SRA used, the compressive strength of the refe-
rence mortar was markedly reduced when SRAs were
applied. Generally, the compressive strength decreased
with the increasing dose of SRAs. The exception was the
mortar with 3 % of SRA C, where a noticeably higher
strength than that of the mortars with 1 % and 2 % of
this additive were observed during the whole testing
period.
The effect of SRAs on drying shrinkage of the
reference mortar is shown in Figure 4. It is clear that all
the SRAs tested modified the drying shrinkage develop-
ment in a similar way. If only 0.5 % of SRA was used,
the drying shrinkage slightly decreased, while a signifi-
cant decrease in shrinkage was observed when the dose
of SRA was raised from 0.5 % to 1.0 %. A further
increase in the SRA dosage only led to a minor decrease
in the drying shrinkage. A noticeably different situation
was observed for the mass loss during the drying of these
mortars (Figure 5), which exhibited a rather gradual
V. BÍLEK et al. EFFECT OF A COMBINATION OF FLY ASH AND SHRINKAGE-REDUCING ADDITIVES ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 813–817 815
Figure 3: Effect of fly ash on the mass changes during drying of
alkali-activated slag mortar
Slika 3: Vpliv lete~ega pepela na zmanj{anje mase med su{enjem
malte iz z alkalijami aktivirane `lindre
Figure 2: Effect of fly ash on the drying shrinkage of alkali-activated
slag mortar
Slika 2: Vpliv lete~ega pepela na skr~ek pri su{enju malte iz z alkali-
jami aktivirane `lindre
Table 1: Effect of SRAs (A, B, C) on the compressive strength development (in MPa) of alkali-activated BFS/FA 50/50 (R) mortar
Tabela 1: Vpliv SRA (A, B, C) na razvoj tla~ne trdnosti malte BFS/FA 50/50 (R), aktivirane z alkalijami
Mixture R A B C
SRA dose (%) 0 0.5 1.0 2.0 3.0 0.5 1.0 2.0 3.0 0.5 1.0 2.0 3.0
1 d 8.1 4.4 5.0 – – 4.7 3.0 – – 4.1 3.8 1.2 5.7
7 d 16 9.9 9.1 2.7 1.5 11 5.7 2.4 1.9 7.9 5.7 4.0 11
28 d 54 40 16 3.6 2.8 44 19 2.2 2.2 40 7.6 4.2 14
increase in the mass loss with the increasing content of
SRAs. The reason for such an increase in the mass loss is
likely to be associated with some retardation effect of all
the SRAs used on the hydration of alkali-activated
BFS/FA mortars, which is indicated by its compressive
strength development. The SRAs also modified the
shrinkage rate during the first days of drying, where a
significant reduction as compared with the reference
mortar was observed. After this the period of most
intensive shrinkage followed, while since approximately
the 10th day of exposure to a dry climate, all the mortars
had almost the same shrinkage profile.
Although its porosity was not measured, we can
expect a significant decrease in the amount of hydration
products of the mortars containing 1.0 % and more SRA
and therefore a higher volume of larger pores filled with
easily evaporable water before the start of the drying.
This raises considerable doubts about the compatibility
V. BÍLEK et al. EFFECT OF A COMBINATION OF FLY ASH AND SHRINKAGE-REDUCING ADDITIVES ...
816 Materiali in tehnologije / Materials and technology 50 (2016) 5, 813–817
Figure 5: Effect of SRAs (A, B, C) on the mass loss of alkali-acti-
vated BFS/FA 50/50 (R) mortar during drying
Slika 5: Vpliv SRA (A, B, C) na zmanj{anje mase pri su{enju malte
BFS/FA 50/50 (R), aktivirane z alkalijami
Figure 4: Effect of SRAs (A, B, C) on the drying shrinkage of alkali-
activated BFS/FA 50/50 (R) mortar
Slika 4: Vpliv SRA (A, B, C) na kr~enje pri su{enju malte BFS/FA
50/50 (R), aktivirane z alkalijami
of the tested SRAs with the studied alkali-activated sys-
tem. Some of the authors mentioned in the introduction
achieved a significant reduction of the drying shrinkage
without any significant changes in the AAS hydration,
but they used SRAs based on polypropylene glycol,
while the SRAs used in this study mainly comprised
low-molecular-weight diols.
5 CONCLUSIONS
This paper looks at the possibilities of drying-shrink-
age reduction for mortars based on AAS by FA and
SRAs. The compressive strength and mass changes
during the drying were also evaluated. From the results
obtained it can be concluded that despite a certain
strength decrease the FA can be successfully used for the
shrinkage reduction of AAS. Considering both the
compressive strength and the drying shrinkage, the
mortar with the ratio of BFS/FA 50/50 seems to be the
most favourable FA. A further decrease in the shrinkage
of this mortar was achieved through the application of
commercial SRAs based mainly on diols and designed
for Portland-cement-based systems. However, all these
SRAs had a very negative impact on the compressive-
strength development, which makes them unsuitable for
the studied alkali-activated BFS/FA system.
Acknowledgement
This outcome was achieved with the financial support
by the project: Materials Research Centre at FCH BUT-
Sustainability and Development, REG LO1211, with
financial support from National Programme for Sustain-
ability I (Ministry of Education, Youth and Sports).
6 REFERENCES
1 A. Férnandez-Jiménez, A. Palomo, D. Revuelta, Alkali activation of
industrial by-products to develop new earth-friendly cements, Proc.
of the 11th Inter. Conf. on Non-Conventional Materials and Tech-
nologies Bath, UK 2009, 1–15
2 F. Puertas, M. Palacios, H. Manzano, J. S. Dolado, A. Rico, J. Rod-
ríguez, A Model for the C-A-S-H gel Formed in Alkali-Activated
Slag Cements, Journal of the European Ceramic Society, 31 (2011),
2043–2056, doi:10.1016/j.jeurceramsoc.2011.04.036
3 F. G. Collins, J. G. Sanjayan, Cracking tendency of alkali-activated
slag concrete subjected to restrained shrinkage, Cement and Concrete
Research, 30 (2000), 791–798, doi:10.1016/S0008-8846(00)00243-X
4 A. A. M. Neto, M. A. Cincotto, W. Repette, Drying and autogenous
shrinkage of pastes and mortars with activated slag cement, Cement
and Concrete Research, 38 (2008), 565–574. doi:10.1016/
j.cemconres.2007.11.002
5 F. Collins, J. G. Sanjayan, Effect of pore size distribution on drying
shrinking of alkali-activated slag concrete, Cement and Concrete
Research, 30 (2000), 1401–1406, doi:10.1016/s0008-8846(00)
00327-6
6 F. Rajabipour, G. Sant, W. J. Weiss, Interactions Between Shrinkage
Reducing Admixtures (SRA) and Cement Paste’s Pore Solution,
Cement and Concrete Research, 38 (2008), 606–615, doi:10.1016/
j.cemconres.2007.12.005
7 F. Wittmann, Heresies On Shrinkage And Creep Mechanisms (Key-
note Lecture), Proc. of the Eighth International Conference on Creep,
Shrinkage and Durability of Concrete and Concrete Structures
(CONCREEP 8), London 2009, 3–10, doi:10.1201/
9780203882955.pt1
8 M. J. Setzer, The solid-liquid gel-system of hardened cement paste,
Proc. of the Eighth International Conference on Creep, Shrinkage
and Durability of Concrete and Concrete Structures (CONCREEP 8),
London 2009, 237–243, doi:10.1201/9780203882955.ch28
9 C. Bilim, O. Karahan, C. D. Atiº, S. Ýlkentapar, Influence of admix-
tures on the properties of alkali-activated slag mortars subjected to
different curing conditions, Materials and Design, 44 (2013),
540–547, doi:10.1016/j.matdes.2012.08.049
10 M. Palacios, F. Puertas, Effect of shrinkage-reducing admixtures on
the properties of alkali-activated slag mortars and pastes, Cement and
Concrete Research, 37 (2007), 691–702, doi:10.1016/j.cemconres.
