VSEBINA – CONTENTS
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Kinetics of weathered-crust elution-deposited rare-earth ore in a leaching process
Kinetika postopka lu`enja elucijsko nanesene preperele skorje rude redkih zemelj
L. Zhang, X. Deng, W. Li, Y. Ding, R. Chi, X. Zuo . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 145
Microstructural and phase analysis of CuAlNi shape-memory alloy after continuous casting
Mikrostrukturna in fazna analiza spominske zlitine CuAlNi po kontinuirnem litju
M. Goji}, S. Ko`uh, I. An`el, G. Lojen, I. Ivani}, B. Kosec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 149
Characterization and determination of mechanical properties of YBCO superconducting thin films with manganese using the
TFA-MOD method
Karakterizacija in dolo~itev mehanskih lastnosti superprevodne tanke plasti YBCO z manganom po metodi TFA-MOD
O. Culha, I. Birlik, M. Toparli, E. Celik, S. Engel, B. Holzapfel. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153
Distortion of the substructure of a 20-ft shipping container exposed to zinc hot-dip galvanizing
Spreminjanje mer podstrukture pri 20 ft transportnem kontejnerju pri vro~em potopnem cinkanju
I. Ivanovi}, A. Sedmak, R. Rudolf, L. Gusel, B. Gruji} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 161
Optimization of the drilling parameters for the cutting forces in B4C-reinforced Al-7XXX-series alloys based on the Taguchi
method
Optimiranje parametrov vrtanja za sile vrtanja pri zlitinah Al-7XXX, oja~anih z B4C s Taguchijevo metodo
A. Taþkesen, K. Kütükde . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 169
Wear properties of AISI 4140 steel modified with electrolytic-plasma technology
Obrabne lastnosti jekla AISI 4140, modificiranega s tehnologijo elektrolitske plazme
A. Ayday, M. Durman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 177
Tribological behavior of a plasma-sprayed Al2O3-TiO2-Cr2O3 coating
Tribolo{ko pona{anje s plazmo napr{enega Al2O3-TiO2-Cr2O3 nanosa
Y. Sert, N. Toplan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 181
Characterization of selected phase-change materials for a proposed use in building applications
Karakterizacija izbranih materialov s fazno premeno za predlagano uporabo v gradbeni{tvu
M. Ostrý, R. Pøikryl, P. Charvát. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 185
Improvement of the damping properties of carbon-fibre-reinforced laminated plastics using damping layers
Izbolj{anje du{enja z ogljikovimi vlakni oja~ane laminirane plastike z uporabo plasti za du{enje
R. Kottner, J. Vacík, R. Zem~ík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 189
AA413.0 and AA1050 joined with friction-stir welding
Spajanje zlitine AA413.0 in AA1050 z gnetenjem
S. Kastelic, J. Tu{ek, D. Klob~ar, J. Medved, P. Mrvar. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 195
Increasing tool life during turning with a variable depth of cut
Pove~anje zdr`ljivosti orodja pri stru`enju z variabilno globino reza
M. Sadílek, R. ^ep, Z. Sadílková, J. Valí~ek, L. Petøkovská . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 199
Raman investigation of sol-gel anticorrosion coatings on electronic boards
Ramanske raziskave sol-gel protikorozijskih prevlek na elektronskih vezjih
A. Rauter, M. Ko`elj, L. Slemenik Per{e, A. [urca Vuk, B. Orel, B. Bengû, O. Sunetci . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 205
Numerical study of Rayleigh-Bénard natural-convection heat-transfer characteristics of water-based Au nanofluids
Numeri~na analiza prenosa toplote nanoteko~in voda-Au v razmerah Reyleigh-Bénardove naravne konvekcije
P. Ternik, R. Rudolf, Z. @uni~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 211
Alumothermic reduction of ilmenite in a steel melt
Alumotermi~na redukcija ilmenita v jekleni talini
J. Burja, F. Tehovnik, J. Lamut, M. Knap . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 217
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 47(2)143–258(2013)
MATER.
TEHNOL.
LETNIK
VOLUME 47
[TEV.
NO. 2
STR.
P. 143–258
LJUBLJANA
SLOVENIJA
MAR.–APR.
2013
Analysis of corrosion properties of melt spun Nd-Fe-B ribbons coated by alumina coatings
Analiza korozijskih lastnosti hitro strjenih Nd-Fe-B-trakov, opla{~enih z aluminijevim oksidom
D. Sojer, I. [kulj, S. Kobe, J. Kova~, P. J. McGuiness . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 223
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Influence of the microstructure on machining a central housing made of pearlite grey cast iron
Vpliv mikrostrukture na obdelovalnost centralnega ohi{ja iz perlitne sive litine
N. [trekelj, M. Nuni}, I. Nagli~, B. Markoli . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 229
The wet-chemical synthesis of functionalized Zn1–xOMnx quantum dots utilizable in optical biosensors
Mokra kemijska sinteza funkcionaliziranih kvantnih delcev Zn1–xOMnx, uporabnih v opti~nih biosenzorjih
M. Alizadeh, R. Salimi, H. Sameie, A. A. Sarabi, A. A. Sabbagh Alvani, M. R. Tahriri . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 235
Influence of aluminium-alloy remelting on the structure and mechanical properties
Vpliv ve~kratnega pretaljevanja aluminijevih zlitin na strukturo in mehanske lastnosti
M. Cagala, M. Bøuska, P. Lichý, J. Beòo, N. [pirutová . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 239
Tensile properties of cold-drawn low-carbon steel wires under different process parameters
Natezne lastnosti hladno vle~ene malooglji~ne jeklene `ice pri razli~nih parametrih procesa
C. S. Çetinarslan, A. Güzey. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 245
Impeller-blade failure analysis
Preiskava po{kodbe lopatice rotorja
R. Celin, F. Tehovnik, F. Vodopivec, B. @u`ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 253
L. ZHANG et al.: KINETICS OF WEATHERED-CRUST ELUTION-DEPOSITED RARE-EARTH ORE ...
KINETICS OF WEATHERED-CRUST
ELUTION-DEPOSITED RARE-EARTH ORE IN
A LEACHING PROCESS
KINETIKA POSTOPKA LU@ENJA ELUCIJSKO NANESENE
PREPERELE SKORJE RUDE REDKIH ZEMELJ
Lili Zhang1,2, Xiangyi Deng1,2, Wen Li2, Yigang Ding1, Ru’an Chi1, Xiaohua Zuo2
1Key Laboratory for Green Chemical Process of Hubei Province and the Ministry of Education, Wuhan Institute of Technology,
430073 Wuhan, China
2School of Chemical and Materials Engineering, Hubei Polytechnic University, 435003 Huangshi, China
deng29606@sina.com, wenl@ualberta.ca
Prejem rokopisa – received: 2012-07-23; sprejem za objavo – accepted for publication: 2012-09-13
The leaching reaction kinetics of weathered-crust, elution-deposited rare earth with mixed ammonium salts was studied. The
influence of the concentration of the reagents and the particle size of the ore on the leaching rate was investigated. The results
showed that the diffusion process and the leaching rate could be improved by increasing the reagent concentration and
decreasing the leaching flowing rate and particle size. The leaching process could be explained using the shrinking-core model,
which could be controlled with the diffusion rate of the reacting reagents in a porous solid layer. The leaching rate followed the
equation 1–2/3 – (1 – )2/3 = 7.126 × 10–4C0.3038R0.1942t.
Keywords: leaching reaction, leaching rate, rare-earth ore, leaching process
Prou~evana je kinetika reakcije lu`enja preperele skorje elucijsko nanesene rude redkih zemelj. Preiskovan je bil u~inek
koncentracije reagentov in velikosti zrn rude na hitrost lu`enja. Rezultati so pokazali, da je mogo~e proces difuzije in hitrost
lu`enja pove~ati s pove~anjem koncentracije reagentov ter z zmanj{anjem hitrosti lu`enja in zmanj{anjem velikosti delcev.
Postopek lu`enja se lahko razlo`i z modelom kr~enja jedra, ki ga je mogo~e kontrolirati s kontroliranjem hitrosti difuzije
reagentov, ki reagirajo v poroznem trdnem sloju. Hitrost lu`enja je skladna z ena~bo
1 – 2/3 – (1 – )2/3 = 7,126 × 10–4 C0,3038R0,1942t.
Klju~ne besede: reakcija lu`enja, hitrost lu`enja, ruda redke zemlje, postopek lu`enja
1 INTRODUCTION
Weathered-crust, elution-deposited, rare-earth ore is
China’s unique rare-earth mineral resource.1–3 There are
many advantages of the ore: a widespread distribution of
rich reserves, a low radioactivity, it is rich in the middle
and heavy rare earths, it is an easily extracted rare earth,
it is simple to process by leaching giving high-quality
products, etc. The development and utilization of the ore
in the world have a significant influence. In China,
special attention has been paid to rare-earth mineral
resources in recent years. Therefore, high efficiency and
the comprehensive exploitation of the weathered-crust,
elution-deposited, rare-earth ore was intensively investi-
gated over the past decade.
The leaching of the rare-earth ore is a liquid-solid
multiphase reaction process, where the most common
reaction is carried out in two ways: by using the
integral-scaled model and the shrinking-core model. In
the latter, a mixed ammonium salt solution is used as a
leaching agent, causing the particle size of the ore to
change only a little during the leaching reaction. For this
reason, the leaching process of the rare-earth ore is
usually based on the shrinking-core model. The leaching
rate is connected with the concentration, the temperature,
the surface area of the solid phase, etc. The leaching
process can be controlled by the outer diffusion, the
inner diffusion, or the chemical reaction.4,5 The
influences of the concentration of the reagents and the
particle size of the ore on the leaching rate were
investigated to achieve a high rare-earth concentration, a
low consumption of the leaching reagent and a high
leaching rate.
2 EXPERIMENTAL WORK
2.1 Analysis of the ore samples
Ore samples are mostly random and taken from
non-cemented sands; they are of a pale flesh-red color,
containing clay minerals, quartz sand, rock-forming
feldspar etc., obtained from Dingnan County of Jiangxi
Province. They contain about 40–70 % of clay minerals,
such as halloysite, illite, kaolinite and a small quantity of
montmorillonite. Some 90 % of rare-earth ions adsorp-
Materiali in tehnologije / Materials and technology 47 (2013) 2, 145–148 145
UDK 621.794.4:546.65 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)145(2013)
Table1: The main components of the ore
Tabela 1: Glavne sestavine rude
main
components REO SiO2 Al2O3 Fe2O3 CaO others
content (%) 0.1146 61.80 14.28 3.190 0.470 20.15
tion is in a state of kaolinite and mica. The chemical
composition of the ore is presented in Table 1.
2.2 Process and Method
Prior to the leaching experiments, the ore sample was
sieved through a +20 mesh, 20–60 mesh, 60–100 mesh,
100–140 mesh, +140 mesh to obtain five natural grain
grades. Then the synthetic ore samples, according to the
proportion of the natural grain grades, were weighed into
a glass flask. The leaching agent with a certain liquid-
solid ratio and a flow velocity puts the rare-earth ore into
consideration. Then there is continuous leaching and the
collected leaching solution. The RE3+ concentration was
determined using the EDTA titration method,6 measuring
the volume of the leaching reagent. The leaching process
was evaluated by the RE leaching rate, which was
calculated according to the following formula:
 = /0 (1)
where  and 0 are the amount of leaching-out RE3+ and
the total RE3+ content of the sample ore.
2.3 Leaching mechanism for rare earth
The leaching process of the weathered-crust, elu-
tion-deposited, rare-earth ore is a kind of ion-exchange-
able process between the positive ions in the solution and
the clay minerals.7,8 The chemical reaction equation is as
follows:
[Al4(Si4O10)(OH)8]m·RE S( )
3 + + 3nNH 4
+ (aq) 
 [Al4(Si4O10)(OH)8]m·(NH 4
+ )3n(S) + nRE aq( )
3 + (2)
The weathered-crust, elution-deposited, rare-earth
ore is composed of ore particles, while the leaching
process of the ore is a typical liquid-solid heterogeneous
reaction. The leaching process can be described by the
shrinking-core model and subdivided into five steps, as
follows:9
• Diffusion of the leaching reagent (NH4+) through the
film surrounding the particle to the surface of the clay
minerals (Outer diffusion);
• Penetration and diffusion of NH4+ to the surface of
the un-reacted core (Inner diffusion);
• Reaction of RE3+ with NH4+ (Chemical reaction);
• Diffusion of the RE3+ exchanged though the remain-
der back to the exterior surface of the clay minerals
(Inner diffusion);
• Diffusion of the RE3+ exchanged though the exterior
surface back into the solution of the fluid (Outer
diffusion).
The kinetic control model of the RE leaching process
can be divided into four models.10,11
1. Chemical reaction control:
1 – (1 – )1/3 = k1t
2. Diffusion through the liquid film control:
1 – (1 – )1/3 = k1t
3. Diffusion through the porous ore matrix control:
1 –
2
3
a – (1 – a)2/3 = k3t
4. Mixed control:
1 – (1 – a)1/3 =
k k
k k
C M
r p
1 2
1 2
0
0+
where, k1, k2, k3 are the constants for the different con-
trol steps, respectively. a, t, C0, r0, p and M represent the
rare-earth leaching rate, the leaching time, the initial
concentration of the leaching reagent, the initial radius
of the ore particle, the mole density of the ore particle
and the mass of the ore particle, respectively.
3 RESULTS AND DISCUSSION
According to the optimum process of mixed ammo-
nium salts, the leaching rare-earth ore experiment showed
that the leaching rate of the rare-earth ore can reach up to
94.05 % when 2.0 % NH4NO3 and (NH4)2SO4 with a
quality ratio of 7 : 3, solid-liquid ratio of 0.5 : 1, and
flow rate of 0.5 mL/min. The rare-earth leaching dyna-
mics was researched on the basis of this process to
determine the influence of the factors of the leaching rate
and the control steps.
3.1 Effect of the leaching-reagent concentration
The initial average particle size of 0.2414 mm for the
rare-earth ore was leached by different concentrations of
mixed ammonium salts under the conditions of NH4NO3
and (NH4)2SO4 with a quality ratio of 7 : 3 , a solid-
liquid ratio of 0.5 : 1, and a flow rate of 0.5 mL/min. The
influence of the different concentrations of reagents on
the leaching rate was investigated. According to the
reaction system, the reaction rate is proportional to the
concentration of the product, increasing the leaching-
L. ZHANG et al.: KINETICS OF WEATHERED-CRUST ELUTION-DEPOSITED RARE-EARTH ORE ...
146 Materiali in tehnologije / Materials and technology 47 (2013) 2, 145–148
Figure 1: Effect of the concentration of reagents on rare-earth
leaching rate
Slika 1: Vpliv koncentracije reagentov na hitrost lu`enja redke zemlje
agent concentration is conducive to an improvement of
the leaching reaction rate, thereby increasing the rare-
earth leaching rate. Figure 1 shows how the leaching
rate of the rare earth increases as the initial concentration
of ammonium salt increases, when the leaching agent
concentration of 2.0 % in 275 min for the rare-earth
leaching rate is 92.79 %, with the mass fraction of
leaching agent increases, the unit volume of ammonium
nitrate leaching agent rate increases, leading to the main
leaching agent in a large number of Al3+ impurities, some
of which will cover the surface of the rare-earth ore and
hinder RE3+ leaching.
The date of the rare-earth leaching rate was substitu-
ted into the shrinking-core model in Figure 2. When the
mass fraction of the leaching agent is greater than or equal
to 2 %, it satisfies the equation 1 – 2/3 – (1 – )2/3 = kt.
Indicating that the leaching process step of the rare-earth
ore is inner-diffusion controlled, and it can be used in the
model equation:
1 – 2/3 – (1 – )2/3 = k'CaRbt (3)
where  is the rare-earth leaching rate (%); C is the con-
centration of the leaching reagent (g/L); R is the initial
radius of the ore particle (mm) and t is the leaching time
(min). This equation reflects the influence of leaching
concentration and the particle size of the rare earth on
the leaching rate.
Figure 2 can be drawn for different concentrations of
the NH4NO3 and (NH4) 2SO4 compounds of the rare-
earth leaching agent and the apparent rate constant k
value, assuming that the apparent rate constant is propor-
tional to the power function of the concentration of the
leaching agent, for which ln k = B + a ln C, and a least
squares linear fit with the slope requirements.
Figure 3 shows a linear relationship between ln k and
ln C, and the apparent reaction order is 0.3038, so that a
= 0.3038.
3.2 Effect of the particle size of the ore
A rare-earth ore with a different particle size was
leached by mixed ammonium salts under the condition
of 2 % NH4NO3 and (NH4)2SO4 with a quality ratio of 7 :
3, a solid-liquid ratio of 0.5 : 1, and a flow rate of 0.5
mL/min. The influence of the different particle size on
the leaching rate was investigated.
Figure 4 presents the effect of particle size on the
rare-earth leaching. It shows that the rare-earth leaching
rate of ore particles with a smaller size is greater. It is
known that the smaller the particle size, the more pores,
and at the same time this shortens the length of the pore,
thereby reducing the resistance of the inner diffusion and
shortens the spread time with the effect of speeding up
the rare-earth leaching. The leaching process could be
L. ZHANG et al.: KINETICS OF WEATHERED-CRUST ELUTION-DEPOSITED RARE-EARTH ORE ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 145–148 147
Figure 3: Relation between ln C and ln k
Slika 3: Odvisnost med ln C in ln k
Figure 4: Effect of particle size on rare-earth leaching rate
Slika 4: U~inek velikosti zrn na hitrost lu`enja redke zemlje
Figure 2: Leaching kinetic data for different concentrations of rare
earth
Slika 2: Podatki za kinetiko lu`enja pri razli~nih koncentracijah redke
zemlje
explained with the shrinking-core model according to the
ore particle-size experiment in Figure 5.
As shown in Figure 6, a good linearity exists bet-
ween ln kd and ln R, and the leaching process could be
controlled by the diffusion rate of the reacting reagents
in a porous solid layer. Furthermore, ln kd = 6.4483 +
0.1942 ln R, so that b = 0.1942.
a and b were presented in the equation 1 – 2/3 – (1 –
)2/3 = k'CaRbt, changing one factor, fixing others at the
same time, so that k' is 7.126 × 10–4. As a result, the
equation of the mixed ammonium salts leaching the
rare-earth ore is:
1 – 2/3 – (1 – )2/3 = 7.126 × 10–4 C0.3038R0.1942t (4)
4 CONCLUSIONS
The leaching kinetics of weathered-crust, elution-
deposited, rare earth with mixed ammonium salts was
investigated. The shrinking-core model with inner
diffusion control was used to describe the leaching pro-
cess of the rare earth. It was summarized as follows: the
leaching rate increases with the increasing leaching
reagent concentration and a decreasing of the particle
size. The kinetics of the weathered-crust, elution-depo-
sited, rare-earth leaching equation can be expressed as:
1 – 2/3 – (1 – )2/3 = 7.126 × 10–4 C0.3038R0.1942t.
Acknowledgement
This work was financially supported by National
Natural Science Foundation of China (50974098) and
Innovation team Fund of the Ministry of Education (IRT
0974).
5 REFERENCES
1 R. Chi, J. Tian, Weathered crust elution-deposited rare earth ores
[M], Nova Science Publishers, Inc., New York, 2008
2 R. Chi, Z. Guocai, Rare earth partitioning of granitoid weathering
crust in southern China [J], Transactions of Nonferrous Metal
Society of China, 8 (1998) 4, 693–704
3 R. Chi, J. Tian, Review of weathered crust rare earth ore [J], Journal
of the Chinese rare earth society, 25 (2007), 641–650
4 T. Xunzhong, L. Maonan, In-situ Leach Mining Of Ion-Absorbed
Rare Earth Mineral [J], Mining research and development, 17 (1997)
2, 1–4
5 M. Dingcheng, Metallurgical kinetic study [M], Central South Uni-
versity of Technology Publisher, 1987
6 F. Zhaoheng, Leaching [M], Metallurgical Industry Publishers, 2007
7 C. Ruan, D. Zuxu, O. Zhigao, W. Yuanxin, W. Cunwen, Correlation
analysis on partition of rare earth in ion-exchangeable phase from
weathered crust ores [J], Transactions of Nonferrous Metal Society
of China, 16 (2006), 1421
8 Y. Huiqin, O. Y. K. Xian, R. Guohua, A Study on Leaching Rare
Earth from the Weathered Elution deposited Rare Earth Ore with
Compound Leaching Reagent [J], Jiangxi Science, 23 (2005) 6,
721–726
9 J. Tian, Y. Jingqun, Kinetic and mass transfer study on leaching a
south china rare earth ore [J], Rare Metal, 22 (1996) 5, 330–342
10 J. Tian, L. Shengliang, Y. Jingqun, Kinetic study on leaching a south
china rare earth ore [J], Engineering Chemistry and Metallurgy, 16
(1995) 3, 354–357
11 H. Y. Sohn, M. E. Wadsworth, Extraction metallurgy rate process
[M], Metallurgical Industry Publishers, 1983
L. ZHANG et al.: KINETICS OF WEATHERED-CRUST ELUTION-DEPOSITED RARE-EARTH ORE ...
148 Materiali in tehnologije / Materials and technology 47 (2013) 2, 145–148
Figure 6: Relation between ln kd and ln R
Slika 6: Odvisnost med ln kd in ln R
Figure 5: Leaching kinetic data for rare earths of different particle
size
Slika 5: Podatki o kinetiki lu`enja redke zemlje pri razli~nih veliko-
stih zrn
M. GOJI] et al.: MICROSTRUCTURAL AND PHASE ANALYSIS OF CuAlNi SHAPE-MEMORY ALLOY ...
MICROSTRUCTURAL AND PHASE ANALYSIS OF
CuAlNi SHAPE-MEMORY ALLOY AFTER CONTINUOUS
CASTING
MIKROSTRUKTURNA IN FAZNA ANALIZA SPOMINSKE ZLITINE
CuAlNi PO KONTINUIRNEM LITJU
Mirko Goji}1, Stjepan Ko`uh1, Ivan An`el2, Gorazd Lojen2, Ivana Ivani}1, Borut Kosec3
1University of Zagreb, Faculty of Metallurgy, Aleja narodnih heroja 3, 44103 Sisak, Croatia
2University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia
3University of Ljubljana, Faculty of Natural Sciences and Engineering, A{ker~eva cesta 12, 1000 Ljubljana, Slovenia
gojic@simet.hr
Prejem rokopisa – received: 2012-07-24; sprejem za objavo – accepted for publication: 2012-09-14
The results of the characterization of a CuAlNi shape-memory alloy after continuous casting technology are shown. Using this
procedure a bar with a diameter of 8 mm was manufactured. After solidification of the alloy the microstructure characterization
was carried out using optic microscopy (OM), scanning electron microscopy (SEM), differential scanning calorimetry (DSC)
and X-ray diffraction (XRD) methods. Our results showed that the as-cast alloy consisted of the parent 1 and 1’ martensite
phases. The martensite phase primary as the needle-like inside grains was observed. Martensite laths have different orientations
inside particular grains. It was found that the average grains size is 98.78 μm. The grain diameter near to the external surface is
higher than in the center. The average hardness of the alloy was 275 HV1.