2006.11.021
11 T. Bakharev, J. G. Sanjayan, Y. B. Cheng, Effect of Admixtures on
Properties of Alkali-Activated Slag Concrete, Cement and Concrete
Research, 30 (2000), 1367–1374, doi:10.1016/S0008-8846(00)
00349-5
12 N. Marjanovi}, M. Komljenovi}, Z. Ba{~arevi}, V. Nikoli}, R. Petro-
vi}, Physical-mechanical and microstructural properties of alkali-
activated fly ash–blast furnace slag blends, Ceramics International,
41 (2015), 1421–1435, doi:10.1016/j.ceramint.2014.09.075
13 J. L. Provis, R. J. Myers, C. White, V. Rose, J. S. J. Van Deventer,
X-ray Microtomography Shows Pore Structure and Tortuosity in
Alkali-Activated Binders, Cement and Concrete Research, 42 (2012),
855–864, doi:10.1016/j.cemconres.2012.03.004
V. BÍLEK et al. EFFECT OF A COMBINATION OF FLY ASH AND SHRINKAGE-REDUCING ADDITIVES ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 813–817 817

O. BÍLEK et al.: CUTTING-TOOL PERFORMANCE IN THE END MILLING OF CARBON-FIBER-REINFORCED PLASTICS
819–822
CUTTING-TOOL PERFORMANCE IN THE END MILLING OF
CARBON-FIBER-REINFORCED PLASTICS
ZMOGLJIVOST REZILNEGA ORODJA PRI REZKANJU PLASTIKE,
OJA^ANE Z OGLJIKOVIMI VLAKNI
Ondøej Bílek, Soòa Rusnáková, Milan @aludek
Faculty of Technology, TBU Zlín, Nám. T.G.M 5555, 760 01 Zlín, Czech Republic
bilek@ft.utb.cz
Prejem rokopisa – received: 2015-06-30; sprejem za objavo – accepted for publication: 2015-09-09
doi:10.17222/mit.2015.153
Carbon-fiber-reinforced plastics (CFRPs) are no longer the exclusive domain of aerospace industries and nowadays find more
applications in the automotive and consumer industries. The growing volume of these materials increases the need for their
efficient machining. However, the machinability of CFRPs is a rather complex task due to the heterogeneity and the consider-
able number of parameters influencing the cutting process. The cutting tool has long been recognized as an important factor
influencing both surface quality and dimensional accuracy. Hence, in this paper, tool performance is investigated in the end
milling of CFRPs. The effect of tool geometry and coatings on the cutting forces, dimensional accuracy and surface quality were
experimentally examined for side and slot milling.
Keywords: CFRP, composites, end mill, surface quality, cutting forces
Plastika, oja~ana z ogljikovimi vlakni (CFRP), dosedaj ekskluzivna domena letalske industrije, se v zadnjem ~asu bolj pogosto
uporablja tudi v strojni{tvu, v avtomobilski industriji in v potro{ni{tvu. Z nara{~anjem uporabe teh materialov, nara{~a tudi
potreba po u~inkoviti strojni obdelavi. Vendar pa je strojna obdelava CFRP-a precej kompleksna naloga zaradi heterogenosti in
{tevilnih parametrov, ki vplivajo na proces rezanja. @e dolgo je znano, da je rezilno orodje pomemben faktor, ki vpliva na
kvaliteto povr{ine in dimenzijsko natan~nost. Zato je v ~lanku preiskovana zmogljivost orodja pri rezkanju plastike, oja~ane z
ogljikovimi vlakni. Pri bo~nem in utorovnem rezkanju je bil preiskovan vpliv geometrije orodja in povr{inske prevleke na sile
rezanja, dimenzijsko zanesljivost in na kvaliteto povr{ine.
Klju~ne besede: CFRP, kompoziti, ~elni rezkar, kvaliteta povr{ine, sile rezanja
1 INTRODUCTION
Milling is the key manufacturing operation for the
successful machining of fiber-reinforced plastic parts,
which ensures dimensional and qualitative requirements.
Milling is used, as a rule, as an edge-trimming operation,
or a slotting/routing process to produce complex con-
tours.
Composite materials such as carbon fiber-reinforced
plastics (CFRPs) are heterogeneous materials composed
of dissimilar constituents, where the resulting mechani-
cal properties are more than the sum of the characte-
ristics of the individual components. A plastic matrix of
CFRPs is reinforced by carbon fibers entirely, resulting
in a high specific strength and stiffness throughout the
lifetime of the components, together with electrical and
magnetic resistance. This makes them interesting for
applications in the automotive and aircraft industries, in
railway and ship building, medical applications and the
space industry.1,2
CFRP composites replace traditional metallic
materials; however, their machinability is fundamentally
different and the cutting mechanism is still under inve-
stigation. The machining properties are influenced by the
heterogeneity and anisotropy of the composite material.
The mechanical properties of the matrix are strongly
Materiali in tehnologije / Materials and technology 50 (2016) 5, 819–822 819
UDK 620.168:621.926:678.077 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 50(5)819(2016)
Table 1: Characteristics of the end mill
Tabela 1: Zna~ilnosti steblastih rezkarjev
Tool cutting zone Cutting diameter (mm) Helix angle (°) Rake angle (°) Coating
T1 Down-cut spiral with chip-breaker 5.81±0.07 15 0 PCD
T2 Four flute up-cut spiral 5.96±0.02 10 6 PCD
T3 Cross-pitch spiral 5.92±0.01 25 and 27 7 PCD
T4 Cross-pitch spiral 5.91±0.02 25 and 27 7 uncoated
T5 Up-cut spiral with chip-breaker 5.96±0.03 47 10 Multilayered
T6 Right- and left-hand spiral 5.94±0.02 13 30 Multilayered
T7 Cross-pitch spiral 5.97±0.02 10 0 PCD
T8 Three flute upcut spiral 5.91±0.03 15 30 PCD
temperature dependent, while the carbon fibers can
withstand temperatures up to 3000 °C, and during
machining cause considerable tool wear.3
Machining CFRPs brings with it certain difficulties,
such as layer delaminations, fiber pull-out, uncut fibers,
degradation and burning of the plastic matrix.4–7 Tool
manufacturers and scientists recommend carbide tools
with a polycrystalline diamond (PCD) coating for CFRP
material machining. These tools ensure high productivity
with a sufficient tool life. Unfortunately, the cost of PCD
tools is six times higher than uncoated tools, whereas
tool life is only three times longer.8 End mills usually
have a very special design suitable for the milling of
CFRPs.9 Helical end mills with two flutes are suitable for
rigid or back-supported parts.10 Compression end mills
with two sets of left-hand and right-hand flutes limit the
delamination and burrs on the cut edges.11 Furthermore,
the machining of flexible parts by multi-tooth end mills
eliminates the cutting force along the Z-axis, thereby re-
ducing the deflection and vibration of the part.8 In-bet-
ween are helical mills with a chip-breaker geometry12
that shears fibers and shortens chips with the simulta-
neous reduction of vibrations.
In this work we compare the performance of end
mills on the resulting surface quality, dimensional accu-
racy and cutting forces in the slotting and side milling of
CFRPs.
2 EXPERIMENTAL PROCEDURE
The composite samples used in the experimental
investigation were produced by vacuum infusion techno-
logy. The carbon fiber and the epoxy were supplied by
KORDCARBON CZ. The fabric was Toray HS 3K 200
tex, a balanced twill bidirectional (0–90°) weave and
380 μm thick. The epoxy used in the fabrication was the
commercially available Havel L285 epoxy system.
Fourteen layers of fabric were used to fabricate the
work samples with an average thickness of 4 mm. Over-
all, the CFRP composite average flexure strength was
501 MPa, the modulus of elasticity was 43,300 GPa and
the tensile strength was 378 MPa.
The experiment was carried out by using eight car-
bide end mills acquired from different producers (SECO,
WNT, KTOOLS). The required tool cutting diameter
was 6 mm. In this article the end mills are referred to as
T1-T8 and were concerned for special applications such
as the milling of composites with a wide range of com-
positions. Their geometry is summarized in Table 1.
Except for the T4 tool, the end mills were coated to with-
stand the machining of a highly abrasive work material.
A CNC milling machine C-442 HWT with a 1.0-kW
spindle power and high-speed spindle of 416.16 Hz was
used to perform the experiments. Part of the program
was created to ensure the repeatability and accuracy of
the experiment. The machining was divided into two
consecutive methods, first slot milling (Figure 1a) and
next upcut side milling (Figure 1b) of the CFRP work
material. The machining conditions are given in Table 2
for both methods. A coolant was not used and the tool
cutting length was 150 mm in both methods. A tool-wear
evaluation was carried out at the end of every set of
experiments. First, each tool was milling a slot and
subsequently a side of work material, so that the tool dis-
placement towards the bottom surface of the work
material was 2 mm.
Table 2: Cutting conditions
Tabela 2: Pogoji rezanja
Method
Cutting
speed
vc (m/min)
Feed rate
vf (mm/min)
Depth of
cut
ap (mm)
Width of
cut ae (mm)
Slot milling 100 550 1.25 ae = D
Side milling 100 550 2.0 4.0
Cutting forces were measured on the strain-gauge
dynamometer in the direction of the feed-rate vector
(force Ff) and perpendicularly to Ff (force Ffn). The
signal from dynamometer was processed by a Spider8
data-acquisition system and the evaluations were made
through a PC with Conmes Spider software, as illus-
trated in Figure 2.
The dimensional accuracy of the machined slot was
evaluated using a Mitutoyo caliper. The surface quality
was measured with a portable surface-roughness tester
Mitutoyo SJ-301, according to ISO 4287 on the high-
O. BÍLEK et al.: CUTTING-TOOL PERFORMANCE IN THE END MILLING OF CARBON-FIBER-REINFORCED PLASTICS
820 Materiali in tehnologije / Materials and technology 50 (2016) 5, 819–822
Figure 1: a) slot milling and b) side milling
Slika 1: a) rezkanje utora in b) bo~no rezkanje
lighted areas at the bottom of the slot (Figure 1a) and
the side surface (Figure 1b). The measurements repeated
were 15 times at different surface locations in the
direction of the feed-rate vector that was identified as the
direction of highest roughness. The roughness parameter
Ra was considered as an evaluation parameter for the
surface quality according to ISO.
3 RESULTS AND DISCUSSION
3.1 Cutting forces
All the force measurements were worth a maximum
of 4.86 % of variance. The results were graphically
summarized in Figure 3 and 4. There is a significant
difference between the end mills, moreover not the only
one is suitable both as side and slot milling. Tool T6 with
right- and left-hand spirals reached the lowest values Ff
when slot milling (Figure 3) and the second best value
as for T4 tool without coating. Higher values of Ff and
Ffn were achieved by the same T3 tool with the coating.
This may be due to increased adhesion of work material
to the PCD coating material.
Tools T1, T3, T8 keep along the average. Conversely,
the highest values of Ff achieved tool T2 and T7,
although with different flute geometries, but with the
same helix angle of 10°, despite the normal force Ffn
was the lowest one. The helix angle probably improves
the chip removal at the bottom of the slot; however, the
tool is more power loaded in the feed direction. In the
case of higher cutting depths we can expect a higher tool
bend, but without a buckling effect.
When side milling, the non-significant Ffn force was
not evaluated, yet was mentioned in Figure 4. The force
Ff is more than 2 times higher than when slot milling.
The smallest Ff value comes from the T4 tool without
coating, the worst was the T6 tool with a right- and
left-hand hand spiral. The reason for this may be in the
compression effect of contra-rotating spirals, becoming
for these tools. Although the flutes’ geometry for the T1,
T2, T5 and T8 tools was dissimilar, the feed-force cha-
racteristics were almost identical and quite balanced.
3.2 Surface roughness
The dynamometric investigation was followed by
surface-roughness measurements of the Ra parameter.
O. BÍLEK et al.: CUTTING-TOOL PERFORMANCE IN THE END MILLING OF CARBON-FIBER-REINFORCED PLASTICS
Materiali in tehnologije / Materials and technology 50 (2016) 5, 819–822 821
Figure 4: Cutting forces for side milling of CFRP work material
Slika 4: Sile rezanja pri bo~nem rezkanju CFRP obdelovanca
Figure 2: Experimental set up for milling CFRPs
Slika 2: Eksperimentalni sestav za rezkanje CFRP
Figure 5: Surface roughness Ra for slot and side milling
Slika 5: Hrapavost povr{ine Ra pri rezkanju utora in pri bo~nem
rezkanju
Figure 3: Cutting forces for slot milling of CFRP work material
Slika 3: Sile rezanja pri rezkanju utora v CFRP obdelovanec
The results are shown in Figure 5 for the side and slot
milling.
The lowest average value of Ra = 1.38±0.40 μm was
on the side surface made by the uncoated T4 tool and the
same value Ra = 1.38±0.29 μm is at the bottom of the slot
by the T8 tool. The T5 tool had the highest and most
unsatisfactory effect for side milling, where the rough-
ness value was Ra = 6.19±1.57 μm. It is the only tool
with an upcut spiral in the experiment and is actually a
production tool recommended for rough machining.
Approximately half values of Ra = 3.45±1.30 μm, but
highest for slot milling, was made with the cross-pitched
T7 tool with a 0° rake angle.
3.3 Dimensional accuracy
Since end mills should have a 6 mm diameter, it was
expected that the resulting dimension of the slot would
be given with a certain accuracy close to that value. As it
turned out, this was not entirely true for slot milling of
CFRP work material, as can be seen in Figure 6.
In most cases, the resulting dimension of the slot is
smaller than the desired nominal size. On the one hand,
it could be given by the true tool diameter (Table 1) that
was smaller than the producer specification. In contrst,
for the T5 tool with an upcut spiral we found that the slot
was about 0.13 mm wider. This excludes the eventuality
that a smaller tool diameter was made on purpose, taking
into account the CFRP material’s behavior. Finally, the
T8 tool made it possible to achieve a superior slot
accuracy, but otherwise lagged behind in terms of cutting
efficiency and surface quality.
4 CONCLUSIONS
In this work the cutting performance of the end mills
in the CFRPs milling has been presented, taking into
account the force, surface roughness and dimensional
measurements. After the testing of eight end mills seve-
ral conclusions can be drawn:
• an upcut spiral can cause considerably worse surface
quality together with an increment of the slot dimen-
sion,
• PCD-coated tools for CFRPs machining has a lower
cutting efficiency due to a different friction charac-
teristic than the uncoated tools,
• tools with left-hand and right-hand spirals recom-
mended for side milling, result in up to twice the
cutting force,
• to ensure compliance with the required dimensions,
accuracy and stability it is necessary to include an
offset (by part program, cutting conditions, thermal
field) for the CFRPs tool individually.
5 REFERENCES
1 E. Uhlmann, F. Sammler, S. Richarz, F. Heitmüller, M. Bilz,
Machining of Carbon Fibre Reinforced Plastics, Procedia CIRP, 24
(2014), 19–24, doi:10.1016/j.procir.2014.07.135
2 S. Rusnakova, L. Fojtl, M. Zaludek, V. Rusnak, Design of material
composition and technology verification for composite front end
cabs, Manufacturing Technology, 14 (2014), 607–611
3 O. Pecat, R. Rentsch, E. Brinksmeier, Influence of milling process
parameters on the surface integrity of CFRP, Procedia CIRP, 1
(2012), 466–470, doi:10.1016/j.procir.2012.04.083
4 M. H. El-Hofy, S. L. Sooa, D. K. Aspinwalla, W. M. Simb, D. Pear-
sonc, P. Hardend, Factors affecting workpiece surface integrity in
slotting of CFRP, Procedia Eng., 19 (2011), 94–99, doi:10.1016/
j.proeng.2011.11.085
5 W. Hintze Wolfgang, D. Hartmann, C. Schütte, Occurrence and
propagation of delamination during the machining of carbon fibre
reinforced plastics (CFRPs) - An experimental study, Compos. Sci.
Technol., 71 (2011), 1719–1726, doi:10.1016/j.compscitech.
2011.08.002
6 Y. Karpat, O. Bahtiyar, B. Deer, Mechanistic force modeling for
milling of unidirectional carbon fiber reinforced polymer laminates,
Int. J. Mach. Tools Manuf., 56 (2012), 79–93, doi:10.1016/
j.ijmachtools.2012.01.001
7 R. Teti, Machining of Composite Materials, CIRP Annals – Manu-
facturing Technology, 51 (2002), 611–634, doi:10.1016/S0007-
8506(07)61703-X
8 L. N. López de Lacalle, A. Lamikiz, F. J. Campa, A. Fdz.
Valdivielso, I. Etxeberria, Design and Test of a Multitooth Tool for
CFRP Milling, J. Compos. Mater., 43 (2009), 3275–3290,
doi:10.1177/0021998309345354
9 Y. G. Wang, X. P. Yan, X. G. Chen, C. Y. Sun, G. Liu, Cutting Per-
formance of Carbon Fiber Reinforced Plastics Using PCD Tool,
Advanced Materials Research, 215 (2011), 14–18, doi:10.4028/
www.scientific.net/AMR.215.14
10 J. P. Davim, P. Reis, Damage and dimensional precision on milling
carbon fiber-reinforced plastics using design experiments, J. Mater.
Process. Technol., 160 (2005), 160–167, doi:10.1016/j.jmatprotec.