Keywords: shape memory alloys, martensite, continuous casting, grain size
Prikazani so rezultati karakterizacije spominske zlitine CuAlNi po postopku kontinuirnega litja. Po navedeni proceduri so bile
izdelane palice premera 8 mm. Po strjevanju zlitine je bila izvedena karakterizacija mikrostrukture z uporabo metod opti~ne
mikroskopije (OM), vrsti~ne elektronske mikroskopije (SEM), diferencialne vrsti~ne kalorimetrije (DSC) in rentgenske
strukturne analize (XRD). Rezultati so pokazali, da lita zlitina sestoji iz izhodne faze 1 in martenzitne faze 1’. Opa`ena je bila
martenzitna faza v obliki iglic v primarnih zrnih. Letve martenzita imajo razli~ne orientacije v posameznih zrnih. Ugotovili
smo, da je srednja velikost zrn 98,78 μm. Premer zrn v v bli`ini zunanje povr{ine je ve~ji kot v centru. Srednja trdota zlitine je
bila 275 HV1.
Klju~ne besede: spominske zlitine, martenzit, kontinuirno litje, velikost zrna
1 INTRODUCTION
Shape-memory alloys (SMAs) demonstrate the
ability to return to some previously defined shape or size
when they are exposed to the appropriate thermal treat-
ment. The condition necessary to enable the memory
effect is the presence of a reversible phase transfor-
mation of austenite to martensite. Such phase trans-
formations can be obtained by mechanical (loading) or
thermal methods (cooling and heating). The main types
of SMAs are Ni-Ti (nitinol), Cu-based and Fe-based
alloys1–8. The main advantage of Cu-based SMAs is their
low price compared to other SMAs. The properties of
Cu-Al-Ni alloys are superior to those of Cu-Zn-Al alloys
due to their wide range of useful transformation
temperatures and small hysteresis. Cu-Al-Ni alloys can
be applied at higher temperatures (close to 200 °C).
Generally, ternary Cu-based shape-memory alloys
show a very large grain size. This problem can be solved
by the addition of appropriate refining elements (Zr, Ti,
B etc.) due to the formation of precipitates that limit the
grain size and grain growth9–11 and/or by applying the
technology of rapid solidification. Generally, one of
reasons for using the technique of rapid solidification is
to obtain a small grain size for the SMAs12,13. The grain
sizes obtained in the Cu-base are of the order of 10 μm in
alloys produced by powder metallurgy and by rapid
solidification14. Melt-spinning is the most commonly
used technique for the production of ribbons15,16. In
recent years the continuous casting technique is one of
the technologies for production of SMAs due to the
special competitive growth mechanism of the crystals
and formation of a cast product with a favorable
texture17,18.
In the present paper the microstructure of Cu-Al-Ni
SMAs obtained directly from the melt by continuous
casting techniques are shown. The main aim of this
paper was to obtain a homogenous martensite micro-
structure by solidification without any heat-treatment
procedure.
2 EXPERIMENTAL
The Cu-Al-Ni shape-memory alloy (a bar of 8 mm
diameter) was fabricated by means of the continuous
casting technique using a device for vertical continuous
casting that is connected with a vacuum induction
furnace. The heating temperature was 1230 °C. The
Materiali in tehnologije / Materials and technology 47 (2013) 2, 149–152 149
UDK 621.74.047:537.533.35 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)149(2013)
process of casting including the initial melting and
solidification was performed using a vacuum or
protective atmosphere. The characterization of the alloy
was carried out by optical microscopy (OM), scanning
electron microscopy (SEM) equipped with energy-
dispersive spectroscopy (EDS), differential scanning
calorimetry (DSC) and X-ray diffraction (XRD)
methods. For microstructural observations, the samples
were grinded (120–800 grade paper) and polished (0.5
μm Al2O3). Later on, the samples are etched in a solution
composed of 2.5 g FeCl3 and 48 ml methanol in 10 ml
HCl. The procedure for etching consisted of etching for
2 min, inter polishing for 2 min, and later etching for 1
min. The grain-size measurements were carried out by
OM using the grain cutting line method. Hardness tests
were carried out using the Vickers method (HV1). The
differential scanning calorimetry (DSC) measurements
were employed using a device under an argon atmo-
sphere in the temperature range from room temperature
to 300 °C. The rates of heating and cooling were
10 K/min. In order to determine the phase composition
the X-ray diffraction (XRD) measurements were per-
formed. CuK radiation was used.
3 RESULTS AND DISCUSSION
The chemical composition analysis of the alloy was
done using EDS analysis (Figure 1). The results showed
that the chemical composition of the alloy was 82.73 %
Cu, 13.16 % Al and 4.11 % Ni (mass fractions). The
microstructures of the lateral and longitudinal cross-
sections of the bar after continuous casting are characte-
ristic of a continuously casted bar (Figure 2). Under the
surface there was a layer of fine equiaxiad grains
followed by a region of long fringe crystals oriented
towards to the centre of the cross-section where again
equiaxiad grains appeared (Figures 2a and b). The
orientation of the crystals was changed. The fringe
crystals were formed at an angle with an average value
of around 60° with the longitudinal bar axis. In accord-
ing to Lojen et al.19 the maximum achieved velocity of
the solidification front was about 2.1 mm/s. It is possible
that the crystallization front that simultaneously proceeds
from the outer part of the bars to its centre causes a
preferential orientation of the growing grains. The
average hardness of the alloy was 275 HV1.
Figure 3 shows OM micrographs of the CuAlNi
alloy after the continuous casting procedure. The grain
boundaries are clearly visualized. As can be seen, the
micrographs of the specimens show the typical marten-
site microstructure. Martensite laths have different orien-
tations into particular grains. The grain size depends on
the place from which the samples were taken. The grain
diameter near to the external surface is higher than in the
center.
An average grain size of 98.78 μm was observed for
as-cast specimens. The number of grains per was 20.3
mm–2. For an average grain size from 50 μm to 100 μm
the fracture strain in the martensite phase is of the order
of 10 %, which is sufficient for shape-memory appli-
cations. The grain size of the rapidly solidified alloys is
determined by the amount of undercooling prior to the
crystallization. Our results for bars after continuous
casting showed that the sizes of grains vary from the
surface towards the center of the bars. The increase in
the grain diameter causes an increase of the Ms tem-
M. GOJI] et al.: MICROSTRUCTURAL AND PHASE ANALYSIS OF CuAlNi SHAPE-MEMORY ALLOY ...
150 Materiali in tehnologije / Materials and technology 47 (2013) 2, 149–152
Figure 1: EDS spectrum of CuAlNi shape-memory alloy
Slika 1: EDS-spekter spominske zlitine CuAlNi
Figure 2: a) Lateral and b) longitudinal cross-section of bar after
continuous casting
Slika 2: a) Pre~ni in b) vzdol`ni prerez kontinuirno ulite palice
perature. A similar behavior was already observed for
the Cu-Al-Ni-Mn and Cu-Al-Ni-Mn-Ti shape-memory
alloys20 obtained using the melt-spinning technique. The
temperature Ms depends on the grain size according to
the relation Ms  d1/2, where Ms is the difference
between the temperature Ms of the melt-spun ribbons
(small grains) and the bulk alloy (large grains) and d is
the mean grain diameter.
The martensitic microstructure was confirmed with
the SEM micrographs (Figure 4). This microstructure is
the result of the beta-phase of the Cu-Al-Ni alloys
transforming into the martensite phase by cooling below
the Ms temperature. The martensite is formed primarily
as the needle-like shape. In some fields the V-shape
M. GOJI] et al.: MICROSTRUCTURAL AND PHASE ANALYSIS OF CuAlNi SHAPE-MEMORY ALLOY ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 149–152 151
Figure 3: OM micrographs of the CuAlNi shape-memory alloy: a), b)
near the external surface and c), d) in the center
Slika 3: OM mikrostrukture spominske zlitine CuAlNi: a), b) v bli`ini
zunanje povr{ine in c), d) v centru
Figure 5: DSC curves of CuAlNi bar during cooling and heating
Slika 5: DSC-krivulje palic CuAlNi med ohlajanjem in ogrevanjem
Figure 4: SEM micrographs of the CuAlNi shape-memory alloy: a),
b) in longitudinal and c), d) lateral cross-section: a, c – SEI; b, d – BEI
Slika 4: SEM mikrostrukture spominske zlitine CuAlNi: a), b) v
vzdol`nem in c), d) pre~nem prerezu: a, c – SEI; b, d – BEI
martensite was observed. This is a typical self-accommo-
dating, zig-zag, martensite morphology, which is
characteristic for the 1’ martensite in the CuAlNi
alloy21. The parallel bands in the martensite can be
considered as twin-like martensite.
Figure 5 shows DSC curves during heating and
cooling. The temperature phase transformation are: Ms =
215 °C, Mf = 169 °C, As = 229 °C and Af = 191 °C. It can
be seen that the CuAlNi alloy is a possible candidate for
high-temperature usage. Figure 6 shows the X-ray
profile of the as-casted alloy. Our XRD analysis showed
that the alloy after continuous casting consisted of the
parent 1 and 1’ martensite phases.
4 CONCLUSIONS
The microstructure of a CuAlNi shape-memory alloy
after continuous casting consisted of the parent 1 and 1’
martensite phases. The martensite laths have different
orientations into particular grains. The grain size
depends on the place from which the samples are taken.
The grain diameter near to the external surface is higher
than in the center. It was found that the average grain
size was 98.78 μm in particular grains. The temperature
phase transformations determined by DSC measure-
ments were: Ms = 215 °C, Mf = 169 °C, As = 229 °C and
Af = 191 °C. The average hardness of the alloy was 275
HV1.
Acknowledgements
This work was supported by EUREKA Project E!
3704 RSSMA "Rapidly Solidified Shape Memory
Alloys" by the Ministry of Science, Education and
Sports of the Republic of Croatia.
5 REFERENCES
1 M. Goji}, Metalurgija, 31 (1992) 2/3, 77–82
2 K. Otsuka, C. M. Wayman, Shape memory alloys, Cambridge Uni-
versity Press, Cambridge, 1998, 97–116
3 D. ]ori}, M. Franz, Welding, 50 (2007) 5/6, 179–187
4 N. M. Lohan, B. Pricop, J. G. Bujoreanu, N. Cimpoesu, Int. J. Mat.
Res., 102 (2011) 11, 1345–1351
5 H. Kubo, H. Otsuka, S. Farjami, T. Maruyama, Scripta Materialia, 55
(2006) 11, 1059–1062
6 K. Otsuda, X. Ren, Progress in Materials Scienc, 50 (2005), 511–678
7 T. Hu, L. Chen, S. L. Wu, C. L. Chu, L. M. Wang, K. W. K. Yeung,
P. K. Chu, Scripta Materialia, 64 (2011), 1011–1014
8 M. Goji}, S. Ko`uh, B. Kosec, I. An`el, Properties and Application
of Shape Memory Alloys, Proceedings 9th Scientific Research Sym-
posium with International Participation: Metallic and Nonmetallic
Materials-production-properties and application, University of
Zenica, Faculty of Metalurgy and Materials Science, Zenica, 2012,
13–26
9 C. E. Sobrero, P. La Roca, A. Roatta, R. E. Bolmaro, J. Mlarría,
Materials Science and Engineering A, 536 (2012), 207–215
10 V. Sampath, Effect of Zr on Microstructure and Transformation
Temperatures of Cu-Al-Ni Shape memory Alloys, Proc. Int. Conf. on
Smart Materials, Bangalore, India, 2005, SC-148–156
11 P. Zhang, A. Ma, S. Lu, G. Liu, P. Lin, J. Jiang, C. Chu, Materials
and Design, 32 (2011), 348–352
12 J. V. Wood, P. H. Shingu, Metallurgical Transactions, 15A (1984) 4,
471–480
13 M. Izadinia, K. Dehghani, Trans. Nonferrous Met. Soc. China, 21
(2011), 2037–2043
14 G. N. Sure, L. C. Brown, Metallurgical Transactions, 15A (1984) 8,
1984–1613
15 A. C. Kneissl, E. Unterweger, G. Lojen, Advanced Engineering
Materials, 8 (2006) 11, 1113–1118
16 T. Goryczka, Archives of Metallurgy and Materials, 54 (2009) 3,
755–763
17 M. Goji}, L. Vrsalovi}, S. Ko`uh, A. Kneissl, I. An`el, S. Gudi}, B.
Kosec, M. Kli{ki}, Journal of Alloys and Compounds, 509 (2011),
9782–9790
18 Z. Wang, X. F. Liu, J. X. Xie, Materials Science and Enginnering A,
532 (2012), 536–542
19 G. Lojen, A. C. Kneissl, M. Goji}, R. Rudolf, M. ^oli}, I. An`el,
Livarski vestnik, 57 (2010) 4, 172–193
20 J. Dutkiewicz, T. Czeppe, J. Morgiel, Materials Scinece and Engi-
neering A, 573–275 (1999), 703–707
21 U. Sari, I. Aksoy, Journal of Alloys and Compaunds, 417 (2006),
138–142
M. GOJI] et al.: MICROSTRUCTURAL AND PHASE ANALYSIS OF CuAlNi SHAPE-MEMORY ALLOY ...
152 Materiali in tehnologije / Materials and technology 47 (2013) 2, 149–152
Figure 6: XRD spectrum of CuAlNi shape-memory alloy after con-
tinuous casting
Slika 6: XRD-spekter spominske zlitine CuAlNi po kontinuirnem litju
O. CULHA et al.: CHARACTERIZATION AND DETERMINATION OF MECHANICAL PROPERTIES OF YBCO ...
CHARACTERIZATION AND DETERMINATION OF
MECHANICAL PROPERTIES OF YBCO
SUPERCONDUCTING THIN FILMS WITH MANGANESE
USING THE TFA-MOD METHOD
KARAKTERIZACIJA IN DOLO^ITEV MEHANSKIH LASTNOSTI
SUPERPREVODNE TANKE PLASTI YBCO Z MANGANOM PO
METODI TFA-MOD
Osman Culha1, Isil Birlik2, Mustafa Toparli2, Erdal Celik2, Sebastian Engel3, Bernhard Holzapfel3
1Celal Bayar University, Department of Materials Engineering, Muradiye Campus, Manisa, Turkey
2Dokuz Eylul University, Department of Metallurgy and Materials Engineering, Tinaztepe Campus, Buca, Izmir, Turkey
3Leibniz-Institut für Festkörper- und Werkstoffforschung (IFW), Solid-State and Materials Research, Helmholtzstraße 20, 01069 Dresden,
Germany
isil.kayatekin@deu.edu.tr
Prejem rokopisa – received: 2012-07-26; sprejem za objavo – accepted for publication: 2012-10-04
The aim of this study is to determine the microstructure, superconducting and mechanical properties of YBa2Cu3O6.56 (YBCO)
and YBCO thin films with a manganese (Mn) addition. All the YBCO superconducting films (undoped and Mn-doped) were
dip-coated onto (001) SrTiO3 (STO) single-crystal substrates with a metalorganic deposition using the trifluoroacetate
(TFA-MOD) technique. The phase analysis, microstructure, surface morphologies and critical temperature (Tc) of the
superconducting thin films were determined with an X-ray diffractometer (XRD), a scanning electron microscope (SEM), an
atomic force microscope (AFM) and an inductive Tc measurement system. Since the main issue of this study is to determine the
mechanical-property variations of the superconducting thin films with/without a Mn addition, the adhesion strength of these
films on a STO substrate was tested with a Shimadzu scratch tester. Depending on the Mn addition, the critical forces of pure
films increase from 56.23 mN, 58.63 mN and 60.11 mN for pure YBCO, YBCO with 0.05 g and 0.10 g of Mn. Furthermore,
Young’s modulus and the hardness of the undoped and Mn-doped YBCO thin films were measured with a CSM Berkovich
nanoindenter using the load-unload sensing analysis under a 0.3 mN applied load.
Keywords: superconducting films, sol-gel synthesis, mechanical properties, nanoindentation
Namen te {tudije je dolo~iti mikrostrukturo, superprevodne in mehanske lastnosti YBa2Cu3O6,56 (YBCO) in tanke plasti YBCO z
dodatkom mangana (Mn). Vse superprevodne plasti YBCO (nedopirane in dopirane z Mn) so bile nanesene na monokristalni
substrat (001) SrTiO3 (STO) s kovinoorganskim nanosom s trifluoracetatno tehniko (TFA-MOD). Fazna analiza, mikrostruktura,
morfologija povr{ine in kriti~na temperatura (Tc) superprevodne tanke plasti so bile dolo~ene z rentgenskim difraktometrom
(XRD), vrsti~nim elektronskim mikroskopom (SEM), mikroskopom na atomsko silo (AFM) in induktivnim merilnim sistemom
za Tc. Ker je bila glavna naloga te {tudije dolo~anje spreminjanja mehanskih lastnosti YBCO z dodatkom Mn in brez njega, je
bila preizku{ena adhezivnost te plasti na STO-podlago s Shimadzu preizku{evalnikom za razenje. Odvisno od dodatka Mn
kriti~na sila v ~isti plasti nara{~a od 56,23 mN, 58,63 mN in 60,11 mN za ~isti YBCO, YBCO z 0,05 g in 0,10 g Mn. Poleg tega
sta bila izmerjena tudi Youngov modul in trdota nedopirane in z Mn dopirane tanke plasti YBCO z nanomerilnikom trdote CSM
Berkovich z analizo zaznavanja obremenjeno – neobremenjeno pri uporabljeni obremenitvi 0,3 mN.
Klju~ne besede: superprevodna plast, sol-gel sinteza, mehanske lastnosti, nanopreizku{anje trdote
1 INTRODUCTION
After the discovery of high-Tc superconductors
(HTS), their various applications have been tested in
various areas. These materials are supposed to increase
the performance of devices such as magnetic resonance
imaging (MRI) in medicine, energy-storage systems in a
transformer, magnetic separators, levitation, nuclear
magnetic resonance (NMR), generators, engines, cables,
superconducting wires and tapes, accelerators, electro-
magnets, electronic transistors and bolometers.1,2
HTS thin films with a sharp resistive transition, high
critical current density Jc and low flux noise offer the
potential for such applications.3 Extensive studies are
currently being carried out worldwide on YBa2Cu3O6.56
(YBCO) films grown on different single-crystal and
metal-based substrates4. Many YBCO thin films have
been developed using different deposition processes.
Most of them use high vacuum techniques such as pulse
laser deposition (PLD) and magnetron sputtering that
can create high critical current densities on YBCO thin
films. Nevertheless, they require significant start-up
costs for long-length coated-conductor production5. On
the other hand, thin films prepared with non-vacuum
techniques like metal-organic decomposition using
triflouroacetic acid (TFA-MOD) which is a sol-gel-
related method, show similar superconducting properties
and are relatively simple and inexpensive6. High-quality
YBCO films with high Jc can be fabricated in a
TFA-MOD process7. Finding the optimum process para-
meters for a coating solution can be challenging, but
once the coating solution is found, it is very easy to
obtain high Jc YBCO superconductors with supreme
reproducibility. Although the TFA-MOD process using
Materiali in tehnologije / Materials and technology 47 (2013) 2, 153–160 153
UDK 532.6:620.17 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)153(2013)
metal acetates as the starting materials is more cost
effective than vacuum processes, the highly purified
metal acetates are expensive and thus it is desirable to
find a more economic route. Recently, several attempts
to use oxide powders, such as the commercially available
YBCO powder, as the starting materials have been
reported and they have shown high Jc for the YBCO
films compared to the other methods.8
In this paper we present a new approach by com-
bining the superior properties of a solvent, especially
2,4-pentanedionate, and a commercially available YBCO
powder with TFA, acetone and propionic acid as the
preliminary study. Therefore, YBCO superconducting
films were produced with BaMnO3 from the solutions
prepared with the cheap and commercially available
YBCO powders and Mn 2,4-pentanedionate, TFA pro-
pionic acid and acetone. In practical applications, the
importance of the mechanical properties of YBCO-based
films cannot be ignored. Superconducting films with
poor mechanical properties are useless, even if they
possess good transport and flux-pinning properties.
Since an addition of particles as pinning centers results
in important changes on the microstructure of thin films,
their effect on micromechanical properties such as
Young’s modulus, hardness and adhesion strength also
have to be investigated with respect to the particles type
and quantity.
2 THEORETICAL BACKGROUND FOR
MECHANICAL PROPERTIES
The scratch test is used to measure the interfacial
adhesion for a range of different coatings.9–19 During the
test, a diamond indenter is drawn over the surface of a
sample tested under a normal force, which is increased
either stepwise or continuously until the critical normal
force is reached, at which a well defined coating failure
occurs.9 Then, force is taken as a measurement of the
adhesion between the coating and the substrate. The
onset of a coating failure can be monitored with optical
microscopy, acoustic emission (AE) and friction-force
measurements10. It has been suggested that the coating
adhesive failure is directly associated with a sudden
increase in the friction force.11,12 It is generally accepted
that the scratch test is suitable for the coatings of a
thickness ranging from 0.1 μm to 20 μm and this covers a
large number of engineering applications.13
The adhesion strength (F) of the films was calculated
by using equation (1) for the average value of three
measurements. The scratch was examined with an
optical microscope and the critical load, at which the
coating was removed, was determined. The adhesion
strength, F (in MPa), was calculated by using equation
(1):20
[ ]F
H
R H W W
=
−( ) /
/
	 2
1 2
c c
(1)
where H is the Brinell hardness value (kg mm–2) of the
SrTiO3 (STO) substrate and R is the radius of the stylus
(μm) and Wc is the critical force.
On the other hand, indentation tests are used to deter-
mine elasto-plastic properties such as Young’s modulus,
yield strength and the strain-hardening exponent21 of thin
films. Young’s modulus may be inferred from the
unloading indentation load-depth curve and the yield
strength from the maximum indentation load. In
addition, a method to extract the flow stress and the
strain-hardening exponent using indentation data has
been researched.22–27 The indentation hardness of mate-
rials is measured in several ways by forcing an indenter
having a specific geometry (ball, cone, and pyramid) into
a specimen’s surface.28 The conventional microhardness
value can be determined with an optical measurement of
the residual impression left behind upon a load release.
The development of depth-sensing indentation equip-
ment has allowed an easy and reliable determination of
two of the most commonly measured mechanical
properties of materials: the hardness and Young’s
modulus.26,29
Two mechanical properties, namely, elastic modulus
(E) and microhardness (H) can be obtained from the load
and penetration-depth data. A typical load-penetration-
depth curve can be investigated using a related
reference.30 During an indenter loading, the test material
is subjected to both elastic and plastic deformations. One
of the challenges in studying mechanical properties of
thin films is that the traditional methods used to evaluate
mechanical properties of bulk materials are not
applicable for thin films and so far there is no standard
test method for the evaluation of mechanical properties
of thin films.31 New methods, such as depth-sensing
nanoindentation, microbridge test, uniaxial tensile test
and ultrasonic method are being developed32–35 for the
measurement of mechanical properties of thin films.
Among these methods, the depth-sensing nanoinden-
tation technique provides a continuous record of the
variation of indentation load with the penetration depth
into a specimen and this technique has been an area of
considerable attention in recent years due to its high
resolution at a low load scale. Currently, the nano-
indentation technique is being applied to determine
hardness and Young’s modulus, while limited studies are
available to develop an effective method to obtain the
plastic properties of thin films with a nanoindentation
technique. For example, Nix34 has utilized a nanoinden-
tation technique to study the strength properties of thin
films. Giannakopoulos and Suresh32 have developed a
step-by-step method to obtain the mechanical properties
of materials from the nanoindentation experimental data.
The indentation response of a thin film on a substrate is a
complex function of the elastic and plastic properties of
both the film and the substrate and it is essential to
understand how the intrinsic mechanical properties of a
film can be determined from the overall mechanical
response of a film/substrate system. As the values of
O. CULHA et al.: CHARACTERIZATION AND DETERMINATION OF MECHANICAL PROPERTIES OF YBCO ...
154 Materiali in tehnologije / Materials and technology 47 (2013) 2, 153–160
elastic modulus and hardness, determined from inden-
tations, should not depend on the value of h (indentation
depth) and, therefore, on the value of the maximum load,
the indentation depth should not exceed 10–20 % of the
coating thickness, otherwise the results will be affected
by the properties of the substrate.21–30
3 EXPERIMENTAL PROCEDURE
3.1 Preparation of the solutions
Y-Ba-Cu-O based solution was prepared from a
commercial YBa2Cu3O6.56 powder (yttrium-barium-
copper oxide) with propionic acid, trifluoroacetic acid
(TFA), acetone and 2,4-pentanedionate under atmo-
spheric conditions at room temperature. The 8.3045 g
YBCO powder was weighted out in order to prepare a
0.25 M and 50 ml solution. After 10 ml of propionic acid
was added to the YBCO powder, the mixture was being
dissolved in 25 ml of TFA at 45 °C for 60 min using an
ultrasonic mixer (Sonorex digital 10P). The solvents
were being removed at 100 °C for 60 min with a hot
plate to yield a blue-sticky-glassy residue with a high
viscosity. The solution of the film was made by dissolv-
ing the residue with up to 50 ml of TFA. The solvents
were being evaporated from the solution again at 60 °C
for 60 min until a highly viscous solution of a trans-
parent-blue colour was obtained. After adding up to 50
ml of acetone into the solution, 15 ml of propionic acid
and 5 ml of TFA were incorporated into the obtained
viscous solution, and then a standard transparent solution
was prepared. Finally, Mn alkoxide was separately added
into 5 ml of the transparent solution with a low content
of 2,4-pentanedionate as presented by the flow chart in
Figure 1.
3.2 Coating process
Initially, (100) STO single-crystal substrates with the
dimensions of 10 mm × 10 mm × 0.75 mm were rinsed
in acetone using a standard ultrasonic cleaner. After that,
the solutions were deposited on the substrates during a
dip-coating process with a withdrawal speed of 0.3 cm/s
in a vacuum atmosphere. The dip coating involves a
formation of a film through a liquid-entrainment process
that may be either batch or continuous in nature. The
general steps include an immersion of the substrate into
the dip-coating solution, a start-up, during which a
withdrawal of the substrate from the solution begins, a
film deposition, a solvent evaporation, and a continued
drainage as the substrate is completely removed from the
liquid bath. The film thickness formed in dip coating is
mainly governed by the viscous drag, gravitational
forces, and the surface tension.
The deposited gel films were converted to an
epitaxial pure YBCO and a YBCO film with BaMnO3
nanoparticles through a combination of the calcining and
heat-treatment procedures. The gel film was dried from
80 °C to 406 °C in 12 % humidified oxygen. After the
calcining was performed at 406 °C for 0.4 h in 12 %
humidified nitrogen, the film was being heated up to 811
°C for 1 h in 12 % humidified nitrogen and the fired film
was consequently heat treated at 465 °C for 2.2 h in a
dry oxygen atmosphere. In order to obtain highly
textured thin films on the STO substrate, the oxygen
content was 500 ml O2 during the heat-treatment process
at 465 °C for 2.2 h as expressed in Figure 1.
O. CULHA et al.: CHARACTERIZATION AND DETERMINATION OF MECHANICAL PROPERTIES OF YBCO ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 153–160 155
Figure 1: Flow chart of the TFA-MOD technique
Slika 1: Potek postopka TFA-MOD
3.3 Characterization process
The structural development of the produced thin
films was investigated using X-ray diffraction (XRD-
Rigaku D/MAX-2200/PC) patterns, recorded using the
Co K irradiation (wavelength, 
 = 0.178897 nm) and
the scanning range was between 2 = 10° and 90°. The
surface topographies and additional particle effects on
the microstructure of the films were examined with a
high resolution SEM (JEOL JSM 6060). The surface
morphologies of the films were measured with an atomic
force microscope (AFM). The resistivity-temperature
behaviors of the superconductive thin films were
determined and the critical temperatures, Tc, were
obtained depending on the additional particle content.
3.4 Mechanical properties of the films
The load on a Rockwell C diamond with a tip radius
of the stylus (R) of 15 μm was linearly increased from 0
mN to 98 mN at the loading speed of 1 mN/s for the
scratch test. In addition, the scratch speed of the
diamond tip was 2 μm/s. The testing temperature and
humidity percentage were 20.3 °C and 50 %, respec-
tively. The scratch was examined with an optical micro-
scope and the critical force (Wc) value, at which the
coating removed from the substrate was determined three
times for each sample. The hardness and Young’s
modulus of the produced thin films were measured under
0.3 mN of the applied peak load three times with the
CSM Berkovich nanoindentation tester (the loading-
unloading test mode) to determine additional particle
effects on mechanical properties.
4 RESULTS AND DISCUSSION
4.1 Characterization of the films
Figure 2 (a–c) shows the XRD patterns of pure
YBCO, YBCO with 0.05 g and 0.10 g Mn on the STO
single-crystal substrate obtained with the TFA-MOD
method. XRD patterns showed that the pure YBCO film
has (001) major diffraction peaks corresponding to the
O. CULHA et al.: CHARACTERIZATION AND DETERMINATION OF MECHANICAL PROPERTIES OF YBCO ...
156 Materiali in tehnologije / Materials and technology 47 (2013) 2, 153–160
Figure 2: XRD patterns of: a) pure YBCO, b) YBCO with 0.05 g and
c) 0.10 g BaMnO3 nanoparticles
Slika 2: Rentgenski pra{kovni posnetki (XRD): a) ~isti YBCO, b)
YBCO z 0,05 g in c) 0,10 g BaMnO3 nanodelci
Figure 3: AFM study of YBCO thin films with: a) 0 g, b) 0.05 g and
c) 0.10 g BaMnO3 nanoparticles
Slika 3: AFM-{tudija YBCO tanke plasti z: a) 0 g, b) 0,05 g in c) 0,10
g BaMnO3 nanodelci
(00l) parallel plane. It is worth mentioning that YBCO
films with a high intensity were grown on the STO
substrate using the TFA-MOD method. In addition, no
second phases such as Y2Cu2O5, BaF2 and CuO were
found in the case, where only a pure YBCO phase was
formed. Apart from that, BaMnO3 perovskite peaks with
a low intensity were determined on account of the
Mn-doping effect. However, XRD peaks from the
BaMnO3 perovskite second phase were hardly detected
even for the film that was prepared using a precursor
solution containing a 0.10 g Mn dopant.
The effects of BaMnO3 nanoparticles on the micro-
structure of YBCO thin films were identified in Figure
3. According to the AFM results, all the films show
typical CuOx precipitatates with 100–200 nm diameters,
which are regularly found on the surface of TFA-MOD
samples. When the Mn content increased from 0 g to
0.10 g, the surface-roughness values of pure YBCO thin
films and the YBCO thin films with BaMnO3 nano-
particles changed from 21 nm to 33 nm depending on the
Mn doping content. Nevertheless, a clear change in the
YBCO surface was observed.
Figures 4, 5 and 6 show SEM micrographs of pure
YBCO and the YBCO films with BaMnO3 nanoparticles.
Figure 4 shows surface topographies of the pure YBCO
film with the a-axis which decrease the superconducting
properties. These films were produced using the standard
YBCO transparent solution. However, the decreasing
superconducting properties with the a-axis were elimi-
nated using the Mn doping in the standard YBCO
precursor solution as clearly depicted in Figures 5 and 6.
SEM micrographs indicate that structural defects can be
reacted as nanodots or nanoparticles (due to the Mn
addition) along the c-axis of the YBCO film. These
properties result in an enhanced pinning over the pure
YBCO film. Since Mn reacts with Ba and a BaMnO3
perovskite structure forms in the YBCO film during the
heat-treatment process, the microstructures of supercon-
ducting thin films were changed as expected.
The dependences of the inductively measured critical
transition temperature (Tc) and transition width (Tc) on
the amount of BaMnO3 in the structure are shown in
Figure 7. It can be seen that there is sharp decrease in
the resistivity near 90 K where the critical-temperature
value of pure YBCO is 90.4 K. When the quantity of
additional particles increases from 0.05 g to 0.10 g, the
critical temperature changes from 90.2 K to 90 K. These
additional particles do not affect the critical temperature
and only behave as impurities or second phases for the
flux-pinning properties.
O. CULHA et al.: CHARACTERIZATION AND DETERMINATION OF MECHANICAL PROPERTIES OF YBCO ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 153–160 157
Figure 6: Surface morphology of an YBCO film with BaMnO3
nanoparticles produced using the YBCO transparent solution with
0.10 g Mn doping
Slika 6: Morfologija povr{ine YBCO tanke plasti z BaMnO3 nano-
delci, izdelane s prozorno raztopino YBCO, dopirano z 0,10 g Mn
Figure 4: Surface morphology of a pure YBCO film with the a-axis
produced using the standard YBCO transparent solution
Slika 4: Morfologija povr{ine ~iste plasti YBCO z a-osjo, izdelane s
standardno prozorno raztopino YBCO
Figure 5: Surface morphology of an YBCO film with BaMnO3 nano-
particles produced using the YBCO transparent solution with 0.05 g
Mn doping
Slika 5: Morfologija povr{ine YBCO tanke plasti z BaMnO3 nano-
delci, izdelane s prozorno raztopino YBCO, dopirano z 0,05 g Mn
4.2 Determination of mechanical properties
The analysis of the scratch test gives critical-force
values of the produced thin films. These critical-force
values correspond to the first peak in the cartridge output
percentage – the test force graphical curve obtained from
the scratch-test machine. The bond strengths, critical-
force and Brinell-hardness values of pure YBCO and
YBCO with BaMnO3 nanoparticles are presented in
Table 1. The hardness value of the substrate was con-
verted to Brinell hardness (H) with the Standard
Hardness Conversion Tables for Metal and the adhesion
strength (F) of the coatings was calculated as a MPa unit
using equation (1).20 It is clearly seen from Table 1 that
pure YBCO, the YBCO thin films with 0.05 g and 0.1 g
of BaMnO3 nanoparticles have the critical forces of
56.23 mN, 58.63 mN and 60.11 mN, respectively.
Therefore, the calculated adhesion strength of the films
increases from 160 MPa to 173 MPa depending on the
BaMnO3 formation in the microstructure.
As expected, the mechanical strength against the
internal forces caused by the flux pinning rather than the
magnitude of the critical current density, may become
the factor limiting the performance of HTS thin films for
high-current density applications. For the bulk pure
YBCO material, this problem has been already discussed
in 36. In this study, an increasing Jc and an improvement
in the flux-pinning properties of the YBCO films with
the BaMeO3 perovskite nanoparticles (Me : Mn), as the
pinning centers, on SrTiO3 (STO) are aimed at with the
trifluoroacetic acid-metal organic deposition (TFA-
MOD). Since the general purpose is to determine the
mechanical properties such as Young’s modulus and the
hardness of a pure YBCO thin film and the YBCO thin
films with a Mn addition (Mn reacts as BaMnO3), an
instrumented nanoindentation test was applied to both
samples. This time, the importance of the applied load
and the indentation size is apparent. As presented in
Table 2, the ratios of the indentation depths and the film
thicknesses are applicable for an instrumented inden-
tation with an applied load of 300 μN. A smoothing
procedure was applied to all of the instrumented
indentation results of the samples. Another parameter,
which can affect the indentation test, is the surface
roughness. It has a very active role in an indentation
experiment at the nanoscale. If the surface roughness is
larger than the maximum indentation depth, the curve
has very scattered data characteristic for the loading and
unloading parts.
O. CULHA et al.: CHARACTERIZATION AND DETERMINATION OF MECHANICAL PROPERTIES OF YBCO ...
158 Materiali in tehnologije / Materials and technology 47 (2013) 2, 153–160
Table 1: Brinell hardness, critical force and bond strengths of YBCO
and YBCO with Mn
Tabela1: Trdota po Brinellu, kriti~na sila in trdnost vezave YBCO in
YBCO z Mn
Sub-
strate
Brinell
hardness
substrate
(HB)
Indenter
radius
(μm)
Material
Average
critical
force
(mN)
Adhesion
strength
(MPa)
STO 143 15
Pure YBCO 56 ± 1.2 160 ± 1.1
YBCO with
0.05 g Mn 59 ± 1.8 168 ± 1.7
YBCO with
0.1 g Mn 60 ± 1.6 173 ± 1.5
Table 2: Maximum depth, residual depth and film thickness of pure
YBCO and YBCO with additional particles
Tabela 2: Maksimalna globina, preostala globina in debelina plasti
~istega YBCO in YBCO z dodatnimi delci
Material Force(μN)
Maximum
depth
(nm)
Residual
depth
(nm)
Film
thickness
(nm)
YBCO 300 40.24 ± 6.4 30.12 ± 7.8 292 ± 9
YBCO with
0.05 g Mn 300 42.32 ± 9.4 32.11 ± 8.7 297 ± 5
YBCO with
0.10 g Mn 300 42.45 ± 5.9 33.53 ± 4.4 294 ± 6
Figure 7: Critical-temperature values of YBCO thin films with: a) 0 g,
b) 0.05 g and c) 0.10 g BaMnO3 nanoparticles
Slika 7: Vrednosti kriti~ne temperature YBCO tanke plasti z: a) 0 g,
b) 0,05 g in c) 0,10 g BaMnO3 nanodelci
The YBCO-based thin film has suitable properties for
ceramic materials, such as hardness and stiffness,
together with the tendency to fracture. However, the
references about the mechanical properties of this
material, particularly the yield strength and the stress-
strain curve, are scarce. The mechanical properties
(hardness, Young’s modulus and fracture toughness) of
the YBCO samples have been examined with the tech-
niques such as ultrasound37, X-ray diffraction38 and
nanoindentation39. The reported values of Young’s
modulus for Y-123 are within the range of E = 40–200
GPa. This large scatter may be due to the residual
porosity and a poor contact between the grains40. Other
authors,41 also using nanoindentation and applying the
loads between 30 mN and 100 mN, reported the values
of E = 171–181 GPa for the YBCO samples textured
with the Bridgman technique, which is in agreement
with Johansen. The nanohardness values in the range of
7.8–8.0 GPa at the maximum loads of 30 mN were
recently reported in several researches for the bulk,
single-crystal YBCO.38–41 Roa et al. (2007) found a hard-
ness value of (8.9 ± 0.1) GPa obtained with a nano-
indentation on the YBCO samples textured with the
Bridgman technique.
After obtaining the loading-unloading (load-displace-
ment) curves, force-time and depth-time graphs for pure
YBCO and the YBCO films with 0.05 g and 0.1 g Mn
(reacting as BaMnO3) under a applied peak load of 300
μN, some characteristic indentation parameters were
listed in Tables 2 and 3. According to the indentation
results in Table 2, the maximum and residual depths of
the YBCO-based thin films under the same indentation
loads were increased from (40.24 ± 6.4) nm to (42.45 ±
5.9) nm and from (30.12 ± 7.8) nm to (33.53 ± 4.4) nm,
respectively. Furthermore, as represented in Table 3, the
elastic modulus and indentation hardness of the
YBCO-based thin films were decreased from (88.54 ±
3.1) GPa to (79.11 ± 1.9) GPa and from (12.51 ± 5.1)
GPa to (5.75 ± 1.1) GPa under a load of 300 μN.
Although the indentation load was fixed at 300 μN, the
indentation hardness was decreased as if the indentation
size affected the mechanical properties. Since the
hardness is accepted as an inherent material property, it
should not vary with the indentation load and size but
may change with different phase formations. The
indentation hardness was decreased with the formation
of the BaMnO3 content in the YBCO thin-film structure.
According to this explanation, it can be concluded that
the pure YBCO thin film is harder and more brittle than
the BaMnO3 additional ones. As listed in Table 3, the
indentation hardness of the YBCO-based thin films
decreased with an increased Mn content in the structure.
The elastic-modulus variation of the YBCO-based thin
films is also shown in Table 3. Although the hardness
was very sensitive to the maximum indentation depth
and the thickness/indentation-depth ratio of the samples
and it changed from 12.51 GPa to 5.75 GPa, the elastic
modulus of the YBCO-based thin films did not show a
sharp decrease. However, as listed in Table 3, the elastic
modulus of pure YBCO, YBCO with 0.05 g Mn, and
YBCO with 0.10 g Mn was calculated to be 88.54 GPa,
83.41 GPa, and 79.11 GPa, respectively. After consider-
ing the hardness and elastic modulus of the YBCO-based
thin films with AFM, SEM and the critical temperature
value, Tc, it was found that the second phase of BaMnO3
did not continuously improve the mechanical and super-
conducting properties. When the Mn content increased
from 0 g to 0.15 g, the BaMnO3 phases occurred and
made the structure more ductile than the pure one.