2004.06.003
11 P. Masek, P. Zeman, Effective milling of composites with thermo-
plastic matrix, MM Prumyslove spektrum, 6 (2013), 60–64
12 S. Legutko, G. Krolczyk, J. Krolczyk, Quality evaluation of surface
layer in highly accurate manufacturing, Manuf. Technol., 14 (2014),
50–56
O. BÍLEK et al.: CUTTING-TOOL PERFORMANCE IN THE END MILLING OF CARBON-FIBER-REINFORCED PLASTICS
822 Materiali in tehnologije / Materials and technology 50 (2016) 5, 819–822
Figure 6: Dimensional accuracy of the width of the slot
Slika 6: Dimenzijska natan~nost {irine utora
M. DOŒPIA£, M. NABIA£EK: INFLUENCE OF SOLIDIFICATION SPEED ON THE STRUCTURE AND MAGNETIC ...
823–827
INFLUENCE OF SOLIDIFICATION SPEED ON THE STRUCTURE
AND MAGNETIC PROPERTIES OF Nd10Fe81Zr1B6 IN THE
AS-CAST STATE
VPLIV HITROSTI STRJEVANJA NA STRUKTURO IN MAGNETNE
LASTNOSTI ZLITINE Nd10Fe81Zr1B6 V LITEM STANJU
Marcin Doœpia³, Marcin Nabia³ek
Czestochowa University of Technology, Institute of Physics, Armii Krajowej Av. 19, 42-200 Czestochowa, Poland
mdospial@wp.pl
Prejem rokopisa – received: 2015-07-01; sprejem za objavo – accepted for publication: 2015-09-15
doi:10.17222/mit.2015.174
The paper presents results of the structure and magnetic properties of the Nd10–xTbxFe81Zr1B6 (x = 0, 2) alloy in the as-cast state.
The samples were produced using the melt-spinning method. The solidification speed was controlled indirectly by changing the
linear velocity of a copper drum. Based on magnetic and structural studies, it was found that samples obtained while the linear
velocity of the copper drum was equal to 20 m/s had good hard-magnetic properties. The substitution of 2 % of Nd by Tb led to
grain growth of both the -Fe (9 nm and 24 nm for x = 0 and x = 2) and Nd2Fe14B phases (41 nm and 70 nm for x = 0 and x = 2,
respectively). The grain growth to sizes higher than the exchange interaction distance in the sample with Tb resulted in a
bimodal shape of the demagnetization curve. The samples obtained at higher linear speeds of the copper drum were composed
of amorphous matrices with small amounts of crystalline phases and had weak, soft-magnetic properties.
Keywords: permanent magnets, nanocomposities, X-ray diffraction, exchange interactions
^lanek predstavlja rezultate {tudije strukturnih in magnetnih lastnosti zlitine Nd10–xTbxFe81Zr1B6 (x = 0, 2) v litem stanju. Vzorci
so bili pripravljeni z ulivanjem na hitro vrte~em se valju. Hitrost strjevanja je bila posredno kontrolirana s spreminjanjem
linearne hitrosti bakrenega valja. Magnetne in strukturne {tudije so pokazale, da imajo vzorci, dobljeni pri linearni hitrosti
bakrenega valja 20 m/s, dobre trdomagnetne lastnosti. Nadomestilo z 2 % atomskega dele`a Nd s Tb, je povzro~ilo rast zrn obeh
faz -Fe (9 nm in 24 nm pri x = 0 in x = 2) in Nd2Fe14B (41 nm in 70 nm pri x = 0 in x = 2). Rast zrn, do velikosti, ve~je od
izmenjalne interakcijske razdalje v vzorcu s Tb, se je pokazala v bimodalni obliki demagnetizacijske krivulje. Vzorci, dobljeni
pri ve~jih linearnih hitrostih bakrenega valja, so bili sestavljeni iz amorfne osnove z majhnim dele`em kristalnih faz in so imeli
slabe mehko magnetne lastnosti.
Klju~ne besede: permanentni magneti, nanokompoziti, rentgenska difrakcija, izmenjalne interakcije
1 INTRODUCTION
Alloys based on rare earths, iron and boron are
currently some of the most popular groups of materials
with hard magnetic properties. These materials are
widely used in the electro-engineering industry. In recent
years, many scientific and industrial units have been
carrying out research on methods for curing their mag-
netic properties.1–5 The studies were mainly focused on
slight modifications to the chemical composition and
modifications to the production process parameters.
The most effective process, which guarantees repeat-
ability of the functional properties of produced alloys, is
producing amorphous ribbons and then tailoring their
properties by an appropriately matched heat treatment.6
However, the application of the heat treatment generates
additional costs. An alternative approach is solidification
of the liquid alloy with a slow speed. This leads to a
material with hard magnetic properties in the as-cast
state. A drawback of this method is the poor repeatability
of the functional parameters.
Modifications to the composition by a partial sub-
stitution of Nd by Tb can lead to two outcomes: if an
excess amount of Tb is used, it can result in the for-
mation of an amorphous matrix dividing the Nd2Fe14B
grains, thereby leading to a higher coercivity and also a
significantly reduced remanence;7,8 if smaller amounts of
rare earths with respect to the stoichiometric Nd2Fe14B
composition are used, it may result in the formation of
nanocomposites consisting of both Nd2Fe14B and -Fe
grains. The proper selection of the production process
parameters leads to grains of soft magnetic phase with
sizes similar or smaller than the exchange-interaction
length. This allows obtaining hard magnetic materials
with high saturation magnetization and a significant
remanence.8,9
Motivated by the above idea, we studied the influ-
ence of cooling rate on the curing of magnetic properties
and the structure of Nd10–xTbxFe81Zr1B6 (x = 0 and x = 2)
alloys, in the as-cast state, and the results of the study we
report in this paper.
2 EXPERIMENTAL DETAILS
The research material, i.e., Nd10–xTbxFe81Zr1B6 (x = 0,
2) melt, was obtained from components of the following
Materiali in tehnologije / Materials and technology 50 (2016) 5, 823–827 823
UDK 620.1:66.017:621.78 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 50(5)823(2016)
purity: Fe – 99.99 %; Zr – 99.99 %; Nd – 99.95 %, Tb –
99.95 %, B – 99.99 %. The crystalline ingots were seve-
ral times remelted to ensure the best mixing of elements
constituting the investigated material. Thin tapes were
produced by a unidirectional solidification of liquid
metal on a rotating copper cylinder, with three different
linear velocities (20, 30 and 35) m/s. The microstructure
of the samples, in the as-cast state, was examined using a
Bruker D8 Advance X-ray diffractometer with a charac-
teristic Cu-K radiation source (0.154056 nm). The
samples were scanned in the 2 angle range from 30° to
120° with a resolution of 0.02° and an exposure time of
3 s per step. Two samples were chosen for further analy-
ses, both qualitative and quantitative. Phase-composition
analysis was performed using the Rietveld method. The
magnetic studies were performed using a LakeShore
vibrating-sample magnetometer in an external magnetic
field up to 2 T. The demagnetization coefficient depen-
dent on the shape of the sample was not taken into
account. All the measurements were performed at room
temperature.
3 RESULTS AND DISCUSSION
Figure 1 shows the X-ray diffraction patterns ob-
tained for Nd10–xTbxFe81Zr1B6 (x = 0, 2) samples in the
form of tapes, in the as-cast state.
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824 Materiali in tehnologije / Materials and technology 50 (2016) 5, 823–827
Figure 1: X-ray diffraction patterns of: a) Nd10Fe81Zr1B6, b) Nd8Tb2Fe81Zr1B6 alloys obtained at different cooling rates, for samples in the
as-cast state
Slika 1: Rentgenska difrakcija: a) Nd10Fe81Zr1B6, b) Nd8Tb2Fe81Zr1B6 zlitin, dobljenih pri razli~nih hitrostih ohlajanja, za vzorce v litem stanju
Figure 2: Quantitative match of the phase composition obtained using the Rietveld refinement method for: a), c) Nd10Fe81Zr1B6, d), f)
Nd8Tb2Fe81Zr1B6 alloys solidified at 20 m/s linear velocity of copper drum, where a), d) experimental curve, b), e) matched pattern and c), f)
difference plot
Slika 2: Kvantitativno ujemanje sestave faz, dobljene z Rietveldovo metodo izpopolnitve: a), c) Nd10Fe81Zr1B6, d), f) Nd8Tb2Fe81Zr1B6 zlitini
strjeni pri linearni hitrosti bakrenega valja 20 m/s, kjer sta a), d) eksperimentalni krivulji, b), e) ujemanje in c), f) diagram razlik
In all the diffraction patterns obtained for the samples
cooled at higher rates, corresponding to the drum linear
speeds of 30 m/s and 35 m/s, at a 2 angle of about 43° a
broad, diffused halo characteristic for materials contain-
ing an iron-based amorphous phase is present. Also, it is
possible to distinguish the presence of narrower peaks
originating from small inclusions of crystalline phases.