According to the Tc measurement, a sharp decrease in Tc
could be considered and the Tc values of the
YBCO-based films were very small up to the addition of
0.15 g Mn. In this case Tc was 4.2 K for the YBCO thin
film with a 0.15 g Mn addition (BaMnO3) and a sharp
decrease in the resistivity could not be seen easily as
presented Figure 7.
Table 3: Indentation-experiment results for pure YBCO and YBCO
with additional particles
Tabela 3: Preizkus trdote ~istega YBCO in YBCO z dodanimi delci
Material Force(μN)
Hardness
(HV)
Indentation
hardness
(GPa)
Young’s
modulus
(GPa)
YBCO 300 695 ± 28 12.51 ± 4.8 88.54 ± 3.1
YBCO with
0.05 g Mn 300 525 ± 16 8.21 ± 1.2 83.41 ± 1.8
YBCO with
0.10 g Mn 300 495 ± 21 5.75 ± 1.1 79.11 ± 1.9
5 CONCLUSION
In this study, YBCO films with/without the Mn
solutions were prepared by dissolving commercially
available YBCO powders in propionic acid, trifluoro-
acetic acid (TFA) and acetone. The prepared solutions
were dip coated on the STO single-crystal substrates.
The following microstructural and mechanical results
were obtained:
• XRD patterns show that the produced YBCO films
have (001) and parallel plane reflections for pure
YBCO and the YBCO with the BaMnO3 thin film.
• According to the AFM study, topographic properties
and the roughness of thin films were increased with
an increase in the amount of nanoparticles (acting as
a pinning center) in the film structure.
• SEM micrographs indicate that the structural defects
consisted of the nanodots or nanoparticles of
BaMnO3 along the c-axis of the YBCO film. These
properties resulted in an enhanced pinning over the
pure YBCO film.
• It can be seen from the Tc analysis that there is a
sharp decrease in the resistivity near 90 K.
• Critical force values of pure YBCO and of the
YBCO-based thin films with 0.05 g and 0.1 g
BaMnO3 were found to be (56.23, 58.63 and 60.11)
O. CULHA et al.: CHARACTERIZATION AND DETERMINATION OF MECHANICAL PROPERTIES OF YBCO ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 153–160 159
mN, respectively. The calculated adhesion strength of
the films increased from 160 MPa to 173 MPa
depending on the BaMnO3 addition.
• The calculated Young’s modulus of YBCO thin films
decreased with the BaMnO3 formation. The same
effects can be seen in the hardness variation of pure
YBCO and the YBCO thin films with the BaMnO3
nanoparticles.
6 REFERENCES
1 N. Hari Babu, K. Iida, D. A. Cardwell, Enhanced magnetic flux
pinning in nanocomposite Y-Ba-Cu-O superconductors, Physica C,
445–448 (2006), 353–356
2 Y. Yoshida, K. Matsumoto, M. Miura, Y. Ichino, Y. Takai, A. Ichi-
nose, M. Mukaida, S. Horii, Controlled nanoparticulate flux pinning
structures in RE1+xBa2–xCu3Oy films, Physica C, 445–448 (2006),
637–642
3 B. Lakew, J. C. Brasunas, S. Aslam, D. E. Pugel, High Tc, transition-
edge superconducting (TES) bolometer on a monolithic sapphire
membrane-construction and performance, Sensor. Actuat. A-Phys.,
114 (2004), 36–40
4 B. Dwir, L. Pavesi, J. H. James, B. Keilelt, D. Pavuna, F. K. Reinhart,
A simple high temperature superconducting thin film optical bolo-
meter, Supercond. Sci. Technol., 2 (1989), 314–316
5 Y. A. Jee, M. Li, B. Ma, V. A. Maroni, B. L. Fisher, U. Balachan-
dran, Comparison of texture development and superconducting pro-
perties of YBCO thin films prepared by TFA and PLD processes,
Physica C, 356 (2001), 297–303
6 Y. Yamada, S. Kim, T. Araki, Y. Takahashi, T. Yuasa, H. Kurosaki,
Critical current density and related microstructures of TFA-MOD
YBCO coated conductors, Physica C, 357 (2001), 1007–1010
7 X. M. Cui, B. W. Tao, J. Xiong, X. Z. Liu, J. Zhu, Y. R. Li, Effect of
annealing time on the structure and properties of YBCO films by the
TFA–MOD method, Physica C, 432 (2005), 147–152
8 S. Y. Lee, S. A. Song, B. J. Kim, J. A. Park, H. J. Kim, G. W. Hong,
Effect of precursor composition on Jc enhancement of YBCO film
prepared by TFA-MOD method, Physica C, 445–448 (2006),
578–581
9 J. Lia, W. Beres, Three-dimensional finite element modelling of the
scratch test for a TiN coated titanium alloy substrate, Wear, 260
(2006), 1232–1242
10 P. Hedenqvist, S. Hogmark, Experiences from scratch testing of tri-
bological PVD coatings, Tribol. Int., 30 (1997), 507–516
11 J. Valli, U. Mäkelä, Applications of the scratch test method for coat-
ing adhesion assessment, Wear, 115 (1987), 215–221
12 J. Sekler, P. A. Steinmann, H. E. Hintermann, The scratch test: diffe-
rent critical load determination techniques, Surf. Coat. Technol., 36
(1988), 519–529
13 K. Holmbert, A. Matthewst, H. Ronkainen, Coatings tribology-con-
tact mechanisms and surface design, Tribol. Int., 31 (1998), 107–120
14 K. Holmbert, The basic material parameters that control friction and
wear of coated surfaces under sliding, Tribologia-Fin. J. Tribol., 19
(2000), 3–18
15 P. A. Steinmann, Y. Tardy, H. E. Hintermann, Adhesion testing by
the scratch test method: the influence of intrinsic and extrinsic para-
meters on the critical load, Thin Solid Films, 154 (1987), 333–349
16 P. J. Burnett, D. S. Rickerby, The relationship between hardness and
scratch adhesion, Thin Solid Films, 154 (1987), 403–416
17 S. J. Bull, Failure modes in scratch adhesion testing, Surf. Coat.
Technol., 50 (1991), 25–32
18 S. J. Bull, Spallation failure maps from scratch testing, Mater. High
Temp., 13 (1995), 169–174
19 S. J. Bull, Failure mode maps in the thin film scratch adhesion test,
Tribol. Int., 30 (1997), 491–498
20 S. T. Gonczy, N. Randall, An ASTM standard for quantitative scratch
adhesion testing of thin, hard ceramic coatings, Int. J. Appl. Ceram.
Tech., 2 (2005), 422–428
21 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
22 A. E. Giannakopoulos, S. Suresh, Determination of elastoplastic pro-
perties by instrumented sharp indentation, Scripta Mater., 40 (1999),
1191–1198
23 A. E. Giannakopoulos, P. L. Larsson, R. Vestergaard, Analysis of
vickers indentation, Int. J. Solids Struct., 31 (1994), 2679–2708
24 Y. F. Gao, H. T. Xu, W. C. Oliver, G. M. Pharr, Effective elastic
modulus of film-on-substrate systems under normal and tangential
contact, J. Mech. Phys. Solids, 56 (2008), 402–416
25 S. Shim, H. Bei, E. P. George, G. M. Pharr, A different type of inden-
tation size effect, Scripta Materialia, 59 (2008), 1095–1098
26 Y. Huang, F. Zhang, K. C. Hwang, W. D. Nix, G. M. Pharr, G. Feng,
A model of size effects in nano-indentation, J. Mech. Phys. Solids,
54 (2006), 1668–1686
27 G. M. Pharr, Measurement of mechanical properties by ultra-low
load indentation, Mat. Sci. Eng. A- Struct., 253 (1998), 151–159
28 Y. Gogotsi, T. Miletich, M. Gardner, M. Rosenborg, Microinden-
tation device for in situ study of pressure-induced phase trans-
formations, Rev. Sci. Instrum., 70 (1999), 4612–4617
29 W. Zhu, P. J. M. Bartos, Application of depth-sensing microinden-
tation testing to study of interfacial transition zone in reinforced
concrete, Cement Concrete Res., 30 (2000), 1299–1304
30 O. Uzun, U. Kolemen, S. Celebi, N. Guclu, Modulus and hardness
evaluation of polycrystalline superconductors by dynamic micro-
indentation technique, J. Eur. Ceram. Soc., 25 (2005), 969–977
31 Z. Shan, S. K. Sitaraman, Elastic–plastic characterization of thin
films using nanoindentation technique, Thin Solid films, 437 (2003),
176–181
32 A. E. Giannakopoulos, S. Suresh, Scr. Mater., 40 (1999), 1191
33 R. P. Vinci, J. J. Vlassak, Annu. Rev. Mater. Sci., 26 (1996), 431
34 W. D. Nix, Mater. Sci. Eng. A, 237 (1997), 37
35 L. De Fazio, S. Syngellakis, R. J. K. Wood, F. M. Fugiule, G.
Sciume, Nanoindentation of CVD diamond: comparison of an FE
model with analytical and experimental data, Diam. Relat. Mater., 10
(2001), 765–769
36 J. Lia, W. Beres, Three-dimensional finite element modelling of the
scratch test for a TiN coated titanium alloy substrate, Wear, 260
(2006), 1232–1242
37 F. Sandiumenge, T. Puig, J. Rabier, J. Plain, X. Obradors, Optimi-
zation of flux pinning in bulk melt textured 1–2–3 superconductors;
Bringing dislocations under control, Adv. Mater., 12 (2000), 375
38 S. Block, G. J. Piermarini, R. G. Munro, W. Wong-Ng, The bulk
modulus and Young’s modulus of the superconductor Ba2Cu3YO7,
Adv. Ceram. Mater., 2 (1987), 601
39 H. T. Johansen, Flux-pinning-induced stress and magnetostriction in
bulk superconductors, Superconducting Science and Technology, 13
(2000), 121–137
40 H. M. Ledbetter, M. W. Austin, S. A. Kim, M. Lei, Elastic constants
and Debye temperature of polycrystalline yttrium barium copper
oxide (YBa2Cu3O7-x), J. Mater. Res., 2 (1987), 786
41 J. J. Roa, X. G. Capdevila, M. Martinez, F. Espiell, M. Segarra,
Nanohardness and Young’s modulus of YBCO samples textured by
the Bridgman technique, Nanotechnology, 18 (2007), 385
O. CULHA et al.: CHARACTERIZATION AND DETERMINATION OF MECHANICAL PROPERTIES OF YBCO ...
160 Materiali in tehnologije / Materials and technology 47 (2013) 2, 153–160
I. IVANOVI] et al.: DISTORTION OF THE SUBSTRUCTURE OF A 20-ft SHIPPING CONTAINER ...
DISTORTION OF THE SUBSTRUCTURE OF A 20-ft
SHIPPING CONTAINER EXPOSED TO ZINC HOT-DIP
GALVANIZING
SPREMINJANJE MER PODSTRUKTURE PRI 20 ft
TRANSPORTNEM KONTEJNERJU PRI VRO^EM POTOPNEM
CINKANJU
Ivana Ivanovi}1, Aleksandar Sedmak2, Rebeka Rudolf3, Leo Gusel3, Biljana Gruji}1
1Innovation Center, Faculty of Mechanical Engineering, University of Belgrade, Kraljice Marije 16, 11000 Belgrade, Republic of Serbia
2Faculty of Mechanical Engineering, University of Belgrade, Kraljice Marije 16, 11000 Belgrade, Republic of Serbia
3University of Maribor, Faculty of Mechanical Engineering, Smetanova ulica 17, 2000 Maribor, Slovenia
iivanovic@mas.bg.ac.rs
Prejem rokopisa – received: 2012-08-21; sprejem za objavo – accepted for publication: 2012-10-03
The main goal of this study was to build a model for a numerical simulation of hot-dip galvanizing of a 20-ft ISO shipping
container. For that purpose, the basic transient thermo-mechanical problem of a steel structure under the influence of the
temperature characteristic for a zinc hot-dip galvanizing bath was analyzed. Numerical calculations were performed for a simple
part and for the complex substructure of the container. Calculations were carried out on the Salome-Meca platform using a
Netgen mesh generator and a Code_Aster finite-element solver.
Keywords: shipping container, structural steel, zinc hot-dip galvanizing, transient heat transfer, distortion
Glavni namen te {tudije je bila izdelava modela za numeri~no simulacijo vro~ega potopnega cinkanja 20 ft ISO transportnega
kontejnerja. Zato smo analizirali osnovni termomehanski problem jeklene konstrukcije zaradi vpliva toplote med vro~im
potopnim cinkanjem. Numeri~ni izra~uni so bili izvr{eni za enostavni del in za kompleksno podstrukturo kontejnerja. Izra~uni
so bili izvr{eni na platformi Salome-Meca z uporabo generatorja mre`e Netgen in programske opreme Code_Aster na osnovi
metode kon~nih elementov.
Klju~ne besede: transportni kontejner, konstrukcijsko jeklo, vro~e potopno cinkanje, prenos toplote, deformacija konstrukcije
1 INTRODUCTION
This study is a part of a wider research launched in
order to develop an innovative 20-ft ISO shipping con-
tainer with a longer lifetime. Since shipping containers
have to face hard handling and hard weather conditions,
zinc hot-dip galvanizing, if possible, in the form of an
integrated welded structure, may represent the best way
to achieve this goal.
The dimensions of the 20-ft ISO shipping container
are 6.058 m × 2.348 m × 2.591 m with a net weight of
approximately 2 400 kg. The skeleton of the container is
composed of various steel profiles made of 235–420
graded structural steel. Profiles are welded to the cubical
corner fittings thus forming an orthogonal skeleton. The
wall and roof panels are composed of welded corrugated
steel sheets. The panels are welded continually to the rest
of the structure. The material thickness varies depending
upon the type of the element.
The entire skeleton will be taken into consideration
as a numerical model, despite the fact that the 20-ft
shipping container is larger than any existing galvanizing
bath, which means that in actual conditions it cannot be
galvanized as a whole.1 On the other hand, it seems that
it would be much more efficient to complete most of the
construction before galvanizing. Welding of galvanized
steel is possible, but it requires a special approach that
differs greatly from the standard steel welding
techniques.2
The temperature of the zinc bath is slightly above 450
°C. If other requirements for the preparation of an object
for galvanizing are fulfilled, this temperature should not
have a great influence on the mechanical properties of
the structural steel in the mentioned structural steel
grades.1,3 Problems were expected to arise due to an
extremely complex structure. For example, one of the
most important requirements for galvanizing is that the
thickness of the material of different parts that are going
to be galvanized together must be as uniform as possible.
Large differences in the thicknesses will result in
different heating and cooling of the material, and will
increase the risk of undesirable large distortions. In the
case of the shipping container, the thickness ranges from
28 mm in some regions of the corner fittings, 4.5 mm to
12 mm for the profiles in the skeleton and, finally, to 2
mm for the wall panel.
Problems will also arise in several areas of the
structure where the hollow sections are completely
sealed by the other parts of the structure. These sections
must be replaced with open sections or vent/drain holes
must be provided. An even greater problem is the
Materiali in tehnologije / Materials and technology 47 (2013) 2, 161–167 161
UDK 519.61/.64:669.58:669.14.018.291 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)161(2013)
numerous corners of the element connections that also
need to be supplied with vent/drain holes.
When it comes to numerical simulation, there is very
little information on previous studies dealing with zinc
hot-dip galvanizing, especially those dealing with the
fluid dynamics of zinc at the temperature of the galva-
nizing bath, and the fluid dynamics of the galvanizing
bath. Some insight into the field could be gained from a
thermal analysis carried out, during the solidification, on
the moving surface in a finite bath given in 4. Concerning
a transient three-dimensional thermo-mechanical nume-
rical analysis, there is a great number of papers dealing
with this subject within the welding researches.5–8 The
basic thermo-mechanical analysis, at some level similar
to this one, is performed and described in 9,10.
This is a highly idealized analysis of a real-time
problem, and numerous real-time conditions are ignored
deliberately.
2 GENERAL-MODELING DESCRIPTION
2.1 Geometry of the model
The skeleton used as a model for numerical
simulations of the zinc hot-dip galvanizing of a 20-ft ISO
shipping container is presented in Figure 1. The skeleton
model is composed of two front-corner posts, two outer
rear-corner posts, two bottom side rails, two top side
rails, a door sill, the lower part of the door header, a front
sill, a front-top end rail, three types of floor cross-
members, two forklift-pocket top plates and, finally, four
top- and four bottom-corner fittings.
It was already mentioned that some elements in the
construction should be omitted or replaced in order to
avoid closed hollow sections. Such sections are undesi-
rable in galvanizing, but also need special treatment
during a numerical simulation. The initial examples, the
most sensitive to changes, are the corner posts. The
front-corner post is planned to be a 6 mm thick,
S235JR-steel closed profile as illustrated in Figure 2
(left). When welded to the corner fittings, this profile
will form a closed hollow section. Therefore, in simu-
lations, the closed profile is replaced by a standard-
shaped open profile of the same thickness as illustrated
in Figure 2 (right). To simplify an already complex
geometry, the edge filleting is omitted here, as well as
for all the other elements.
The rear-corner-post profile consists of an inner and
outer part. A standard 6 mm thick outer profile, planned
to be made of the S420NL steel, is kept in the numerical
model. The inner part could not be omitted from the real
skeleton, but it is assumed that it can be added after-
wards. The notches for door-hinge-pin lugs, which are
usually made on the outer part of the rear-corner post,
are omitted to simplify the geometry.
Bottom side rails (4.5 mm S420NL steel) are
introduced without the two forklift-pocket openings.
Only the standard U-shaped lower part of the door
header (4 mm S355NL steel) is included. For the two top
side rails and for the front-top end rail (4 mm S355NL
steel), the standard square tubes are replaced with the
channel-shaped steel sections.
The elements of the skeleton are connected to the
construction with corner fittings. The dimensions of the
corner fittings are much smaller than the dimensions of
the other elements, 178 mm × 162 mm × 118 mm, but
the thickness of the walls is much higher, and it ranges
from 10 mm to 28 mm. They have a specific geometry
since they are designed for lifting, stacking and securing
the container. The real form of the top-corner fitting is
illustrated in Figure 3 (left). A simplified form, attached
I. IVANOVI] et al.: DISTORTION OF THE SUBSTRUCTURE OF A 20-ft SHIPPING CONTAINER ...
162 Materiali in tehnologije / Materials and technology 47 (2013) 2, 161–167
Figure 2: Cross sections of the front-corner-post profiles
Slika 2: Prereza profilov kotnih drogov v sprednjem delu
Figure 1: Skeleton of the 20-ft ISO shipping container
Slika 1: Ogrodje 20 ft ISO ladijskega kontejnerja
Figure 3: Top-corner-fitting geometry
Slika 3: Geometrija gornjega vogala
to the cables in a way that will be used for numerical
simulations, is presented in Figure 3 (right).
2.2 Numerical procedures
All pre-processing and post-processing work is
performed using an open-source Salome-Meca platform.
From the mesh module in Salome-Meca, an open-source
Netgen mesh generator is applied to generate an auto-
matic three-dimensional tetrahedral mesh. An example
of the mesh generated on the top-corner fitting, with the
maximum edge size of 9 mm, resulting in 14 358 tetra-
hedral elements, is illustrated in Figure 4.
The molten zinc bath is kept at a constant tempe-
rature of 450 °C during all numerical simulations. The
top-corner fitting and the whole skeleton are chosen
separately for the simulations. The bottom side of each
model is placed in an xy plane. The plane is placed so as
to be parallel to the surface of the zinc bath. The models
are suspended with four steel cables (Figure 3 (left) and
Figure 4). In this manner, for the purpose of the
numerical model of the skeleton, the mechanism for
vertical lifting of the container from the top corner
fittings is simplified and replaced only by cables. Cable
thickness is selected so that its elongation is reduced to a
minimum.
A transient thermomechanical analysis is performed
with the open-source finite-element solver Code_Aster.
The thermomechanical properties of the steel are kept
constant during the simulations. In the linear thermal
analysis, convection boundary conditions are imposed on
all the surfaces. An immersion is simulated at a constant
velocity of 0.004 m/s and with two different constant
heat-transfer coefficients, one at the immersed surfaces
and the other at the surfaces outside the zinc bath. The
initial temperature of the object and the temperature of
the surroundings are the same, 25 °C. The thermal loads
and the weight of the model are introduced in a
non-linear static analysis. The cables are fixed at one end
and connected to the selected points of the object (the
corner fitting or skeleton) at the other end.
3 RESULTS AND DISCUSSIONS
Simulations were performed on a desktop PC with an
Intel Core i5-2300 CPU running at 2.8 GHz and with 4
GB RAM. Because of the limited computer resources, a
simple top-corner fitting model was chosen for general
model verifications and validations.