Similar patterns were described in the literature as
semi-crystalline structures.10,11
The diffraction patterns of the samples solidified at a
drum speed of 20 m/s were composed of a number of
narrow peaks characteristic for a crystalline structure. In
this case, no diffused halo was observed, which indicates
the absence of any amorphous phase. A comparison of
the experimental diffraction peaks with the EVA data-
base let us conclude that two crystalline phases were
present in the sample: -Fe and Nd2Fe14B. These results
are in agreement with other reports concerned on
Nd2Fe14B alloys with an over-stoichiometric iron con-
tent.7,12–13
To ascertain quantitatively the phase compositions of
the samples without an amorphous matrix, Rietveld
analysis was performed (Figure 2), and the matching
results are summarized in Tables 1 and 2.
Table 1: The qualitative phase composition of Nd10–xTbxFe81Zr1B6
(where x = 0 or 2) samples solidified with the linear velocity of the
copper drum of 20 m/s, determined using the Rietveld refinement
method
Tabela 1: Kvalitativna sestava faz Nd10–xTbxFe81Zr1B6, (kjer je x = 0
ali 2), strjenih vzorcev z linearno hitrostjo bakrenega valja 20 m/s,
dolo~ena z uporabo Rietveldove metode izpopolnitve
Phase composition
Alloy -Fe (% vol) RE2Fe14B (% vol)
Nd10Fe81Zr1B6 19.15 80.85
Nd8Tb2Fe81Zr1B6 19.40 80.60
Table 2: Matching coefficients of the quantitative phase composition
of Nd10–xTbxFe81Zr1B6 (where x = 0 or 2) samples solidified with the
linear velocity of the copper drum of 20 m/s, determined using the
Rietveld refinement method
Tabela 2: Koeficienti ujemanja kvantitativne sestave faze
Nd10–xTbxFe81Zr1B6 (kjer je x = 0 ali 2) strjenih vzorcev z linearno
hitrostjo bakrenega valja 20 m/s, dolo~eno z uporabo Rietveldove
metode izpopolnitve
R – match parameters
Alloy Rp Rwp Rexp
Nd10Fe81Zr1B6 46.11 197.03 0.18
Nd8Tb2Fe81Zr1B6 44.15 192.79 0.11
The average grain size was determined on the basis
of Bragg’s equation for the 10 and 3 most intense peaks
of the Nd2Fe14B and -Fe, respectively:14
Δ hkl B
hkl
B
hklK
D
d
d
( ) cos( ) sin( )2 2 
 
⋅ =
⋅
+ ⋅ (1)
where D is the average grain size, K is a shape factor of
0.89,  is the X-ray wavelength, and hkl is the half-
width of the peak, d/d is the relative deformation of
the crystalline lattice and  B
hkl is the Bragg angle.
The above-presented equation describes the effect of
the grain size and the relative deformation of the cry-
stalline lattice resulting from diffraction-peak broad-
ening. This relationship is linear and the grain size is
determined from the intersection of the ordinate axis.
The average grain size of the -Fe and Nd2Fe14B pha-
ses was 9 nm and 41 nm, respectively, for Nd10Fe81Zr1B6,
and 24 nm and 70 nm for the alloy Nd8Tb2Fe81Zr1B6 for
samples solidified at a drum linear velocity of 20 m/s.
Due to the low intensity of the peaks originating from
the crystalline phases, for samples solidified at higher
linear velocities of the copper drum, the particle size was
not determined.
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Materiali in tehnologije / Materials and technology 50 (2016) 5, 823–827 825
Figure 3: The magnetic hysteresis loops and the initial magnetization curves for the: a), c) Nd10Fe81Zr1B6 and d), f) Nd8Tb2Fe81Zr1B6 alloys,
solidified at linear speeds of the copper drum equal to: a), d) 20 m/s, b), e) 30 m/s, c), f) 35 m/s
Slika 3: Magnetne histerezne zanke in za~etne krivulje magnetizacije pri zlitinah: a), c) Nd10Fe81Zr1B6 in d), f) Nd8Tb2Fe81Zr1B6 , strjenih pri
linearni hitrosti bakrenega valja: a), d) 20 m/s, b), e) 30 m/s, c), f) 35 m/s
Figure 3 shows the magnetic hysteresis loops and the
initial magnetization curves for all the studied samples.
From the analysis of the initial magnetization curves
and the main hysteresis loops the magnetic parameters,
such as coercivity (JHC), remanence (μ0MR) and satura-
tion of the magnetization (μ0Ms), were determined (Table
3). Additionally, for the materials with the best hard
magnetic properties the ratio MR/MS was also calculated.
It allows estimating the extent to which, in the nano-
composite, the phases with hard and soft magnetic pro-
perties are coupled by means of exchange interac-
tions.12,15,16 An additional indicator proving information
about the existence of such interactions in the nano-
composite comprising the phases of different magnetic
hardness can be the shape of the magnetic hysteresis
loop in combination with the results of the phase com-
position and grain size studies.15–17
Table 3: The basic magnetic parameters of the studied
Nd10–xTbxFe81Zr1B6 (where x =0 or 2) tapes in the as-cast state, deter-
mined from the initial magnetization curves and magnetic hysteresis
loops
Tabela 3: Osnovne magnetne lastnosti prou~evanih
Nd10–xTbxFe81Zr1B6 (kjer je x = 0 ali 2) trakov v litem stanju, dolo-
~enih iz za~etnih magnetizacijskih krivulj in magnetnih histereznih
zank
Linear
velocity v
(m/s)
coercivity
JHC
(T)
remanence
μ0MR
(T)
MR/MS
ratio
(a.u.)
Saturation of the
magnetization
μ0MS (T)
Nd10Fe81Zr1B6
20 0,7 0,82 0,77 1,06
30 9×10–4 – – 1,02
35 4×10–4 – – 1,15
Nd8Tb2Fe81Zr1B6
20 1,03 0,81 0,73 1,10
30 54×10–4 – – 0,81
35 13×10–4 – – 0,93
Based on the data collected in Table 3, it was found
that the material solidified at higher linear speeds of the
rotating copper drum (30 m/s and 35 m/s) was charac-
terized by poor, soft-magnetic properties, which was
related to the polycrystalline structure. Inclusions of
nanocrystalline grains in the amorphous matrix led to the
formation of additional pinning sites, which increases the
coercivity of the sample. Tapes produced with the lowest
cooling rates (20 m/s drum speed) were characterized by
good hard magnetic properties. The substitution of Nd
by Tb in the Nd10–xTbxFe81Zr1B6 alloy, quenched with the
lowest cooling rate resulted in an improvement of the
saturation magnetization and in a significant increase in
the coercivity. The increase in the saturation magneti-
zation was associated with a slightly higher content of
-Fe phase in the alloy composition.
The shape of the magnetic hysteresis loop of a
Nd10Fe81Zr1B6 sample with hard magnetic properties was
smooth, in spite of the presence of the two phases with
different magnetic hardnesses. Additionally, a high
MR/MS ratio (greater than 0.5) and a moderate grain size
of both phases are evidence of a strong coupling between
the grains of different phases, through exchange interac-
tions. The change of 2 % in the fractions of Nd by Tb in
the atomic composition, while maintaining the same
cooling rate, caused the formation of a bimodal shape of
the hysteresis loop. Juxtaposing this with the much larger
average grain size of both phases (almost two times
higher) and a lower MR/MS ratio, it can be concluded that
not all the soft magnetic -Fe grains were covered by ex-
change interactions. In this case, an increase in the
cooling rate by increasing the linear speed of the rotating
copper drum to 22 m/s should result in hardening of the
magnetic properties, because a Nd8Tb2Fe81Zr1B6 alloy
with smaller grain sizes would be obtained and phases
with different magnetic hardnesses would interact with
each other by exchange coupling.
The hysteresis loops for the samples solidified at
higher cooling speeds were of a wasp shape, typically
encountered in materials containing small inclusions of
hard magnetic phase distributed in the amorphous ma-
trix. The exception was the hysteresis loop obtained for
the Nd10Fe81Zr1B6 alloy that was solidified at the fastest
cooling rate. It had a typical shape characteristic for a
magnetically soft material. In this case, the low content
of crystalline phase only slightly affected the shape of
the hysteresis loop, mainly by increasing the coercivity
field (relative to a fully amorphous alloy).
4 CONCLUSIONS
On the basis of X-ray studies, it was found that the
samples made with the 20 m/s linear velocity of the
copper cylinder were crystalline and consisted of two
phases: -Fe and RE2Fe14B. A qualitative analysis of the
phase composition showed that the contents of both
phases were 19.15 and 19.40 of -Fe phase, 80.85 and
80.60 of Nd2Fe14B phase, for Nd10Fe81Zr1B6 and
Nd8Tb2Fe81Zr1B6 alloys, respectively. The grain size
analysis based on Bragg’s equation showed that the
grains in the alloy containing Tb were almost two times
larger. Alloys obtained with higher linear speeds of the
copper drum (30 m/s and 35 m/s) were mainly amor-
phous and contained only small amounts of crystalline
phase precipitates.
From the magnetic studies it was found that the tapes
produced with the lowest cooling rate were characterized
by the best hard magnetic properties. The application of
higher speeds of the copper drum resulted in their dete-
rioration.
The hysteresis loop of Nd10Fe81Zr1B6 (obtained at 20
m/s) was smooth due to presence of strong exchange
coupling between the phases of different magnetic hard-
ness. The substitution of Nd by Tb in these samples
resulted in an improvement of the saturation magne-
tization and in a significant increase in the coercivity, but
also resulted in the formation of a bimodal shape for the
hysteresis loop. The reason for this was decoupling of
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826 Materiali in tehnologije / Materials and technology 50 (2016) 5, 823–827
the exchange interaction due to an increase of the grain
sizes of both phases. Summarizing, for alloy with Tb
addition, an increase in the linear speed of the rotating
copper drum from 20 m/s to 22 m/s should result in a
hardening of the magnetic properties.
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Materiali in tehnologije / Materials and technology 50 (2016) 5, 823–827 827

M. TORKAR et al.: METALOGRAFSKA PREISKAVA IN KOROZIJSKA ODPORNOST ZVAROV ...
829–834
METALOGRAFSKA PREISKAVA IN KOROZIJSKA ODPORNOST
ZVAROV FERITNEGA NERJAVNEGA JEKLA
METALLOGRAPHIC INVESTIGATION AND CORROSION
RESISTANCE OF WELDS OF FERRITIC STAINLESS STEELS
Matja` Torkar, Aleksandra Kocijan, Roman Celin, Jaka Burja, Bojan Podgornik
In{titut za kovinske materiale in tehnologije, Lepi pot 11, 1000 Ljubljana, Slovenija
matjaz.torkar@imt.si
Prejem rokopisa – received: 2016-04-04; sprejem za objavo – accepted for publication: 2016-05-12
doi:10.17222/mit.2016.059
Preiskovani so bili vzorci zvarov feritnega nerjavnega jekla X6Cr17 (W.Nr. 1.4016). Prikazane so metalografske zna~ilnosti
pre~nega in kro`nega zvara. Dolo~ena je bila hitrost korozije v dveh korozijskih medijih in pri dveh temperaturah. Predstavljeni
so rezultati osnovnega materiala, dveh zvarov in dveh primerjalnih materialov, jekla X2CrTi17 in malooglji~nega jekla DC01
EN10130. Rezultati preiskave so pokazali, da se v zvaru pove~a velikost zrn, po mejah zrn pa je izlo~ena martenzitna faza,
oboje pa vpliva na korozijsko obstojnost. Nobeden od preiskovanih vzorcev ni pokazal izrazite pasivacije povr{ine pri
elektrokemijskem korozijskem preizkusu.
Klju~ne besede: feritno nerjavno jeklo, metalografija, zvar, korozijska odpornost
Samples of welds made from the ferritic stainless steel X6Cr17 (W.Nr. 1.4016) were investigated. Here we present the results of
the metallographic charaterisations of these welds. We determined the corrosion rate in two corrosion media at two tempera-
tures. The results from the base material and two welds are presented and compared with the steel X2CrTi17 and the low-carbon
steel DC01 EN10130. The results revealed enlarged grains in the weld and martensite at the grain boundaries, both have an in-
fluence on the corrosion resistance. None of the investigated samples showed distinctive passivation of the surfaces during the
electrochemical corrosion tests.
Keywords: ferritic stainless steel, metallography, weld, corrosion resistance
1 UVOD
Feritna nerjavna jekla so poceni, cenovno stabilna in
korozijsko odporna jekla. Uporaba teh nerjavnih jekel je
pogosta v avtomobilski industriji, pri izdelavi kuhinjskih
pripomo~kov in naprav ter tudi na drugih podro~jih.
Feritna nerjavna jekla se zaradi njihove dobre toplotne
prevodnosti in majhnega toplotnega raztezka uporabljajo
tako za izdelavo dimnikov, glu{nikov, izpu{nih sistemov,
pritrdilnih elementov, kot tudi za grelne elemente, ki se
jih uporablja v kopelih staljene soli za toplotno obdelavo,
v konstrukcijske namene in podobno.1–5
Pri visokih temperaturah iz taline nastaja faza delta
ferita (), ki se potem ko kristalizira, pri ohlajanju ne
spremeni. Te popolnoma feritne mikrostrukture se ne da
toplotno obdelati. Neprisotnost fazne premene razlo`i,
zakaj so pri ogrevanju ta jekla nagnjena k rasti zrn.
Feritna faza v zvaru in toplotno vplivana cona sta
ob~utljivi na intersticijske elemente, kot sta ogljik in
du{ik, ki pri visokih temperaturah z difuzijo spremenita
feritno fazo v avstenit. Ta avstenit pa se pri ohlajanju po
mejah zrn pretvori v martenzit (krhek). Dodaten problem
je tudi z rastjo zrn v toplotno vplivani coni in v zvaru.6,7
Za zadr`anje duktilnosti feritne strukture, ki vsebuje tudi
martenzit, je potrebno izvr{iti toplotno obdelavo ~im prej
po zaklju~nem varjenju. Ta obdelava za odpravo nape-
tosti se izvede pri maksimalni temperaturi 750-800 °C,
kar je malo pod temperaturo nastanka avstenita v pod-
ro~jih bogatih z ogljikom. Namen te obdelave je
popustiti martenzit, zmanj{ati zaostale napetosti v zvaru
ter pove~ati `ilavost zvarnega spoja in toplotno vplivane
cone.
Za prepre~itev rasti zrn in pojava avstenitizacije je
potrebno varjenje izvr{iti z majhnim vnosom toplote
(1 kJ/mm).8
[tudija korozijske odpornosti zvarov je izredno po-
membna za uspe{no uporabo feritnih nerjavnih jekel.
Poleg obna{anja samega materiala v korozijskem mediju,
je pomembna tudi korozijska odpornost zvarnih spojev.
Literatura9, navaja, da je korozijska odpornost zvarov
odvisna tudi od velikosti zrn in prisotnosti razli~nih faz v
mikrostrukturi zvara.
Namen predstavljene preiskave je z metalografsko
preiskavo ugotoviti mikrostrukturne zna~ilnosti dveh vrst
zvarov ter napraviti primerjavo s preizkusom korozijske
odpornosti osnovnega materiala, zvarov in primerjalnih
materialov v dveh korozijskih medijih.
Dobljeni podatki naj bi slu`ili optimiranju tehno-
logije varjenja feritnega nerjavnega jekla in naj bi prispe-
vali k podalj{anju zdr`ljivosti zvarov v danem
korozijskem mediju.
Materiali in tehnologije / Materials and technology 50 (2016) 5, 829–834 829
UDK 620.1:691.714.018.8:621.791.053 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 50(5)829(2016)
2 OPIS VZORCEV
Vzorci, prikazani na Sliki 1, so bili odrezani iz kadi
in ozna~eni kot "pre~ni zvar", izveden pre~no na dno
kadi in "kro`ni zvar", izveden po obodu kadi. Obod in
dno kadi sta izdelana iz enakega feritnega nerjavnega
jekla X6Cr17, pre~ni in kro`ni zvar pa sta izdelana z
razli~nimi varilniki, zaradi ~esar je potrebno napraviti
tudi primerjavo med posameznimi zvari. Dno kadi se na
sliki 1 vidi slabo, ker je pravokotno na obod kadi in je
ozna~en z belo pu{~ico.