3.1 Top-corner fitting
Corner fittings are the smallest parts of a container
construction with very different wall thicknesses. Points
8, 9, 10 and 11, which were used for the presentations of
the results obtained for different mesh densities, are
illustrated in Figure 5. The thickness of the walls around
the bottom point 9 is 20 mm on the sides and 10 mm at
the bottom. The thickness around the bottom point 8 is
20 mm and 10 mm on the sides. The thickness of the top
wall, where points 11 and 10 are situated, is approxi-
mately 28 mm. In the case of a velocity of 0.004 m/s, the
corner fitting is fully immersed after approximately 30 s.
The temperature results for the selected points are
illustrated in Figure 6a. The results are presented from
the beginning of the simulation until 630 s. It can be seen
that the temperatures of four points differ throughout the
entire simulation. The maximum value of the tempera-
ture difference between the two bottom points is
approximately 33 °C, and is reached at 240 s. The
temperature of the zinc bath was not reached at the end
of the simulation, despite the fact that the corner fitting
was immersed fully for 10 min. The maximum
temperature was reached at point 8 and has the value of
408 °C.
I. IVANOVI] et al.: DISTORTION OF THE SUBSTRUCTURE OF A 20-ft SHIPPING CONTAINER ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 161–167 163
Figure 5: Positions of points 8, 9, 10 and 11
Slika 5: Pozicije to~k 8, 9, 10 in 11
Figure 4: Example of a Netgen three-dimensional tetrahedral mesh
Slika 4: Primer tridimenzionalne tetraedri~ne mre`e Netgen
The results obtained using three different mesh
densities are presented in Figure 6b. The values for the
coarse mesh of 2 655 elements (the maximum edge size
is 20 mm) are significantly different from the other two
sets of values (for the mesh of 14 358 elements the
maximum edge size is 9 mm, and for the mesh of 59 739
elements the maximum edge size is 5 mm). The differen-
ce was around 3.5 °C at the beginning of the simulation,
and around 1 °C at the end of the simulation.
At the beginning of the simulation the temperature at
some points of the mesh fell below the assigned initial
temperature of the model, which is equal to the tempe-
rature of the surroundings. This error decreased with an
increase in the mesh density.
The previous discussions are confirmed with Figure
7, where the temperature distribution, at 30 s (a) and at
630 s (b), is presented for the mesh of 14 358 elements.
At 30 s, when the phase of immersion is just finished,
some parts of the corner fitting are still at the initial
temperature of 25 °C. This is an obvious consequence of
the already mentioned numerical error caused by the
finesse of discretization.
Compared to the temperature of the other previously
selected points (Figure 5), the highest temperature is in
the vicinity of point 8, and is around 80 °C. The tempe-
I. IVANOVI] et al.: DISTORTION OF THE SUBSTRUCTURE OF A 20-ft SHIPPING CONTAINER ...
164 Materiali in tehnologije / Materials and technology 47 (2013) 2, 161–167
Figure 6: a) Temperature at points 8, 9, 10 and 11 from the beginning
of the simulation until 630 s for the mesh of 14 358 elements and b)
the temperature at point 9 for three different mesh densities from 10 s
to 40 s
Slika 6: a) Temperature v to~kah 8, 9, 10 in 11 od za~etka simulacije
do 630. sekunde za mre`o s 14 358 elementi in b) temperature v to~ki
9 za tri gostote mre`e od 10. sekunde do 40. sekunde
Figure 8: Total displacements at the top-corner fitting at 30 s (scale
200) and at 630 s (scale 40) for the mesh of 14 358 elements
Slika 8: Skupni raztezki v vrhnjem kotnem spoju pri 30 s (skala 200)
in pri 630 s (skala 40) za mre`o s 14 358 elementi
Figure 7: Temperature distribution at the top-corner fitting at: a) 30 s
and b) at 630 s for the mesh of 14 358 elements
Slika 7: Razporeditev temperature v vrhnjem kotnem spoju pri: a) 30
s in b) pri 630 s za mre`o s 14 358 elementi
a)
b)
rature difference between these points is even more
evident at the end of the simulation as illustrated in
Figure 7b. At 630 s, the maximum temperature is
approximately 410 °C, and it is 40 °C lower than the
temperature of the zinc bath. The corresponding total
displacements are illustrated in Figure 8.
3.2 Skeleton
In the case of the skeleton, two mesh densities were
used for numerical simulations. The maximum edge size
of the finer mesh was equal to the critical edge size of
the corner fitting (20 mm). The immersion phase takes
approximately 648 s. Because of the limited computer
resources, a coarser mesh of 241 623 elements (the
maximum edge size of 60 mm) was used to reach the end
of the immersion phase. The total displacements and
temperature distributions at (50, 350, and 650) s are
illustrated in Figure 9.
As expected, for this mesh density the minimum
temperature at any time was lower than the initial
temperature of 25 °C. The temperature of the zinc bath
was almost reached at the lower parts of the skeleton by
the end of the simulation. The maximum temperature at
650 s, illustrated in Figure 9 (bottom right), was appro-
ximately 449 °C.
It is not hard to notice large differences in the
temperature at different positions of the skeleton.
Especially at the beginning of the immersion, these
temperature differences can produce large distortions.
For example, from the temperature distribution at 50 s
illustrated in Figure 9 (top right), it can be seen that the
temperature is much higher at the bottom than at the top
of the bottom side rails. The temperature of the bottom-
corner fittings is even lower. The largest distortions are
found in the central area of the floor construction, at the
central cross members, and at the two forklift-pocket top
plates (Figure 9 (top left)). There is an obvious dis-
tortion of the bottom side rails. This distortion affects the
very end of the front corner posts, near the corner
fittings. The influence of the distortion of the floor
construction on the outer part of the rear corner posts is
even more pronounced, but it is also better distributed
along the posts than in the case of the front corner posts.
At 350 s, the top side rails started to deform as a result of
the distortion of the lower half of the skeleton (Figure 9
(center right)).
The maximum value of the total displacements
reached in the first 50 s is around 12.5 mm, in the next
300 s it is increased by only 5 mm, and in the next 300 s
it is increased by an even smaller value of only 3 mm.
The influence of the mesh density, for the mesh of
241 623 elements with the maximum edge size of 60
mm, and the mesh of 490 420 elements with the maxi-
mum edge size of 20 mm, from 80 s to 120 s, at the point
621 with the coordinates (0, 0, 0), is illustrated in Figure
10. A simulation for the finer mesh was executed for the
period of up to 200 s. The temperature difference was
between 3.5 °C at the beginning of the simulation and
2.5 °C at 200 s. The values of the total displacements,
from 80 s to 120 s, were higher for the coarser mesh.
The temperature at different points on the bottom-
corner fitting and in its vicinity is presented in Figure
11. During the simulation, the temperature was the
highest at point 4 067, which is at the bottom of the front
sill, and at point 31 870, which is at the bottom of the
bottom side rail. The thickness of the material of these
profiles is 4.5 mm.
I. IVANOVI] et al.: DISTORTION OF THE SUBSTRUCTURE OF A 20-ft SHIPPING CONTAINER ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 161–167 165
Figure 10: Temperature and total displacements at point (0, 0, 0) for
the two different mesh densities, a mesh of 241 623 elements with the
maximum edge size of 60 mm, and a mesh of 490 420 elements with
the maximum edge size of 20 mm, from 80 s to 120 s
Slika 10: Temperatura in celotni raztezek v to~ki (0, 0, 0) za dve
razli~ni gostoti mre`e: mre`e z 241 623 elementi, maksimalna velikost
roba 60 mm, in mre`o s 490 420 elementi, maksimalna velikost roba
20 mm, od 80 s do 120 s
Figure 9: Total displacements (the scale factor of 20) and temperature
distributions at (50, 350, and 650) s for the skeleton in the case of a
mesh of 241 623 elements (the maximum edge size of 60 mm)
Slika 9: Celotni raztezek (faktor skale 20) in razporeditev temperature
pri (50, 350 in 650) s za ogrodje v primeru mre`e z 241 623 elementi
(maksimalna velikost roba 60 mm)
Points 621 and 610 are on the bottom-corner fitting
and their positions are equivalent to those of points 11
and 10, illustrated in Figure 5. As expected, the tempe-
rature of point 610 was higher than the temperature of
point 621, as presented in the case of the corner fittings.
Both temperatures were lower than the temperatures of
the points on the bottom profiles.
The last presented point is at the front-corner post,
just above the bottom-corner fitting. The thickness of the
material of the front-corner post is 6 mm. This point
starts to heat up 100 s after the above-presented points,
but in the following 100 s it exceeds the temperature of
the two points on the bottom-corner fitting.
The distributions for the finer mesh at 40 s are pre-
sented in Figure 12. While the minimum temperature for
the coarse mesh at 50 s, presented in Figure 9 (top
right), was approximately 21 °C, for the finer mesh this
value was around 22 °C at 40 s (Figure 12b). Also, the
maximum value for the total displacements was almost 1
mm higher at 40 s for the finer mesh than at 50 s for the
coarser mesh (Figure 9 (top left)).
The maximal stresses at 40 s, illustrated in Figure
13, were located at the bottom of the bottom side rails
near the front bottom-corner fittings. The bottom rail is
planned to be made of 4.5 mm thick S420NL steel, and
its maximum stress value is 381 MPa.
4 CONCLUSIONS
Three-dimensional linear thermal and nonlinear static
analyses, of the top-corner fitting as a part of the con-
tainer structure and of the entire skeleton as the container
substructure, were performed to initiate the analysis of
the distortion of a 20-ft ISO shipping container exposed
to the temperature of a zinc hot-dip galvanizing bath.
Numerical simulations were carried out using open-
source CAE software: Salome-Meca, Netgen, and
Code_Aster. The temperature was the only transient
boundary condition. The temperature of the galvanizing
bath, the velocity of immersion, and the thermomecha-
nical properties of steel were kept constant during all the
simulations. Tests were carried out within the limits of
the very poor computer resources. However, satisfactory
initial results were obtained and sufficient information
was gathered to make a general image of the container-
substructure behavior under the influence of the zinc
hot-dip galvanizing-bath temperature.
I. IVANOVI] et al.: DISTORTION OF THE SUBSTRUCTURE OF A 20-ft SHIPPING CONTAINER ...
166 Materiali in tehnologije / Materials and technology 47 (2013) 2, 161–167
Figure 13: Maximum stress at the bottom side rail near the front-
bottom-corner fitting at t = 40 s
Slika 13: Maksimalna napetost v spodnjem stranskem profilu blizu
spodnjega sprednjega vogalnega spoja pri t = 40 s
Figure 11: Temperature at different points of the skeleton, on and
around the bottom-corner fitting, from the beginning of the simulation
to 200 s, for the mesh of 490 420 elements
Slika 11: Temperature v razli~nih to~kah ogrodja na spodnjem
vogalnem spoju in okrog njega od za~etka simulacije do 200 s v
primeru mre`e s 490 420 elementi
Figure 12: a) Total displacements (scale 20) and b) the temperature
distribution at 40 s for the mesh of 490 420 elements
Slika 12: a) Skupni raztezek (skala 20) in b) razporeditev temperature
pri 40 s za mre`o s 490 420 elementi
a)
b)
Acknowledgements
The authors would like to acknowledge the financial
support given under the EUREKA E! 5009 Galvacont
project.
5 REFERENCES
1 IGAG, INGAL Specifiers Manuel, Available: www.ingal.com.au/
igsm.htm
2 G. Livelli, T. Langill, Guidelines for Welding Galvanized Steel, PCI
Journal, (1998), 40–48
3 L. Mraz, J. Lesay, Problems with reliability and safety of hot dip
galvanized steel structures, Soldagem and inspecao, 14 (2009) 2,
184–190
4 H. Zhang, K. M. Moallemi, S. Kumar, Thermal Analysis of the Hot
Dip-Coating Process, Journal of Heat Transfer, 115 (1993), 453–460
5 M. Berkovi}, S. Maksimovi}, A. Sedmak, Analysis of Welded Joints
by Applying the Finite Element Method, Structural Integrity and
Life, 4 (2004) 2, 75–83
6 D. Velji}, M. Perovi}, A. Sedmak, M. Rakin, N. Baji}, B. Medjo, H.
Dascau, Numerical Simulation of the Plunge Stage in Friction Stir
Welding, Structural Integrity and Life, 11 (2011) 2, 131–134
7 D. Velji}, M. Perovi}, A. Sedmak, M. Rakin, M. Trifunovi}, N.
Baji}, D. Baji}, A Coupled Thermo-Mechanical Model of Friction
Stir Welding, Thermal Science, 16 (2012) 2, 527–534
8 M. Perovi}, D. Velji}, M. Rakin, N. Radovi}, A. Sedmak, N. Baji},
Friction-Stir Welding of High-Strength Aluminium Alloys and a
Numerical Simulation of the Plunge Stage, Mater. Tehnol., 46 (2012)
3, 105–111
9 W. J. Rudd, S. W. Wen, P. Langenberg, B. Donnay, A. Voelling, T.
Pinger, M. Feldmann, J. Carpio, J. A. Casado, J. A. Alvarez, F.
Gutierrez-Solana, Failure mechanisms during galvanising, Office for
Official Publications of the European Communities, Luxembourg,
2008
10 M. Feldmann, T. Pinger, D. Schäfer, R. Pope, G. Sedlacek, Hot-dip-
zinc-coating of prefabricated structural steel components, Office for
Official Publications of the European Communities, Luxembourg,
2010
I. IVANOVI] et al.: DISTORTION OF THE SUBSTRUCTURE OF A 20-ft SHIPPING CONTAINER ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 161–167 167

A. TAÞKESEN, K. KÜTÜKDE: OPTIMIZATION OF THE DRILLING PARAMETERS FOR THE CUTTING FORCES ...
OPTIMIZATION OF THE DRILLING PARAMETERS FOR
THE CUTTING FORCES IN B4C-REINFORCED
Al-7XXX-SERIES ALLOYS BASED ON THE TAGUCHI
METHOD
OPTIMIRANJE PARAMETROV VRTANJA ZA SILE VRTANJA PRI
ZLITINAH Al-7XXX, OJA^ANIH Z B4C S TAGUCHIJEVO
METODO
Ahmet Taþkesen1, Kenan Kütükde2
1Gazi University, Department of Manufacturing Engineering, 06500 Teknikokullar, Ankara, Turkey
2Gazi University, Institute of Science and Technology, 06500 Teknikokullar, Ankara, Turkey
taskesen@gazi.edu.tr
Prejem rokopisa – received: 2012-08-23; sprejem za objavo – accepted for publication: 2012-09-28
In this study, drilling tests of aluminum-based composites produced with the powder-metallurgy (PM) technique and reinforced
with boron-carbide (B4C) particles were carried out with three different types of drills under dry cutting conditions. In order to
determine the mechanical properties of the produced composites, hardness and tensile tests were performed. Moreover, the
effects of the machining parameters such as cutting speed, feed rate, particle fraction and cutting-tool material, and of their
interactions on the thrust force and cutting torque were determined with the Taguchi experimental design. Drilling parameters
were optimized in terms of cutting forces (thrust force and torque). Furthermore, an analysis of variance (ANOVA) was
conducted to obtain the degree of the effect of the parameters. The most influential control factors for the cutting forces were
found to be the particle fraction and feed rate. According to the experimental results, the thrust force and cutting torque
increased significantly as the feed rate or the particle content increased. On the other hand, the influence of the drill-bit material
and the interactions of the factors for the cutting forces were quite low.
Keywords: B4C, powder metallurgy, drilling, cutting force, torque, Taguchi method
V tej {tudiji so bili narejeni preizkusi vrtanja kompozita na osnovi aluminija, oja~anega z borovim karbidom (B4C) in izdelanega
po postopku pra{ne metalurgije (PM), s tremi razli~nimi svedri in pri suhem vrtanju. Za dolo~itev mehanskih lastnosti
izdelanega kompozita je bila izmerjena trdota in opravljeni so bili natezni preizkusi. Poleg tega so bili dolo~eni s Taguchijevo
eksperimentalno tehniko parametri obdelave, kot so hitrost rezanja, hitrost podajanja, dele` delcev, material za orodje za rezanje
in njihov vpliv na potisno silo ter navor pri rezanju. Parametri rezanja so bili optimirani glede na sile rezanja (potisna sila in
navor). Poleg tega je bila narejena analiza variance (ANOVA), da bi dobili stopnjo vpliva parametrov. Ugotovljeno je bilo, da sta
najbolj vplivna kontrolna faktorja na sile rezanja dele` delcev in hitrost podajanja. Skladno z rezultati preizkusov potisna sila in
navor mo~no narasteta, ~e se pove~a hitrost podajanja ali pove~a vsebnost delcev. Po drugi strani so razmeroma majhni vplivi
materiala svedra in medsebojni vpliv faktorjev na sile rezanja.
Klju~ne besede: B4C, pra{na metalurgija, vrtanje, sila rezanja, navor, Taguchijeva metoda
1 INTRODUCTION
The composition of many composite materials used
in engineering applications consists of additives pro-
viding a better hardness and resistance and of the matrix
material that holds these substances together as well as
allowing ductility and toughness.1 Due to their high
specific strength, superior wear resistance, low thermal
expansion and lightweight, metal-matrix composites
(MMCs), widely used, especially in aerospace and
automotive industry, have attracted the attention of the
researchers.2–4 However, in spite of these advantages, the
machinability of these composites is difficult.5–12
A drilling process is one of the last production stages
that have to be done before the assembly step. The past
studies relating to the drilling of MMCs have revealed
that Al2O3 and SiC are mostly used as a reinforcement
material in an aluminum composite material.3,5–8,12–14
However, there are no adequate studies on the drilling of
the B4C-reinforced aluminum composites. An inclusion
of the B4C particles as a reinforcement material has the
advantage of having a higher hardness (4200 HV) than
the other ceramics such as SiC (3500 HV) and Al2O3
(2300 HV).2,15
In the previous studies regarding the drilling
machinability of MMCs, it is stated that an increase in
the cutting speed does not significantly affect the thrust
force, and that the most important factor increasing the
thrust force is the feed rate.3,10,16 Moreover, the particle
content in the composite material as well as the drilling
tool are of importance for the drilling of aluminum-
matrix composites; and the lowest drilling forces are
obtained with polycrystalline diamond (PCD) drills.3 In
addition, the coated carbide tools produce more thrust
forces than the uncoated carbide drills.16 Heat-treatment
conditions also have a significant effect on the cutting
forces and the highest tool forces were observed (nearly
twice) when drilling aged the composites.6,8 On the other
Materiali in tehnologije / Materials and technology 47 (2013) 2, 169–176 169
UDK 621.762:621.95:519.233.4 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)169(2013)
hand, while an addition of graphite to a composite
material positively affects both the cutting forces and the
machinability, it adversely affects the strength of the
composite material.2,5,12 With respect to the cutting-tool
material, lower thrust forces are obtained when through-
tool cooling is performed.7 However, the cutting torques
produced with conventional cooling (the cooling method,
in which the cooling fluid is sprayed from the outside to
the cutting zone) are lower than those produced with the
through-tool cooling and dry drilling. Thrust forces also
increase depending on the drilled hole number.7 From the
point of view of the cutting forces, the results of drilling
fiber-reinforced composites are similar to the results of
drilling MMCs.9,11
The Taguchi design method is a useful tool for
determining the effect of machining parameters and their
significance levels. A plan of experiments can be con-
ducted with the Taguchi method with the purpose of
analyzing the data and obtaining the information about
the property of a certain process. This method uses
orthogonal arrays for defining the experiment
plan.4,13,14,16 Its important advantage is the fact that it
saves experimental effort and time, reducing the cost.
Furthermore, the results other than the conducted experi-
ments can be predicted with a great accuracy by using
this method. In recent times, a variety of applications of
the Taguchi method have been performed in many areas.
The aim of this study was to introduce the Taguchi
method in determining the optimum drilling conditions
for the thrust force and drilling torque when drilling an
Al7XXX alloy reinforced with three different mass
fractions of B4C particles. For this purpose, the effect of
the control factors such as spindle speed, feed rate,
particle fraction and cutting tool on the cutting forces
were investigated. Significance levels of individual
factors were determined with ANOVA. The values pre-
dicted with the Taguchi method were compared with the
experimental results.
2 EXPERIMENTAL PROCESS
2.1 Production of Composite Materials
In this work, 7xxx-series aluminum alloy (including
mass fractions: 5 % zinc, 3.5 % copper and 2.5 % of
magnesium) was used as the matrix element. B4C
ceramic powders under 325 meshes were used as the
reinforcement element. To investigate the effects of
different reinforcement fractions on the machinability,
three different weight fractions of B4C particles were
selected as 10 %, 15 % and 25 %. The mixture was cold
pressed in the mold under the pressure of 25 MPa in an
electrical furnace. Then the internal temperature of the
furnace was fixed at 540 °C and the composite materials
were produced by applying the liquid-phase sintering
method for half an hour. Later the produced samples
were subjected to the hardness, tensile and drilling tests.
For the hardness test, three hardness measurements were
performed on each sample by using OKO SEIKI
hardness-measurement equipment and the mean of the
hardness values was used. Tensile tests were also carried
out by placing each specimen into a 60-ton Tinius Olsen
tensile-test device. Tensile-test specimens were prepared
according to the EN 10002-1 standard by turning the
sintered blocks as shown in Figure 1.
2.2 Test Setup and the Drilling Process
For the drilling tests of the produced MMCs, a com-
puter numerically controlled (CNC) vertical machining
center (VMC-550 Johnford Fanuc Series O-M) having
the capacity of 15 kW and 3 500 r/min was used. The
machining conditions and geometrical properties of the
drills are given in Table 1. The cutting forces were
measured for all the drilling experiments with three
previously unused, different, 8-mm drills. Each test was
repeated twice and the mean values were used. A total of
100 holes were drilled in addition to 27 Taguchi
experiments for confirmation purpose. The length of the
drilled composites was 12 mm. A KISTLER 9272 dyna-
mometer was used to measure the thrust force and torque
during the drilling process. Figure 2 shows the sche-
matic image of the drilling setup17. A picture of the
produced composite, attached to a specially developed
and manufactured fixture, after being drilled with the
CNC vertical machining center, is depicted in Figure 3.
After measuring the thrust forces and drilling torques,
the results were recorded into a computer environment
using the KISTLER DynoWare software. The average
value of the measured cutting forces was taken into
account so that the conical section of the tool tip was
completely inside the workpiece. A sample output of the
A. TAÞKESEN, K. KÜTÜKDE: OPTIMIZATION OF THE DRILLING PARAMETERS FOR THE CUTTING FORCES ...
170 Materiali in tehnologije / Materials and technology 47 (2013) 2, 169–176
Figure 1: Tensile-test specimens: a) prepared test specimens, b) tech-
nical drawing of a test specimen
Slika 1: Preizku{anci za natezni preizkus: a) pripravljeni preizku{anci,
b) tehni~na risba preizku{anca
dynamometer showing the variation of the cutting force
and torque is given in Figure 4.
3 RESULTS AND DISCUSSION
3.1 Microstructure and Mechanical Properties
The microstructure of the produced composites is
shown in Figure 5. A homogeneous distribution of the
ceramic particles over the composite alloy can be seen
from this figure. According to the hardness test results,
the average hardnesses of the specimens with mass
fractions 10 % B4C, 15 % B4C and 25 % B4C were meas-
ured as 61 HRB, 79 HRB and 87 HRB, respectively.
These hardnesses were significantly higher than that for
the Al7075 alloy (43 HRB) but close to the Al7075–T6
alloy (87 HRB).18 The hardness of the composites in-
creased as the particle fraction increased due to the hard
nature of the ceramic particles.
The results of the strength and elongation (%) for
each sample are given in Table 2. The highest yield and
A. TAÞKESEN, K. KÜTÜKDE: OPTIMIZATION OF THE DRILLING PARAMETERS FOR THE CUTTING FORCES ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 169–176 171
Figure 4: Typical cutting forces observed when spindle speed = 1500
r/min, feed = 0.3 mm/r, the work piece contains the mass fraction 10
% B4C and the drill used is made of HSS
Slika 4: Zna~ilna sila rezanja pri hitrosti vrtenja vretena = 1500 r/min,
podajanje 0,3 mm/r, obdelovanec je vseboval masni dele` 10 % B4C,
sveder je bil iz HSS-jekla
Figure 3: Drilling setup
Slika 3: Preizkus vrtanja
Table 1: Machining conditions
Tabela 1: Pogoji obdelave
Machine tool Johnford VMC-550 Fanuc Serial O-M CNC controlled vertical machining center
Drills
HSS: 
 8 mm, 135° tool tip angle, spiral, 30° helical angle
Uncoated carbide: 
 8 mm, 140° tool tip angle, spiral, 30° helical angle
TiAlN-coated
carbide: 
 8 mm, 140° tool tip angle, spiral, 30° helical angle
Workpiece materials Mass fractions: 10 % B4C/Al, 15 % B4C/Al and 25 % B4C/Al composite
Cutting parameters Spindle speeds (n): 1000 r/min, 1500 r/min, 2000 r/min, 2500 r/minFeed rates (f) : 0.1 mm/r, 0.2 mm/r, 0.3 mm/r
Figure 2: Schematic presentation of the measuring setup
Slika 2: Shematski prikaz sestava za merjenje
Figure 5: Microstructure of the composite having the mass fraction
15 % B4C
Slika 5: Mikrostruktura kompozita z masnim dele`em 15 % B4C
tensile strength values were obtained when the B4C
particle fraction was 15 %. Generally, an increase in the
particle fraction increases the strength of the composite
material but, at the same time, reduces ductility due to an
increased dislocation density.19 In this study, it was
observed that the ceramic reinforcements added to the
aluminum matrix reduced the ductility of the composite
material and made it more brittle (Table 2). The fact that
the strength of the composite having 25 % particle
fraction was lower than the strength of the composites
having mass fractions 10 % B4C and 15 % B4C can be
attributed to the increase in the interfacial decompo-
sitions between the particles and the matrix.20
Table 2: Strength results of the B4C-reinforced MMC
Tabela 2: Trdnost MMC, oja~anega z B4C
B4C particle
mass fraction
w/%
Yield
strength
MPa
Tensile
strength
MPa
Elongation
%
10 491 527 22.2
15 532 599 6.9
25 328 408 4.8
3.2 Cutting Forces and Torques
According to the experimental results, the effects of
the cutting parameters such as particle fraction, cutting
speed, feed rate and cutting-tool material on the thrust
force and torque were given in Figures 4 and 5. The
cutting forces increased with the particle weight fraction,
and the rate of this increment for HSS tools was higher
than that for the carbide tools with greater particle
fractions. An increase in the weight fraction of the B4C
particles within the aluminum matrix increased the
hardness of the composite causing a rapid tool wear due
to a more intense contact with the cutting edge. There-
fore, increasing both the weight fraction and the area of
hard particles being in contact with the cutting tool
resulted in an increase in the friction and flow strength of
the cutting tool-chip as well as the cutting tool-work-
piece interface. On the other hand, it could be observed
from Figures 6 and 7 that the thrust force and cutting
torque increased with the feed rate, but they decreased
with the cutting speed. Previous researchers stated that
the most important factor affecting the cutting forces was
the feed rate3,8–12,16 and this was confirmed by our study.
Since the chip volume removed per revolution of the
cutting tool increased with an increase in the feed rate,
the thrust force and cutting torque increased as well.21
When experimental results were analyzed in terms of
the cutting-tool material, higher thrust forces, ranging
from 150 N to 250 N, were produced with HSS tools
than with carbide tools while drilling the 25 % B4C
reinforced MMCs. This situation could be attributed to
the hardness of the cutting tool and to the wear mecha-
A. TAÞKESEN, K. KÜTÜKDE: OPTIMIZATION OF THE DRILLING PARAMETERS FOR THE CUTTING FORCES ...
172 Materiali in tehnologije / Materials and technology 47 (2013) 2, 169–176
Figure 7: Effect of machining parameters on the drilling torque: a) particle mass fraction, w/%, b) feed rate, mm/r, c) spindle speed, r/min
Slika 7: Vpliv parametrov obdelave na navor pri vrtanju: a) vsebnost delcev v masnih dele`ih, w/%, b) hitrost podajanja, mm/r, c) hitrost vrtenja
vretena, r/min
Figure 6: Effect of machining parameters on the thrust force: a) particle mass fraction, w/%, b) feed rate, mm/r, c) spindle speed, r/min
Slika 6: Vpliv parametrov obdelave na potisno silo: a) vsebnost delcev v masnih dele`ih, w/%, b) hitrost podajanja, mm/r, c) hitrost vrtenja
vretena, r/min
nisms of the drilling tool. Since HSS tools had a lower
hardness than carbide tools, higher thrust forces were
produced with HSS drills, especially with a higher
particle fraction due to the tool flank wear. When the
particle fraction was less than 25 %, the difference
between the thrust forces produced by HSS and carbide
tools was relatively lower (20 N–40 N). However, the
test results indicated that HSS drills generally produced
less cutting torques than carbide drills as shown in
Figure 7. This condition might be attributed to the point
angle of the drill because the thrust force and cutting
torque increased with an increased point angle.22
Therefore, it was concluded that HSS drills produced
less cutting torques than carbide drills due to having a
point angle smaller by 5° than carbide drills. Consequen-
tially, although lower cutting forces were produced by
PCD diamond tools according to the existing literature3,
carbide tools could be preferred for machining MMCs
taking into account the production-cost balance.
3.3 Optimization with the Taguchi Method
In this section, optimization of drilling parameters
was carried out in terms of drilling forces with the
Taguchi analysis. The importance order of the effects of
each control factor on drilling forces was identified. For
this purpose, the factors selected in the Taguchi
experimental design and the levels of these factors are
shown in Table 3. A four-factor, 27-line and three-level
L27 (313) orthogonal array was chosen since it has the
ability to control the interactions among the factors.23–25
In the Taguchi method, there are three categories
such as "the smallest is better", "he biggest is better" and
"the nominal is better" for the calculation of the
signal/noise (S/N) ratio. In this study, since "the lowest"
thrust-force and cutting-torque values were desired for
the optimization, "the smallest is better" calculation
method was chosen. In the ith experiment, the S/N ratio i
can be calculated using the following equation14,26,27:
 i i
i
n
n
Y= −
⎛
⎝
⎜ ⎞
⎠
⎟
=
∑10 110 2
1
log (1)
where n is the number of replications and Yi is the
measured characteristic value (i.e, the thrust force or
cutting torque). The calculated S/N ratios () of the
thrust forces and cutting torques are given in Table 4.
Table 4: Experimental design with the L27 orthogonal array and the
S/N ratios
Tabela 4: Oblikovanje preizkusov z ortogonalno razporeditvijo L27 in
razmerje S/N
Test
No A B C D
Thrust
force
(N)
S/N
ratio for
thrust
force
Torque
(N cm)
S/N
ratio for
torque
1 1 1 1 1 843.2 –58.52 293 –49.34
2 1 1 2 2 834.3 –58.43 300.9 –49.57
3 1 1 3 3 804.3 –58.11 249.3 –47.93
4 1 2 1 2 966.8 –59.71 478.4 –53.60
5 1 2 2 3 924.1 –59.31 344 –50.73
6 1 2 3 1 880.7 –58.90 294 –49.37
7 1 3 1 3 1113.3 –60.93 434.7 –52.76
8 1 3 2 1 1048 –60.41 446.6 –53.00
9 1 3 3 2 1096 –60.80 291.3 –49.29
10 2 1 1 2 1168 –61.35 500.3 –53.99
11 2 1 2 3 1062 –60.53 398.5 –52.01
12 2 1 3 1 973.3 –59.77 309.1 –49.80
13 2 2 1 3 1244 –61.90 496.7 –53.92
14 2 2 2 1 1132 –61.07 370.6 –51.38
15 2 2 3 2 1089 –60.74 340 –50.63
16 2 3 1 1 1313 –62.37 578.3 –55.24
17 2 3 2 2 1310 –62.35 328 –50.32
18 2 3 3 3 1218 –61.71 388.5 –51.79
19 3 1 1 3 1631 –64.25 627.1 –55.95
20 3 1 2 1 1655 –64.37 434.1 –52.75
21 3 1 3 2 1374 –62.76 415.4 –52.37
22 3 2 1 1 1826 –65.23 496.2 –53.91
23 3 2 2 2 1477 –63.39 492.8 –53.85
24 3 2 3 3 1379 –62.79 547.8 –54.77
25 3 3 1 2 1835 –65.27 674.1 –56.57
26 3 3 2 3 1629 –64.24 641 –56.14
27 3 3 3 1 1441 –63.17 492.2 –53.84
The arithmetic average of S/N ratios for the levels of
each control factor was calculated with respect to the
thrust force and the cutting torque (Table 5). In addition,
after arranging the difference between the maximum and
minimum S/N ratios for each factor in a descending
order, the degree of influence of each factor on the thrust
force or cutting torque was found. Accordingly, the
effective control factors for the thrust force were particle
fraction, feed rate, spindle speed and drill-bit material
(Table 5). The optimum machining parameters for the
thrust force and drilling torque are found at the level
where each factor has the largest S/N ratio.14 Therefore,
the optimal machining conditions for the thrust force
were found to be the particle fraction of 10 %, feed rate
of 0.1 mm/r, spindle speed of 2000 r/min and drill
material of TiAlN-coated carbide. Similarly, the optimal
machining conditions for the drilling torque were found
to be the particle fraction of 10 %, feed rate of 0.1 mm/r,
spindle speed of 2000 r/min and a HSS drill.
The effect graph of each control factor for the thrust
force and drilling torque, according to the mean
responses, was given in Figures 8a and 8b, respectively.
Both Figures 8a and 8b showed that the thrust force and
A. TAÞKESEN, K. KÜTÜKDE: OPTIMIZATION OF THE DRILLING PARAMETERS FOR THE CUTTING FORCES ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 169–176 173
Table 3: Factors and levels used in the experiments
Tabela 3: Faktorji in stopnje, uporabljene pri eksperimentu
Process parameters Units
Levels
Level 1 Level 2 Level 3
Particle fraction
(A) % 10 15 25
Feed rate (B) mm/r 0.1 0.2 0.3
Spindle speed (C) r/min 1000 1500 2000
Drill material (D) HSS Carbide TiAlN-coatedcarbide
drilling torque increased with an increase in the particle
fraction and the feed rate, while the thrust force and
cutting torque decreased with an increase in the spindle
speed. However, the effect of the drill-bit material on the
thrust force was very low.
Table 5: Average S/N ratios for each factor and level with regard to
thrust force and cutting torque
Tabela 5: Povpre~je razmerja S/N za vsak faktor in stopnjo glede na
potisno silo in navor pri rezanju
Fo
r
th
ru
st
fo
rc
e Level A B C D
1 –59.4564* –60.8976* –62.1689 –61.5337
2 –61.3091 –61.4479 –61.5665 –61.6429
3 –63.941 –62.361 –60.9711* –61.53*
Difference 4.4846 1.4634 1.1978 0.1092
rank 1 2 3 4
Fo
r
to
rq
ue
1 –50.6201* –51.5226* –53.9199 –52.0704*
2 –52.1194 –52.4621 –52.194 –52.2421
3 –54.462 –53.2168 –51.0877* –52.8891
Difference 3.8419 1.6942 2.8322 0.8187
rank 1 3 2 4
* = Optimal level
3.4 Analysis of Variance (ANOVA)
The purpose of the analysis of variance was to
determine which parameter significantly affects the
cutting forces.28 ANOVA was performed to find whether
individual factors and their interactions that affect the
cutting forces were meaningful. According to the
ANOVA results presented in Table 6, the most
influential factor for the thrust force was found to be the
particle fraction of 80.23 %. The other important factors
were feed rate (6.72 %) and spindle speed (6.73 %).
Similarly, the most influential factor for the drilling
torque was found to be the particle fraction of 45.99 %,
followed by spindle speed (25.25 %) and feed rate (8.75
%) as seen in Table 7. In addition, the effect of the
drill-bit material on the cutting forces was found to be
small. Ftest values for the cutting forces, with regard to
the factor interactions, were not meaningful since they
were smaller than Ftable values.27 Hence, the statistical
significance of interactions was minimum and it could be
neglected.
Table 6: ANOVA results for the thrust force
Tabela 6: Rezultati ANOVA za potisno silo
Factor DF SS V Ftest PD
Particle fraction (A) 2 1882587 941294 210.5 80.23
Feed rate (B) 2 157716 78858 17.63 6.721
Spindle speed (C) 2 157829 78915 17.65 6.726
Drill material (D) 2 1250 624.9 0.1397 0.053 3
AxB 4 25505 6376 1.426 1.087
AxC 4 79868 19967 4.465 3.404
BxC 4 15006 3751 0.8388 0.639 5
Error 6 26835 4472 1.144
Total 26 2346595 100
DF: Degree of Freedom, SS: Sum of Squares, V: Variance, PD: Per-
centage Distribution. Ftable(0.05;2;6 ) = 5.14, Ftable(0.05;4;6) = 4.53
Table 7: ANOVA results for the drilling torque
Tabela 7: Rezultati ANOVA za navor pri vrtanju
Factor DF SS V Ftest PD
Particle fraction (A) 2 163615 81808 15.92 45.99
Feed rate (B) 2 31123 15562 3.029 8.749
Spindle speed (C) 2 89834 44917 8.742 25.25
Drill material (D) 2 10230 5115 0.995 2.876
AxB 4 11959 2990 0.582 3.362
AxC 4 12466 3116 0.607 3.504
BxC 4 5689 1422 0.277 1.599
Error 6 30829 5138 8.666
Total 26 355745 100
DF: Degree of Freedom, SS: Sum of Squares, V: Variance, PD: Per-
centage Distribution. Ftable(0.05;2;6) = 5.14, Ftable(0.05;4;6) = 4.53
3.5 Confirmation Experiments
The final step of the Taguchi experimental design
process includes confirmation experiments.14,27 For this
aim, the results of the experiments were compared with
the predicted values with the Taguchi method and the
error rates were obtained. S/N ratios predict were pre-
dicted using the following model:14,26
   predict = + −
=
∑m i m
i
k
( )
1
(2)
A. TAÞKESEN, K. KÜTÜKDE: OPTIMIZATION OF THE DRILLING PARAMETERS FOR THE CUTTING FORCES ...
174 Materiali in tehnologije / Materials and technology 47 (2013) 2, 169–176
Figure 8: Mean effect graphs of responses: a) thrust force, b) drilling
torque
Slika 8: Graf u~inka povpre~nih rezultatov: a) potisna sila, b) navor
pri vrtanju
where m is the total mean of the S/N ratios, i is the
mean S/N ratio at the optimum level and k is the number
of the main design parameters that significantly affect
the performance characteristics.
After predicting the S/N ratios other than 27 experi-
ments (with Eq.2), the thrust forces and drilling torques
were calculated using the following equation:26
Y
S N
predict =
−
10 20
(
/
)
(3)
where Ypredict is the thrust force or drilling torque with
regard to the S/N ratio. Figure 9 represents the compa-
rison between the predicted and experimental results
according to the experiment numbers. These results
show that the Taguchi method can be applied success-
fully in predicting the thrust forces with the coefficient
of determination R2 = 0.946 (Figure 7a). The value of
R2 for the prediction of the cutting torques was 0.644,
but the R2 value for 54 of 100 experiments was 0.85.
4 CONCLUSIONS
In this study, aluminum MMCs containing three
different weight fractions of B4C particles were produced
with the PM technique, and drilling experiments were
carried out to study the effects of the machining para-
meters on the thrust force and cutting torque. Moreover,
the optimum drilling parameters were obtained for the
performance characteristics (thrust force and torque)
using the Taguchi analysis. The obtained results can be
summarized as follows:
• An increase in the proportion of the B4C particle
caused a decreased ductility of the material but an
increased hardness of the composite. The highest
tensile strength was obtained with the 15 % B4C par-
ticle fraction.
• According to the experimental results, the cutting
forces significantly increased with an increase in the
B4C fraction and the feed rate but decreased with an
increase in the spindle speed.
• HSS tools produced more thrust forces than the two
carbide tools especially when drilling the composites
with higher particle fractions. On the other hand, the
coated and the solid carbide tools produced similar
thrust forces. However, the coated tools produced
somewhat higher drilling torques than the uncoated
ones.
• With the Taguchi and ANOVA analysis, the effective
factors for the thrust force and drilling torque were
found to be the particle-weight fraction and feed rate,
respectively. Furthermore, the effects of the cutting
tool material and the interactions of the factors on the
thrust force and cutting torque were found to be very
low.
Acknowledgments
This research was supported by the Gazi University
under the Project Number 07/2008-8.
The authors wish to thank the TOBB Economy and
Technology University for providing laboratory opportu-
nities during the course of the research work. The
authors express their gratitude to MÝTAÞ CÝVATA and
Mr. Serdar Iskender for enabling the tensile and impact
tests for the fabricated composite materials.