Za primerjavo korozijskih hitrosti sta bila dodatno
preizku{ena {e vzorec malo oglji~ne plo~evine kvalitete
DC01 po standardu EN10130 in vzorec plo~evine ferit-
nega nerjavnega jekla X2CrTi17. Nominalna kemijska
sestava za preiskovana jekla je podana v Tabeli 1.
3 EKSPERIMENTALNO DELO
Zvari so bili izdelani po postopku TIG (Tungsten in-
ert gas welding), brez dodajanja materiala.
Odrezani so bili vzorci zvarov (jeklo X6Cr17) in pri-
pravljeni so bili za metalografsko preiskavo po standard-
nem postopku z zalivanjem v maso, bru{enjem in
poliranjem. Za metalografsko preiskavo so bili vzorci
jedkani z Vilella jedkalom (5 cm3 HCl, 2 g pikrinske
kisline in 100 cm3 etil alkohola).
Izvedena je bila metalografska preiskava obeh vrst
zvarov; pre~nega in kro`nega zvara. Mikrostruktura je
bila pregledana s pomo~jo svetlobne mikroskopije, s
svetlobnim mikroskopom Nikon Microphot FXA,
opremljenim z videokamero Olympus DP73 in ra~unal-
ni{kim programom za analizo Stream Motion.
Korozijske meritve so potekale na potenciostatu/
galvanostatu EG&G PAR Model 273 s trielektrodno
korozijsko celico, pri ~emer je bila uporabljena tehnika
potenciodinamske polarizacije. Korozijsko so bile preiz-
ku{ene tri osnovne plo~evine (Tabela 1) in dve vrsti
zvara (pre~ni in kro`ni zvar).
4 REZULTATI
4.1 Metalografski pregled zvarov
Iz vzorcev sta bila pripravljena pre~na prereza obeh
vrst zvarov, pre~nega zvara in kro`nega zvara.
Pregled izgleda povr{ine zvara je pokazal, da se pri
obeh vrstah zvarov na za~etku zvara oziroma na mestu
zdru`evanja zvarov pojavi raz{irjeno podro~je (Slika 2).
V pre~nem preseku pre~nega zvara je opazna ~rta, ki
izgleda kot nezveznost v zvaru (Slika 3). ^rta se pri~ne
na povr{ini, te~e pod kotom 45° do sredine zvara, nato
te~e vzporedno s povr{ino in se nato zopet pod kotom
45° obrne proti nasprotni povr{ini zvara.
Metalografski pregled je pokazal, da gre za nespo-
jeno podro~je med plo~evinama in da se kristalna zrna na
prehodu v zvar pove~ajo. Pri feritnem nerjavnem jeklu se
v zvaru zrna lahko pove~ajo, ker to jeklo nima faznih
premen pri ohlajanju in ogrevanju, zato je potrebno
pazljivo izvajati proces varjenja.
Na Sliki 4 je prikazan prehod iz zvara v toplotno
vplivano cono na preseku pre~nega zvara, na Sliki 5 pa
mikrostruktura v sredini enojnega (~elnega) pre~nega
zvara. Iz Slike 5 je razvidno, da so zrna v zvaru ve~ja kot
M. TORKAR et al.: METALOGRAFSKA PREISKAVA IN KOROZIJSKA ODPORNOST ZVAROV ...
830 Materiali in tehnologije / Materials and technology 50 (2016) 5, 829–834
Tabela 1: Nominalna kemijska sestava za preiskovana jekla
Table 1: Nominal chemical composition for investigated steels
Jeklo C (max.) Si(max.)
Mn
(max.)
P
(max.)
S
(max.) Cr
N
(max.)
Ti
(max.)
Al
(max.)
X6Cr17 0,08 1,0 1,0 0,040 0,015 16 – 18 / / /
X2CrTi17 0,025 0,50 0,5 0,040 0,015 16-18 0,015 0,3-0,6 /
DC01 0,12 0,03 0,6 0,045 0,045 / / / 0,02-0,06
Slika 2: Detajl kro`nega zvara z raz{iritvijo, ki jo ka`e pu{~ica
Figure 2: Detail of circular weld with enlargement shown by arrow
Slika 1: Vzorci s pre~nim in kro`nim zvarom. Jeklo X6Cr17. Svetla
pu{~ica ozna~uje dno kadi, ~rna pu{~ica pa kro`ni zvar.
Figure 1: Samples with cross and circular weld. Steel X6Cr17. White
arrow indicates the bottom of the tub, the black arrow indicates the
circular weld.
v osnovni mikrostrukturi jekla X6Cr17, ki je prikazana
na Sliki 6.
Zna~ilnosti mikrostrukture kro`nega zvara so
prikazane na Slikah od 7 do 9.
V posameznih delih zvara so opazne nezveznosti,
kjer ni pri{lo do popolne pretalitve vsega materiala.
M. TORKAR et al.: METALOGRAFSKA PREISKAVA IN KOROZIJSKA ODPORNOST ZVAROV ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 829–834 831
Slika 3: Detajl preseka skozi pre~ni zvar: a) zgornja povr{ina,
b) sredina in c) spodnja povr{ina. Jeklo X6Cr17, jedkano z Vilella
jedkalom.
Figure 3: Detail of cross-section of cross weld: a) upper surface, b)
middle and c) lower surface. Steel X6Cr17, etched with Vilella’s re-
agent.
Slika 6: Mikrostruktura osnovne plo~evine. Jeklo X6Cr17, jedkano z
Vilella jedkalom.
Figure 6: Microstructure of base sheet material. Steel X6Cr17, etched
with Vilella’s reagent.
Slika 5: Mikrostruktura v sredini pre~nega zvara. Jeklo X6Cr17,
jedkano z Vilella jedkalom.
Figure 5: Microstructure in the middle of the cross weld. Steel
X6Cr17, etched with Vilella’s reagent.
Slika 4: Presek pre~nega zvara. Prehod iz zvara v toplotno vplivano
cono. Jeklo X6Cr17, jedkano z Vilella jedkalom.
Figure 4: Cross-section of the cross weld. Transition from weld to
heat-affected zone. Steel X6Cr17, etched with Vilella’s reagent.
Poleg tega so v podro~ju kro`nega zvara, na posameznih
mestih, opazna tudi pove~ana zrna in izlo~ena marten-
zitna faza po mejah zrn. Spreminjanje mikrostrukture
zvara pomeni, da so se spreminjali parametri med po-
stopkom varjenja.
Zna~ilnosti obeh vrst zvarov ka`eta, da bi bilo
potrebno izbolj{ati, oziroma optimirati izvedbo varjenja.
4.2 Korozijske preiskave
Korozijski preskusi so bili izvedeni v dveh medijih,
ki se jih uporablja pri funkcionalnih meritvah na ferit-
nem nerjavnem jeklu, pri temperaturah 60 °C in 90 °C.
Korozijski medij oz. raztopino smo pripravili iz dveh
pra{kov z oznako IEC-A base in Na-perborate tetra-
hydrate in to na naslednji na~in:
Raztopina 1: 10 g IEC/L H2O s pH 10,5 pri temperaturi
60 °C
Raztopina 2: 8 g IEC + 2 g Na-perborata/L H2O pH 11,
pri temperaturi 90 °C
Vse korozijske meritve so potekale na potenciostatu/
galvanostatu z uporabo tehnike potenciodinamske
polarizacije. Ta tehnika postopnega dvigovanja poten-
ciala na enoto ~asa, k bolj pozitivnim vrednostim,
spreminja korozijsko stabilnost kovine. Z nara{~anjem
potenciala postaja material termodinamsko vedno manj
stabilen, kar pomeni, da nara{~a tudi korozijski tok, ki je
merilo za stopnjo korozije. Tako spoznamo splo{no
kvalitativno sliko o obna{anju materiala v nekem mediju.
Iz oblike anodnega dela krivulje lahko sklepamo o koro-
zijski odpornosti, morebitni pasivaciji, potencialu pre-
boja in podobno.10 Iz Taflovega dela krivulje (±250 mV
glede na Ekor) lahko izra~unamo tudi korozijsko hitrost in
polarizacijsko upornost, s tem pa stopnjo korozije.