5 REFERENCES
1 R. G. Budynas, R. Budynas, K. Nisbett, Shigley’s Mechanical Engi-
neering Design, McGraw-Hill, 2010
2 V. Songmene, M. Balazinski, Machinability of graphitic metal matrix
composites as a function of reinforcing particles, CIRP Annals -
Manufacturing Technology, 48 (1999) 1, 77–80
3 M. Ramulu, P. N. Rao, H. Kao, Drilling of (Al2O3)p/6061 metal
matrix composites, Journal of Materials Processing Technology, 124
(2002) 1–2, 244–254
4 S. Basavarajappa, G. Chandramohan, J. P. Davim, Some studies on
drilling of hybrid metal matrix composites based on Taguchi tech-
niques, Journal of Materials Processing Technology, 196 (2008) 1–3,
332–338
5 C. A. Brown, M. K. Surappa, The machinability of a cast aluminium
alloy-graphite particle composite, Materials Science and Engineer-
ing, 102 (1988) 1, 31–37
6 S. Barnes, I. R. Pashby, A. B. Hashim, Effect of heat treatment on the
drilling performance of aluminum/SiC MMC, Applied Composite
Materials, 6 (1999) 2, 121–138
A. TAÞKESEN, K. KÜTÜKDE: OPTIMIZATION OF THE DRILLING PARAMETERS FOR THE CUTTING FORCES ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 169–176 175
Figure 9: Comparison of the predicted and experimental results:
a) thrust force (R2 = 0.946), b) drilling torque (R2 = 0.644)
Slika 9: Primerjava napovedanih in eksperimentalno dolo~enih
rezultatov: a) potisna sila (R2 = 0,946), b) navor pri vrtanju (R2 =
0,644)
7 S. Barnes, I. R. Pashby, Through-tool coolant drilling of aluminum/
SiC metal matrix composite, Journal of Engineering Materials and
Technology, Transactions of the ASME, 122 (2000) 4, 384–388
8 J. P. Davim, A. Monteiro Baptista, Cutting force, tool wear and sur-
face finish in drilling metal matrix composites, Proceedings of the
Institution of Mechanical Engineers, Part E: Journal of Process
Mechanical Engineering, 215 (2001) 2, 177–183
9 A. M. Abrão, J. C. C. Rubio, P. E. Faria, J. P. Davim, The effect of
cutting tool geometry on thrust force and delamination when drilling
glass fibre reinforced plastic composite, Materials and Design, 29
(2008) 2, 508–513
10 M. T. Hayajneh, A. M. Hassan, A. T. Mayyas, Artificial neural net-
work modeling of the drilling process of self-lubricated aluminum/
alumina/graphite hybrid composites synthesized by powder
metallurgy technique, Journal of Alloys and Compounds, 478 (2009)
1–2, 559–565
11 D. Iliescu, D. Gehin, M. E. Gutierrez, F. Girot, Modeling and tool
wear in drilling of CFRP, International Journal of Machine Tools and
Manufacture, 50 (2010) 2, 204–213
12 Y. Altunpak, M. Ay, S. Aslan, Drilling of a hybrid Al/SiC/Gr metal
matrix composites, International Journal of Advanced Manufacturing
Technology, 60 (2012) 5–8, 513–517
13 A. N. Haq, P. Marimuthu, R. Jeyapaul, Multi response optimization
of machining parameters of drilling Al/SiC metal matrix composite
using grey relational analysis in the Taguchi method, International
Journal of Advanced Manufacturing Technology, 37 (2008) 3–4,
250–255
14 G. Tosun, Statistical analysis of process parameters in drilling of
AL/SIC P metal matrix composite, International Journal of Advan-
ced Manufacturing Technology, 55 (2011) 5–8, 477–485
15 A. R. Ahamed, P. Asokan, S. Aravindan, M. K. Prakash, Drilling of
hybrid Al-5%SiCp-5%B4Cp metal matrix composites, International
Journal of Advanced Manufacturing Technology, 49 (2010) 9–12,
871–877
16 S. Basavarajappa, G. Chandramohan, J. P. Davim, M. Prabu, K.
Mukund, M. Ashwin, M. Prasannakumar, Drilling of hybrid alumi-
nium matrix composites, International Journal of Advanced Manu-
facturing Technology, 35 (2008) 11–12, 1244–1250
17 C. C. Tsao, H. Hocheng, Effect of tool wear on delamination in
drilling composite materials, International Journal of Mechanical
Sciences, 49 (2007) 8, 983–988
18 Y. Kazancoglu, U. Esme, M. Bayramoglu, O. Guven, S. Ozgun,
Multi-objective optimization of the cutting forces in turning ope-
rations using the Grey-based Taguchi method, Mater. Tehnol., 45
(2011) 2, 105–110
19 H. Zhang, M. W. Chen, K. T. Ramesh, J. Ye, J. M. Schoenung, E. S.
C. Chin, Tensile behavior and dynamic failure of aluminum
6092/B4C composites, Materials Science and Engineering A, 433
(2006) 1–2, 70–82
20 E. Mohammad Sharifi, F. Karimzadeh, M. H. Enayati, Fabrication
and evaluation of mechanical and tribological properties of boron
carbide reinforced aluminum matrix nanocomposites, Materials and
Design, 32 (2011) 6, 3263–3271
21 J. S. Strenkowski, C. C. Hsieh, A. J. Shih, An analytical finite
element technique for predicting thrust force and torque in drilling,
International Journal of Machine Tools and Manufacture, 44 (2004)
12–13, 1413–1421
22 S. Jayabal, U. Natarajan, Influence of cutting parameters on thrust
force and torque in drilling of E-glass/polyester composites, Indian
Journal of Engineering and Materials Sciences, 17 (2010) 6, 463–470
23 J. P. Davim, Study of drilling metal-matrix composites based on the
Taguchi techniques, Journal of Materials Processing Technology, 132
(2003) 1–3, 250–254
24 C. C. Tsao, Taguchi analysis of drilling quality associated with core
drill in drilling of composite material, International Journal of
Advanced Manufacturing Technology, 32 (2007) 9–10, 877–884
25 C. C. Tsao, H. Hocheng, Evaluation of thrust force and surface
roughness in drilling composite material using Taguchi analysis and
neural network, Journal of Materials Processing Technology, 203
(2008) 1–3, 342–348
26 R. K. Roy, A primer on the Taguchi method / Ranjit K. Roy, Van No-
strand Reinhold, New York, 1990
27 K. Palanikumar, Experimental investigation and optimization in drill-
ing of GFRP composites, Measurement, Journal of the International
Measurement Confederation, 44 (2011) 10, 2138–2148
28 U. Esme, Use of grey based Taguchi method in ball burnishing
process for the optimization of surface roughness and microhardness
of AA 7075 aluminum alloy, Mater. Tehnol., 44 (2010) 3, 129–135
A. TAÞKESEN, K. KÜTÜKDE: OPTIMIZATION OF THE DRILLING PARAMETERS FOR THE CUTTING FORCES ...
176 Materiali in tehnologije / Materials and technology 47 (2013) 2, 169–176
A. AYDAY, M. DURMAN: WEAR PROPERTIES OF AISI 4140 STEEL MODIFIED WITH ELECTROLYTIC-PLASMA ...
WEAR PROPERTIES OF AISI 4140 STEEL MODIFIED
WITH ELECTROLYTIC-PLASMA TECHNOLOGY
OBRABNE LASTNOSTI JEKLA AISI 4140, MODIFICIRANEGA S
TEHNOLOGIJO ELEKTROLITSKE PLAZME
Aysun Ayday, Mehmet Durman
Sakarya University, Faculty of Engineering, Department of Metallurgical and Materials Engineering, 54187 Sakarya, Turkey
aayday@sakarya.edu.tr
Prejem rokopisa – received: 2012-08-29; sprejem za objavo – accepted for publication: 2012-10-09
An electrolytic-plasma treatment (EPT) was applied to the surface of AISI 4140 steel and the wear behavior under dry sliding
conditions was studied for different treatment parameters. The modified samples were characterized before and after the wear
testing using metallographic, SEM-microscope and microhardness techniques. The test results indicate that the wear resistance
of the AISI 4140 steel can be improved by means of electrolytic-plasma technology (EPT). The wear resistance increases with
an increased modified-layer hardness due to a transformation to the martensitic structure.
Keywords: plasma, microhardness, wear
Preu~evane so bile obrabne lastnosti pri suhem drsenju in razli~nih parametrih z elektrolitsko plazmo obdelane povr{ine jekla
AISI 4140. Modificirani vzorci so bili karakterizirani pred preizkusom obrabe in po njem z uporabo metalografije, z vrsti~nim
elektronskim mikroskopom (SEM) in meritvijo trdote. Rezultati so pokazali, da se obrabna odpornost jekla AISI 4140 lahko
pove~a s tehnologijo elektrolitske plazme (EPT). Obrabna odpornost se pove~uje s povi{anjem trdote modificirane plasti zaradi
pretvorbe v martenzitno mikrostrukturo.
Klju~ne besede: plazma, mikrotrdota, obraba
1 INTRODUCTION
As an advanced surface-processing technique, elec-
trolytic-plasma treatment (EPT) has been successfully
used to improve the hardness, the wear resistance and the
corrosion resistance of materials1. When hardening is not
necessary for a whole surface or bulk of material, EPT is
a suitable method for treating a specific location on a
surface2. Electrolytic-plasma (heating-quenching) har-
dening is a standard hardening mechanism involving two
main steps: "austenitizing", during which the material is
heated above the critical temperature for the austenite
formation (but below the melting point) and "quenching"
or cooling down, where austenite is transformed into
martensite. The heating or quenching of medium and
high-carbon steels can change the steel microstructures,
which causes variations in the mechanical and physical
properties and affects the behavior of the steels under
service conditions and operations3,4.
EPT is characterized by several process parameters:
voltage, current, electrolyte, duration, and heating-
quenching rate. All these parameters are strongly corre-
lated to each other and affect the final hardening results;
for this reason process modeling seems to be a good
approach to the process optimization. In this study, the
wear resistance of the electrolytic-plasma-modified AISI
4140 steel was evaluated under dry sliding conditions
and compared with the AISI 4140 steel samples. The
modified samples were characterized before and after the
wear tests with metallographic, SEM microscope and
microhardness techniques.
2 EXPERIMENTAL DETAILS
The test material was the commercial AISI 4140
low-alloy steel with the composition (in mass fractions,
%) of 0.4 C, 0.22 Si, 0.77 Mn, 0.04 S, 0.035 P, 0.8 Cr
and 0.25 Si. The diameter of cylindrical samples was 20
mm and the height was 10 mm. All the samples were
modified with EPT. The EPT voltage, heating and cool-
ing times were 310–260 V, 3 s and 3 s, respectively,
depending upon the thermal cycle and the process
temperature. The sample codes and the EPT parameters
are listed in Table 1.
The morphology of the modification layer was
investigated with a scanning electron microscope (SEM
Joel, JSM 6060-LU). The hardness measurements were
conducted on the cross-sections of the samples with a
Vickers microhardness tester. The test load was 100 g for
the hardness measurements at the cross-sections. The
temperature distribution of the samples from the
plasma-treated side to the internal side was investigated
via thermocouples during the process. The surface
temperature data were collected from the system with the
aid of a computer data-acquisition system.
The wear tests were performed both on the original
AISI 4140 and on the EPT-modified specimens to
determine the optimum process parameters. All the wear
Materiali in tehnologije / Materials and technology 47 (2013) 2, 177–180 177
UDK 669.14:533.9:539.92 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)177(2013)
tests were carried out under dry sliding conditions at
room temperature using a ball-on-disc (CSM tribo-
meter), friction- and wear-test machine. The counterpart
was an Al2O3 ball (
 = 6 mm) according to DIN 50 324
and ASTM G 99-95a. The tests were performed with a
nominal load of 3 N and a sliding speed of 0.10 m/s for
the total sliding distance of 200 m.
3 RESULTS AND DISCUSSION
The cross-sectional SEM images of the EPT-6 sample
showed that it typically consisted of a modified diffusion
zone (Figure 1). This figure shows the microstructure of
the cross-section of the EPT-modified steel. The follow-
ing zones are visible in the diffusion layer: the hardened
zone (HZ), the heat-affected zone (HAZ) and the base
material (BM). During EPT, austenite transforms com-
pletely or partially to martensite and thus the micro-
A. AYDAY, M. DURMAN: WEAR PROPERTIES OF AISI 4140 STEEL MODIFIED WITH ELECTROLYTIC-PLASMA ...
178 Materiali in tehnologije / Materials and technology 47 (2013) 2, 177–180
Table 1: EPT parameters
Tabela 1: EPT-parametri
Parameter code Electrolytic solution Heating (V) Cooling (V) Total time (s) Thermal cycle
EPT–0 Original AISI 4140
EPT–4 Na2CO3; 12 % 310 260 (3 and 3) s × 4 = 24 s 4
EPT–5 Na2CO3; 12 % 310 260 (3 and 3) s × 5 = 30 s 5
EPT–6 Na2CO3; 12 % 310 260 (3 and 3) s × 6 = 36 s 6
Table 2: Maximum surface-hardness, surface-temperature and wear-rate values
Tabela 2: Maksimalna trdota povr{ine, temperatura povr{ine, vrednosti obrabe
Original AISI 4140 (EPT-0) EPT-4 EPT-5 EPT-6
Max. surface hardness (HV0.1) 200 800 900 930
Surface temperature (°C) ø 600 780 835
Wear rate (mm3/(N m)) 9.06 E-05 5.05 E-05 4.87 E-05 4.66 E-05
Figure 1: Cross-sectional appearance of the EPT-modified EPT-6 (R – radius of the modification area, H – HZ depth)
Slika 1: Pre~ni prerez vzorca EPT-6, modificiranega z EPT (R – polmer modificiranega podro~ja, H – globina HZ)
Figure 2: Microhardness profiles of the EPT-modified samples
Slika 2: Profili mikrotrdote vzorcev, modificiranih z EPT
structure of the HZ consists of martensite. An amount of
the retained austenite may be present in this region5. In
the neighborhood of the HZ with the base material, a
narrow heat-affected zone was observed, consisting of
martensite, bainite and some traces of the initial pearlitic
structure. These are the most probable structures accord-
ing to Refs. 5,6. The microstructure of the base material is
composed of ferrite and pearlite. The maximum
microhardness of this hardened zone was 930 HV0.1. The
maximum microhardness of the second zone (HAZ) was
800 HV0.1. From the plasma surface to the base metal,
the hardness values decreased and the phase structure
changed into a ferritic-pearlitic matrix; in this zone the
hardness was measured as 200 HV0.1, which is shown in
Figure 2.
The values obtained for the maximum surface
hardness, surface temperature and wear rate of different
specimen groups are listed in Table 2. It is evident that
EPT markedly improves the wear performance of the
steel and that the degree of improvement depends on the
EPT parameters. This means that the wear resistance
increases with an increase in the thermal cycle. The
thermal cycle was effective at rising the surface tempe-
rature. A high surface temperature affects the modifi-
cation depth and the surface hardness. The wear rate
obtained for the EPT-6 sample was lower than the rates
obtained for the original AISI 4140 (EPT-0) or for the
EPT-4 and EPT-5 samples. This is due to the micro-
structure of the EPT-6 sample, which had a martensitic
structure. The microstructure of the EPT-6 surface, after
the EPT treatment, was finer and more homogenous in
comparison with the surfaces of EPT-4 and EPT-5.
Figure 3 shows the wear surfaces of the original
AISI 4140 (EPT-0), the EPT-4 and EPT-6 samples, tested
at a load of 3 N. EPT-6 shows quite a smooth surface
with shallow, abrasive wear scars due to the high
hardness of the sample. The original AISI 4140 (EPT-0)
steel was tested under a similar wear-test condition. The
plastic deformation was obvious in this case and was
caused by a low surface hardness as shown in Figure 3.
The worn-surface analysis revealed a severe abrasive
wear of the original AISI 4140 (EPT-0) accompanied
with a high degree of plastic deformation (Figure 3
–AISI 4140).
4 CONCLUSIONS
It is evident that the wear rate of steel is increased
significantly by EPT. The degree of improvement
depends on the EPT-process conditions. The modified
layer thickness and the surface hardness increase with
the increasing surface temperature and thermal cycle.
The specimens with the maximum hardness showed the
maximum resistance to wear. Thus, the hardness of the
surface is a very important factor with respect to the
wear rate. The hardness results arise from the micro-
structures of the modified samples that had martensitic
structures. An increase in the thermal cycle increases the
wear resistance of EPT-6 due to a finer and more
homogenous hardened zone. It was observed that the
longer the heating and the cooling times, the greater was
the hardened-layer thickness. The initial microstructure
was fully martensitic, especially in the HAZ zone after
the EPT-6 sample processing. When processing the other
A. AYDAY, M. DURMAN: WEAR PROPERTIES OF AISI 4140 STEEL MODIFIED WITH ELECTROLYTIC-PLASMA ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 177–180 179
Figure 3: SEM micrographs of the original surface (EPT-0) and worn surfaces of AISI 4140, EPT-4 and EPT-6
Slika 3: SEM-posnetki originalne povr{ine (EPT-0) in obrabljene povr{ine AISI 4140, EPT-4 in EPT-6
samples, like EPT-4 and EPT-5, the microstructure
transforms to a martensitic and also bainitic matrix.
Therefore, the maximum microhardness values decrease
from 930 HV to 200 HV during the experiments.
5 REFERENCES
1 Y. N. Tyurin, A. D. Pogrebnyak, Electric Heating Using a Liquid
Electrode, Surface and Coatings Technology, 142–144 (2001),
293–299
2 C. Ye, S. Suslov, B. J. Kim, E. A. Stach, G. J. Cheng, Fatigue
performance improvement in AISI 4140 steel by dynamic strain
aging and dynamic precipitation during warm laser shock peening,
Acta Materialia, 59 (2011), 1014–1025
3 S. A. Jenabali Jahromi, A. Khajeh, B. Mahmoudi, Effect of different
pre-heat treatment processes on the hardness of AISI 410 martensitic
stainless steels surface-treated using pulsed neodymium-doped
yttrium aluminum garnet laser, Materials and Design, 34 (2012),
857–862
4 N. S. Bailey, W. Tan, Y. C. Shin, Predictive modeling and experi-
mental results for residual stresses in laser hardening of AISI 4140
steel by a high power diode laser, Surface & Coatings Technology,
203 (2009), 2003–2012
5 C. Soriano, J. Leunda, J. Lambarri, V. García Navas, C. Sanz, Effect
of laser surface hardening on the microstructure, hardness and
residual stresses of austempered ductile iron grades, Applied Surface
Science, 257 (2011), 7101–7106
6 C. T. Kwok, K. I. Leong, F. T. Cheng, H. C. Man, Microstructural
and corrosion characteristics of laser surface-melted plastics mold
steels, Materials Science and Engineering A, 357 (2003), 94–103
A. AYDAY, M. DURMAN: WEAR PROPERTIES OF AISI 4140 STEEL MODIFIED WITH ELECTROLYTIC-PLASMA ...
180 Materiali in tehnologije / Materials and technology 47 (2013) 2, 177–180
Y. SERT, N. TOPLAN: TRIBOLOGICAL BEHAVIOR OF A PLASMA-SPRAYED Al2O3-TiO2-Cr2O3 COATING
TRIBOLOGICAL BEHAVIOR OF A PLASMA-SPRAYED
Al2O3-TiO2-Cr2O3 COATING
TRIBOLO[KO PONA[ANJE S PLAZMO NAPR[ENEGA
Al2O3-TiO2-Cr2O3 NANOSA
Yeºim Sert, Nil Toplan
Sakarya University, Department of Metallurgical and Materials Engineering, Esentepe Campus, 54187 Sakarya, Turkey
toplan@sakarya.edu.tr
Prejem rokopisa – received: 2012-08-31; sprejem za objavo – accepted for publication: 2012-10-08
Alumina-titania, titania, chromia and chromia-titania coatings, deposited on aluminium substrates with atmospheric,
plasma-spray, coating techniques (APSCTs), were tested on a low-frequency reciprocating-sliding tribometer. The surface wear
of the coatings was investigated with SEM and optical microscopy. The denser Cr2O3 coatings showed a higher wear resistance
than the more porous Al2O3-TiO2 and TiO2 coatings. An increase in the titania content diminishes the coating hardness and the
wear resistance.
Keywords: Cr2O3-TiO2, Al2O3-TiO2 coatings, textile parts, plasma-spray coating, wear
Nanosi Al2O3-TiO2, TiO2, Cr2O3, Cr2O3-TiO2 z atmosferskim plazemskim nana{anjem z brizganjem (APSCT) na podlago iz
aluminija so bili preizku{eni na nizkofrekven~nem protismernem drsnem tribometru. Obrabljena povr{ina nanosa je bila
preiskovana s SEM in svetlobno mikroskopijo. Gostej{i nanos Cr2O3 je pokazal ve~jo odpornost proti obrabi kot bolj porozna
nanosa Al2O3-TiO2 in TiO2. Pove~anje vsebnosti TiO2 zmanj{uje trdoto nanosa in zmanj{a odpornost proti obrabi.
Klju~ne besede: nanosi Cr2O3-TiO2, Al2O3-TiO2, kosi tkanine, plazemski nanos z brizganjem, obraba
1 INTRODUCTION
It is well known that aluminum (Al) alloys have been
considered to be some of the most useful and versatile
materials because of their metallurgical characteristics,
such as high strength-to-weight ratio, and high thermal
conductivity. They are also easy to shape and relatively
inexpensive. However, the low hardness results in poor
tribological characteristics and prevents their wide use,
especially in the situations where a hard surface is need-
ed. To improve the wear resistance, many techniques,
such as metal-matrix composites, plasma spraying, ther-
mal spraying and hard anodizing have been explored1.
The APSCT is an economical and effective method
applied to various machine parts to improve the com-
ponent performance in wear, corrosion, thermal barrier,
and electric insulation. Plasma-sprayed Al2O3-TiO2 has
been widely used as a wear-resistant coating in textile,
machinery and printing industries2–5. Cr2O3 has a wide
range of applications such as green pigments, coating
materials for thermal protection and wear resistance as
well as refractory applications due to the high melting
temperature (about 2435 °C)3,6. The present paper deals
with the wear resistance of the plasma-sprayed alu-
mina-titania, titania, chromia and chromia-titania
coatings that increased the service life of the shutters
(Al-based) used in the textile industry.