Za delovno elektrodo so bili uporabljeni naslednji
vzorci:
Vzorec 1 (dno kadi – X6Cr17)
Vzorec 2 (kro`ni zvar – X6Cr17)
Vzorec 3 (pre~ni zvar – X6Cr17)
Vzorec 4 (malo oglji~no jeklo kvalitete DC01)
M. TORKAR et al.: METALOGRAFSKA PREISKAVA IN KOROZIJSKA ODPORNOST ZVAROV ...
832 Materiali in tehnologije / Materials and technology 50 (2016) 5, 829–834
Slika 9: Mikrostruktura kro`nega zvara z izlo~enim martenzitom po
mejah zrn. Jeklo X6Cr17, jedkano z Vilella jedkalom
Figure 9: Microstructure of circular weld with a martensite phase on
the grain boundaries. Steel X6Cr17, etched with Vilella’s reagent
Slika 7: Prerez prehoda v kro`ni zvar. Jeklo X6Cr17, jedkano z Vilella
jedkalom
Figure 7: Cross-section of transition to a circular weld. Steel X6Cr17,
etched with Vilella’s reagent
Slika 8: a) Napaka v sredini kro`nega zvara in b) rast zrn. Jeklo
X6Cr17, jedkano z Vilella jedkalom
Figure 8: a) Failure in the central part of a circular weld and c) grain
growth. Steel X6Cr17, etched with Vilella’s reagent
Vzorec 5 (jeklo X2CrTi17).
Izpostavljena povr{ina elektrodne reakcije na delovni
elektrodi je bila 1 cm2, referen~na elektroda je bila nasi-
~ena kalomelova elektroda in pomo`na elektroda je bila
grafitna palica.
V Tabeli 1 so podani rezultati potenciodinamskih
meritev za korozijsko hitrost, korozijski potencial, koro-
zijski tok in polarizacijska upornost petih preiskovanih
vzorcev v raztopini 1, v Tabeli 2 pa rezultati za razto-
pino 2.
Na Slikah 10 in 11 so prikazane potenciodinamske
polarizacijske krivulje preiskovanih vzorcev v raztopini 1
pri 60 °C in raztopini 2 pri 90 °C.
Tabela 1: Rezultati potenciodinamskih meritev petih vzorcev v raz-
topini 1, pri 60 °C
Table 1: Results of potentiodynamic measurements of five samples in
solution 1, at 60 °C
Vzorec Korozijskahitrost E (I = 0) Icorr Rp
1 (dno kadi)
X6Cr17
0,0013
mm/leto –151,3 mV 120,0 nA 277,0 k

2 (kro`ni zvar)
X6Cr17
0,036
mm/leto –242,1 mV 3282 μA 13,16 k

3 (pre~ni zvar)
X6Cr17
0,015
mm/leto –281,5 mV 1399 μA 27,26 k

4 (DC01) 0,059mm/leto –340,5 mV 6132 μA 5382 k

5 (X2CrTi17) 0,0013mm/leto –240,8 mV 116,1 nA 310,5 k

Preiskave potenciodinamske polarizacije petih preis-
kovanih vzorcev so pokazale, da sta vzorec 1 (dno kadi –
X6Cr17) in vzorec 5 (X2CrTi17) v raztopini 1 pri 60 °C
korozijsko najbolj odporna, kar je razvidno iz najni`jih
vrednosti korozijskih hitrosti (1,3 μm/leto) in najvi{jih
vrednosti polarizacijske upornosti Rp (277–310 k
).
Pri~akovano je korozijsko najmanj stabilen vzorec 4
(malo oglji~no jeklo), s 45- do 60-krat slab{o korozijsko
odpornostjo, medtem ko sta vzorca 2 in 3 (kro`ni zvar –
X6Cr17 in pre~ni zvar –X6Cr17) dokaj primerljiva, a {e
vedno za en velikostni razred manj korozijsko odporna
kot osnovni material. Domnevamo, da je to posledica
razlike v velikosti zrn in zna~ilnosti strjevalne strukture
zvarov ter pojava martenzitne faze po mejah zrn. Vidimo
lahko tudi, da je pre~ni zvar 2-krat korozijsko bolj
odporen kot kro`ni zvar, kjer se opazi v zvaru marten-
zitna faza po mejah zrn. V raztopini 2 pri 90 °C so
vrednosti korozijske hitrosti 10-krat vi{je, kar je posle-
dica bolj agresivnega medija. V primeru vzorca 4 (malo
oglji~no jeklo) je ta pojav {e posebej izrazit, kjer se
ocenjena korozijska hitrost pove~a kar za dva velikostna
razreda na pribli`no 8 mm/leto. Tudi v raztopini 2 sta
korozijsko najbolj stabilna vzorca 1 (X6Cr17) in 5
(X2CrTi17) z ocenjeno korozijsko hitrostjo 50–60
M. TORKAR et al.: METALOGRAFSKA PREISKAVA IN KOROZIJSKA ODPORNOST ZVAROV ...
Materiali in tehnologije / Materials and technology 50 (2016) 5, 829–834 833
Slika 10: Potenciodinamska krivulja v raztopini 1 pri 60 °C: a) X6Cr17 – dno kadi, b) X6Cr17 – kro`ni zvar, c) X6Cr17 – pre~ni zvar, d) malo-
oglji~no jeklo CD01, EN10130 in e) primerjalno jeklo X2CrTi17
Figure 10: Potentiodynamic curve in solution 1 at 60 °C: a) X6Cr17 – bottom of the tub, b) X6Cr17 – circular weld, c) X6Cr17 – cross weld, d)
low-carbon steel CD01, EN10130 and e) X2CrTi17 comparative steel
Tabela 2: Rezultati potenciodinamskih meritev petih vzorcev v raz-
topini 2 pri 90 °C
Table 2: Results of potentiodynamic measurements of five samples in
solution 2 at 90 °C
Vzorec Korozijskahitrost E (I = 0) Icorr Rp
1 (dno kadi)
X6Cr17
0,059
mm/leto –21,62 mV 5442 μA 6063 k

2 (kro`ni zvar)
X6Cr17
0,411
mm/leto –40,08 mV 37,77 μA 1137 k

3 (pre~ni zvar)
X6Cr17
0,351
mm/leto –78,87 mV 32,22 μA 781,1 

4 (DC01) 8,167mm/leto –434,4 mV 847,4 μA 56,51 

5 (X2CrTi17) 0,049mm/leto –4851 mV 4568 μA 7700 k

μm/leto in polarizacijsko upornostjo 6–7 M
, najmanj
pa vzorec 4. Med vzorcema 2 in 3 pa v raztopini 2 prak-
ti~no ni razlik, {e vedno pa ne dosegata korozijske
odpornosti osnovnega materiala (vzorec 1).
Med korozijskim preskusom nobeden od preisko-
vanih vzorcev ni pokazal izrazite pasivacije povr{ine.
5 ZAKLJU^KI
Na podlagi primerjave narejenih raziskav zvarov in
plo~evin lahko postavimo naslednje zaklju~ke:
Zvari, tako pre~ni kot kro`ni, so izvedeni slabo, saj
material ni pretaljen po vsej dol`ini zvara. Potrebno bo
tehnolo{ko optimiranje tehnologije izvedbe varjenja.
Razlike v mikrostrukturi in velikosti zrn ter nespo-
jena mesta v zvaru ka`ejo, da so se procesni parametri
med varjenjem spreminjali.
Izvr{ene korozijske preiskave v dveh medijih in pri
dveh temperaturah so pokazale, da sta na korozijo v obeh
primerih najbolj odporna dno kadi iz feritnega nerjav-
nega jekla X6Cr17 in primerjalno jeklo X2CrTi17. Na
mestu zvara se korozijska odpornost materiala zmanj{a
za cel velikostni razred.
Slab{a korozijska odpornost kro`nega zvara je
posledica nastanka martenzitne faze po mejah zrn v
zvaru.
Od preiskovanih vzorcev je korozijsko najmanj stabi-
len vzorec malooglji~nega jekla.
Pri nobenem od preiskovanih vzorcev se med koro-
zijskim preizkusom v dveh korozijskih medijih ni poka-
zala izrazita pasivacija povr{ine.
6 LITERATURA
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Slika 11: Potenciodinamska krivulja v razopini 2 pri 90 °C: a) X6Cr17 – dno kadi, b) X6Cr17 – kro`ni zvar, c) X6Cr17 – pre~ni zvar, d) malo-
oglji~no jeklo CD01 po standardu EN10130 in e) primerjalno jeklo X2CrTi17
Figure 11: Potentiodynamic curve in solution 2 at 90 °C: a) bottom of the tub, b) circular weld, c) cross weld, d) low-carbon steel CD01,
EN10130 and e) X2CrTi17 comparative steel