2 EXPERIMENTAL PROCEDURE
The commercial feedstock powders in the mass
fractions 13 % TiO2-Al2O3 (Metco 130), 40 %
TiO2-Al2O3 (Metco 131VF), 100 % TiO2 (Metco 102)
and 100 % Cr2O3 (Metco 106) were supplied by
SULZER METCO Powder Technology. Al2O3, TiO2 and
Cr2O3 powders were premixed to form five different
compositions (Table 1) and these were prepared on an
aluminium alloy (AA1050). The mixtures were ball-
milled for 2 h by using ZrO2 balls and distilled water as
the milling media to provide homogenous mixtures.
After drying the powders were screened and sieved to
Materiali in tehnologije / Materials and technology 47 (2013) 2, 181–183 181
UDK 621.793:533.9:539.92 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)181(2013)
Table 2: Plasma-spray parameters
Tabela 2: Parametri nabrizgavanja s plazmo
Coating parameters (for the 1–5 coded compositions)
Primary-gas flow rate (Ar, L/min) 80
Secondary-gas flow rate (H2, L/min) 15
Carrier-gas flow rate (Ar, L/min) 40
Spray distance (mm) 100
Current (A)
Voltage (V)
500
60–70
Table 1: Coating powders (w/%)
Tabela 1: Prahovi za nana{anje (w/%)
Composition Al2O3 TiO2 Cr2O3
1 87 13 –
2 60 40 –
3 – 100 –
4 – 50 50
5 – – 100
achieve the correct particle-size distribution needed for
plasma spraying. Prior to the deposition process,
aluminium substrates were grit blasted with Al2O3
particles and this was followed by ultrasonic cleaning in
acetone for 15 min. A 40 kW plasma-spraying system
(METCO 3MB) was utilized to produce the coatings
using the parameters summarized in Table 2.
The surface roughness was measured with a surface-
roughness tester (Perthometer M4P) and the average
roughness Ra, defined as the arithmetic mean of the
departures of the profile from the mean line, was used to
quantify the coating-surface roughness. A scanning
electron microscope (SEM) (JEOL JSM-6060LV)
equipped with an energy dispersive X-ray spectrometer
(EDS) was used to examine the microstructures and
chemical compositions of the coatings. An X-ray
diffraction analysis (XRD) was carried out on RIGAKU
DMAX 2200 to determine the phases of the coating(s).
The microhardness values of the specimens were taken
from the cross-sections of the polished samples at the
load of 200 g and after a loading time of 15 s using
LEICA VMHT MOT microhardness equipment. The
wear tests were performed using a low-frequency reci-
procating-sliding tribometer, connected to a computer
monitoring the dynamic coefficient of friction (in both
sliding directions), relative humidity and temperature.
The tests were performed by applying a load of 5 N to a
single-crystal Al2O3 (sapphire) ball with a diameter of
6 mm. The wear specimens had the dimensions of 3 cm
× 3 cm × 3 mm, the shear rate was 0.15 m/s and the
sliding distance was 150 m. The values of the coefficient
of friction were calculated from the normal load and the
friction force was obtained from a digital oscilloscope.
After the wear tests, the morphology of each wear scar
was observed with SEM.
3 RESULTS AND DISCUSSION
Table 3 summarizes five different compositions of
the coating-test results. While the TiO2 coatings had the
highest surface-roughness values, the values for the
Cr2O3 coatings were found to be the lowest. Having a
lower as-sprayed surface roughness is very important for
technological applications because it reduces the number
of post-deposition mechanical treatments necessary7,8.
As a function of the substrate-surface roughness, the
values of porosity and coating roughness increased,
while the increase in the substrate-surface roughness
grow up. The hardness values were also relatively
reduced. The highest hardness, the lowest porosity and
the lowest coating roughness were obtained at the value
of the substrate roughness of 2.346 μm. The hardness
values decrease with the increasing amounts of TiO2 in
the Al2O3-TiO2 coatings. In the Cr2O3-TiO2-based coat-
ings, the hardness values increase with the amount of
Cr2O3. An increase in the porosity amount will result in a
decrease in the hardness of the coating. The lowest
coefficient of friction (μ) was achieved in the 100 %
Cr2O3 coatings. Hardness has a strong influence on wear
resistance. The higher the hardness, the better is the wear
resistance. It is well known that an addition of TiO2 to an
Al2O3 coating is to reduce the melting temperature of the
oxide, thereby producing less porous and better perfor-
mance coatings than the pure Al2O3 coatings. The
melting temperature decreases due to the fact that TiO2
has a lower melting temperature (1854 °C) than Al2O3
(2040 °C) and due to its ability to form a liquid solution
with Al2O3. It is also noted that the trend displayed by
the coating densities is consistent with that exhibited by
the degree of melting, i.e., a high degree of melting (e.g.,
Cr2O3, 2435 °C) results in high density. Increasing the
microhardness leads to the improvements in the wear
resistance of the coatings. The grain size also has an
effect on the wear resistance. The nanocoating has a
higher wear resistance than the commercial coating
although its hardness is lower than that of the commer-
cial coating. A related study on the abrasive wear has
revealed that nanocoatings could have a two-to-four-fold
increase in the wear resistance in comparison with the
commercial coatings2,7.
Table 3: Surface roughness, hardness, density and friction coefficient
of the coatings
Tabela 3: Hrapavost povr{ine, trdota, gostota in koeficient trenja
nanosa
Composition 1 2 3 4 5
Ra/μm 3.553 3.437 4.443 3.011 2.346
Hardness (HV)
Relative density (%)
1028
87.70
899
88.45
812
86.00
1010
90.12
1724
92.51
Average friction
coefficient 0.142 0.190 0.223 0.144 0.074
The XRD analysis of the starting powders showed
that the chromia powder consisted of an eskolaite phase
(Cr2O3) and the alumina-titania powder of -Al2O3 and
anatase. It was also clear from XRD that the chromia
coating consists of eskolaite, the chromia-titania coating
consists of eskolaite and Ti2Cr2O7 and the alumina-titania
coating consists mainly of -Al2O3 with some -Al2O3,
Al2TiO5, a glassy phase and a small amount of rutile. A
very low amount of crystalline TiO2 indicates that it
mostly dissolves in the molten Al2O3 2,8.
The main wear mechanisms of plasma-sprayed cera-
mic coatings were reported to be a plastic deformation,
crack formation and spalling due to fatigue, brittle
fracture and material transfer. In the reciprocated dry
sliding, wear debris was considerably involved in the
wear process in the steady state. The worn surfaces of
the Cr2O3 and TiO2 coatings were observed with SEM at
different magnitudes (Figure 1). In the SEM images the
wear scar of the TiO2 coating was much larger than that
of the Cr2O3 coating. The Cr2O3 coating is the hardest
and the most anisotropic among the plasma-sprayed
ceramics due to its low interlamellar cohesion; the
Al2O3-TiO2 and TiO2 coatings are the most isotropic but
Y. SERT, N. TOPLAN: TRIBOLOGICAL BEHAVIOR OF A PLASMA-SPRAYED Al2O3-TiO2-Cr2O3 COATING
182 Materiali in tehnologije / Materials and technology 47 (2013) 2, 181–183
also less hard and less tough due to the formation of an
alumina-titania glassy phase which favors intersplat
adhesion but turns out to be quite brittle8. The smooth
film, formed due to a large plastic deformation of the
adhered wear debris, strongly influenced the friction and
wear behavior. For the plasma-sprayed Cr2O3 coatings,
similar wear mechanisms were reported under dry
sliding conditions and the role of the wear-protective
film formed by a plastic deformation of the adhered and
compacted debris particles was discussed9. The abrasive
wear mechanism of the coatings does not only depend on
the coating hardness and density, but also on the particle
size, the type of the powder used, the coating micro-
structure, as well as on the microstructural change during
a wear testing. The average coarser powder particle size
causes an appearance of a significant number of
unsmelted particles.
4 CONCLUSIONS
Alumina-titania, titania, chromia and chromia-titania
coatings were deposited with APSCT to increase the
wear behavior of the aluminium-based shutters. While
the friction coefficient and the coating-surface roughness
increased with an increase in the titania content, the
coating density, hardness and wear resistance decreased.
5 REFERENCES
1 K. Huang, X. Lin, C. Xie, T. M. Yue, J. Mater. Sci. Technol., 23
(2007) 2, 201–206
2 M. Wang, L. L. Shaw, Surface and Coatings Technology, 202 (2007)
1, 34–44
3 V. P. Singh, A. Sil, R. Jayaganthan, Wear, 272 (2011), 149–158
4 S. Tao, Z. Yin, X. Zhou, C. Ding, Tribology International, 43 (2010),
69–75
5 S. T. Aruna, N. Balaji, J. Shedthi, V. K. W. Grips, Surface &
Coatings Technology, 208 (2012), 92–100
6 A. Cellard, V. Garnier, G. Fantozzi, G. Baret, P. Fort, Ceramics Inter-
national, 35 (2009), 913–916
7 L. L. Shaw, D. Goberman, R. Ren, M. Gell, S. Jiang, Y. Wang, T. D.
Xiao, P. R. Strutt, Surface and Coatings Technology, 130 (2000) 1,
1–8
8 G. Bolelli, V. Cannillo, L. Lusvarghi, T. Manfredini, Wear, 261
(2006), 1298–1315
9 H. S. Ahn, O. K. Kwon, Wear, 225–229 (1999), 814–824
Y. SERT, N. TOPLAN: TRIBOLOGICAL BEHAVIOR OF A PLASMA-SPRAYED Al2O3-TiO2-Cr2O3 COATING
Materiali in tehnologije / Materials and technology 47 (2013) 2, 181–183 183
Figure 1: a), b), c) Worn surface morphologies of the Cr2O3 and d), e), f) TiO2 coatings
Slika 1: a), b), c) Morfologija obrabljene povr{ine nanosa Cr2O3 in d), e), f) TiO2

M. OSTRÝ et al.: CHARACTERIZATION OF SELECTED PHASE-CHANGE MATERIALS ...
CHARACTERIZATION OF SELECTED PHASE-CHANGE
MATERIALS FOR A PROPOSED USE IN BUILDING
APPLICATIONS
KARAKTERIZACIJA IZBRANIH MATERIALOV S FAZNO
PREMENO ZA PREDLAGANO UPORABO V GRADBENI[TVU
Milan Ostrý1, Radek Pøikryl2, Pavel Charvát3
1Brno University of Technology, Faculty of Civil Engineering, Institute of Building Structures, Veveøí 95, 602 00 Brno, Czech Republic
2Brno University of Technology, Faculty of Chemistry, Purkyòova 464/188, 612 00 Brno, Czech Republic
3Brno University of Technology, Faculty of Mechanical Engineering, Technická 2, 616 69 Brno, Czech Republic
ostry.m@fce.vutbr.cz
Prejem rokopisa – received: 2012-09-02; sprejem za objavo – accepted for publication: 2012-09-28
The generally positive trend of ever-stricter requirements for the thermal insulation properties of building envelopes, leading to
a significant reduction in the heat losses of modern buildings, has also brought about some negative aspects. Modern
light-weight buildings with high-thermal-resistance envelopes are prone to overheating in the summer due to both solar and
internal heat gains. This problem is often solved by installing mechanical cooling (air-conditioning) that leads to an increase in
the energy consumption and, since electricity is mostly used to power the air-conditioning systems, the increase in the energy
consumption for cooling can offset the heating-energy savings in terms of primary energy. A lot of attention has therefore been
paid to the other means of temperature control in buildings, such as night-time ventilation and/or the building-integrated thermal
storage. The phase-change materials that can store a rather large amount of heat in a narrow temperature interval around their
melting point seem to be particularly suitable for this purpose. There are many ways of integrating PCMs into the building
structures as well as the techniques that employ that extra thermal-storage capacity to provide thermal comfort for the
occupants. This paper deals with the results of the laboratory testing of selected organic and inorganic phase-change materials
for integration into building structures. Differential scanning calorimetry was used to obtain the melting ranges and enthalpies of
fusion of the selected materials and thermogravimetry was used to explore the thermal stability (decomposition) of the materials
at higher temperatures.
Keywords: thermal-energy storage (TES), heat-storage medium, phase-change materials (PCMs), organic materials, inorganic
materials, sensible heat storage, latent-heat storage
Splo{ne pozitivne usmeritve v vedno ostrej{e zahteve pri toplotni izolaciji poslopij vodijo k ob~utnemu zmanj{anju toplotnih
izgub modernih zgradb in so prinesle tudi nekaj negativnih vidikov. Moderne, lahke zgradbe, z dobrim izolativnim ovojem so
nagnjene k pregrevanju v poletju zaradi son~ne in notranje toplote. Ta problem se pogosto re{uje s postavitvijo mehanskega
ohlajevanja (air-conditioning), ki povzro~i pove~anje porabe energije, saj je elektrika najbolj pogosto uporabljena za pogon
sistema ohlajanja, vendar pa se s stali{~a primarne energije pove~uje poraba energije za ohlajanje, ki lahko celo prese`e
prihranke pri energiji za ogrevanje. Mnogo pozornosti je treba zato posvetiti drugim sredstvom za kontrolo temperature v
zgradbah, kot so no~na ventilacija in/ali shranjevanje toplote integrirano v zgradbi. Materiali s fazno premeno lahko shranjujejo
relativno velike koli~ine toplote v temperaturnem intervalu okrog njihovega tali{~a in so zato videti posebno primerni za ta
namen. Mnogo na~inov je za vklju~itev PCM-materialov v strukturo zgradbe kot tudi tehnike, ki vklju~ujejo shranjevanje ekstra
toplotne kapacitete za zagotavljanje udobja stanovalcev. ^lanek obravnava rezultate laboratorijskih preizkusov izbranih
organskih in anorganskih materialov s fazno premeno za njihovo vklju~itev v strukturo zgradbe. Diferen~na dinami~na
kalorimetrija je bila uporabljena za dolo~anje podro~ja taljenja in entalpije taljenja izbranih materialov, termogravimetrija pa za
raziskovanje toplotne stabilnosti (dekompozicije) materialov pri povi{anih temperaturah.
Klju~ne besede: shranjevanje toplotne energije (TES), sredstvo za shranjevanje toplote, materiali s fazno premeno (PCM),
organski materiali, anorganski materiali, smiselno shranjevanje toplote, shranjevanje latentne toplote
1 INTRODUCTION
Thermal-energy storage systems have a wide variety
of applications1. Heat (cold) can be stored by heating
(cooling), melting (solidifying), vaporizing (liquefying)
a medium or by reversible thermochemical reactions.
Heat-storage media that undergo a phase change during
the process of storage and release of energy are called
phase-change materials (PCMs)2. The thermal-storage
capacity of PCMs depends on the specific heat in each
state and the latent heat of each phase transformation3. A
large heat of fusion and the transition temperature in a
required range are the two main characteristics deter-
mining the suitability of PCMs for a specific application.
The determination of the selected properties of PCMs is
the most important condition for a correct design of a
new application of PCMs in buildings and for a pre-
diction of the influence of the latent-heat storage on an
indoor environment and energy savings. In practice there
is a lack of reliable information about PCM properties
and, therefore, only the results of validated laboratory
experiments involving selected PCMs can help the
investigators in designing and developing latent-heat
storage systems. However, there are some limitations in
the use of PCMs4:
• PCMs may interact with the building structure and
change the properties of the building materials;
Materiali in tehnologije / Materials and technology 47 (2013) 2, 185–188 185
UDK 536.7:536.65 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)185(2013)
• there is a risk of a leakage of PCMs from the building
structure;
• PCMs have a rather poor thermal conductivity in the
solid state.
These problems are commonly solved with a proper
PCM encapsulation.
Salt hydrates, paraffin waxes, fatty acids and
eutectics of organic and non-organic compounds are the
main categories of PCMs that have been considered for
building use during the recent decades.
2 MATERIAL AND METHODS
Only the solid-liquid phase change of a material can
be used when a material is integrated in a building
structure. In some cases differential scanning calorimetry
(DSC) is a standard method for a thermal analysis of
PCMs. The most widely used scanning mode includes
heating and cooling at a constant rate5. This dynamic
method is used for the investigation of the melting and
solidification enthalpies. Figure 1 shows the solidi-
fication process of PCMs. The evolution of the released
heat flux is the function of the external temperature Text,
when this temperature is following a ramp6. The shape of
the curve depends on the temperature rates. The latent
heat of the phase change is calculated from the area
under the curve and the external temperature rate that is
constant in our case (e.g., 0.1 K min–1, 1 K min–1, 10
K min–1):
l Q t dt Q T
t
T v
Af
t
t
ext
extT
T
Text
f
e
= = =∫ ∫( ) ( )
1
2 0 1∂
∂
(1)
where lf is the latent heat of the phase transformation
(J), Q is the heat flux (W), t is the time (s), To is the
onset temperature (K), Te is the end temperature (K),
Text is the external temperature (K), vText is the external-
temperature rate (K s–1) and Af is the area under the
curve and the external-temperature rate (W K–1).
There are inflections of the plotted curve on each side
of the heat-flow peak temperature Tp. The onset tempe-
rature To and the end temperature Te are the temperatures
corresponding with the intersections between the
tangents at the inflection points and the base line. The
onset temperature and the peak temperature are often
used for the characterization of PCMs.
We have used selected non-commercial and commer-
cial organic- and inorganic-based PCMs in our experi-
ments. The list of selected PCMs is in Table 1. Perkin
Elmer PYRIS1 DSC, equipped with a cooling device
Perkin Elmer Intracooler 2P, was used for determining
the thermal properties (the heat of fusion and the melting
range). All DSC experiments were carried out at the
temperature rate of 0.1 K min–1.
A thermogravimetric apparatus (TGA) Q500 made
by TA Instruments was used for the evaluation of the
thermal stability of PCMs. The airflow rate was set to be
60 ml min–1 and the heating rate was 5 K min–1 from the
room temperature to 600 °C. An open platinum pan was
used as a sample holder. The weight of the samples was
approximately 10 mg. The results of TGA determine the
suitability of these materials in latent-heat-storage appli-
cations because the operating temperature must be below
the thermal-decomposition temperatures of PCMs. The
proposed operating-temperature range for building appli-
cation was estimated to be between 18 °C and 30 °C.7
Table 1: PCMs tested in laboratory experiments
Tabela 1: PCM, preizku{eni v laboratorijskih preizkusih
Sample Organic / inorganic Source
CaCl2·6H2O inorganic noncommercial
Parafol 16-97 organic Sasol
Parafol 18-97 organic Sasol
SP 22 A17 inorganic Rubitherm
SP 25 A8 inorganic Rubitherm
RT 21 organic Rubitherm
RT 27 organic Rubitherm
ThermusolHD26 inorganic Salca
3 RESULTS AND DISCUSSION
Characteristics of all the samples were tested twice.
The results in Tables 2 and 3 represent the average
values from both measurements. As already mentioned,
all the experiments were carried out at the temperature
rate of 0.1 K min–1. Though the experiments carried out
M. OSTRÝ et al.: CHARACTERIZATION OF SELECTED PHASE-CHANGE MATERIALS ...
186 Materiali in tehnologije / Materials and technology 47 (2013) 2, 185–188
Figure 1: Characteristic temperatures for the solidification process
Slika 1: Zna~ilne temperature pri procesu strjevanja
Table 2: Peak temperatures of selected PCMs
Tabela 2: Vrhovi temperature izbranih PCM
Sample
Peak temperature in °C
Melting Solidification
CaCl2·6H2O 29.9 –
Parafol 16-97 18.8 16.1
Parafol 18-97 28.9 27.3
SP 22 A17 22.5 22.4
SP 25 A8 26.6 18.5
RT 21 22.8 22.6
RT 27 27.8 27.6
ThermusolHD26 27.0 21.5
at the temperature rate of 0.1 K min–1 take roughly 10
times more time than the experiments at the rate of 1
K min–1, the slower rate was chosen because it is much
closer to the real daily swing of indoor air temperature in
summer in the rooms with natural ventilation without
air-conditioning. PCM-based heat storage integrated in
building structures is a way of controlling the indoor air
temperature by storing and releasing the thermal energy
from the solar radiation or internal heat gains.
Calcium chloride hexahydrate is a non-commercial
salt-hydrate PCM. Parafol 16-97 is based on hexadecane,
Parafol 18-97 is based on octadecane. Samples SP
22A17 and SP 25 A8 consist of a composition of salt
hydrates and organic compounds. Samples RT 21 and RT
27 are based on n-paraffins and waxes. Thermusol HD26
represents a commercial group of salt-hydrate-based
PCMs.
Table 3: Heat of fusion of tested PCMs
Tabela 3: Talilna toplota preizku{enih PCM
Sample
Heat of fusion in J g–1
Melting Solidification
CaCl2–6H2O 129.0 –
Parafol 16-97 223.6 –227.4
Parafol 18-97 221.3 –215.5
SP 22 A17 12.3 –11.9
SP 25 A8 71.4 –78.0
RT 21 116.7 –106.7
RT 27 139.7 –139.3
ThermusolHD26 132.4 –132.0
The possibility of a regeneration of PCMs (a
rejection of stored heat) at night is very important for the
building applications. The PCMs integrated with
building structures absorb heat gains during the day and
release the absorbed heat at night. If the heat absorbed
during one day is not released at night the ability of
PCMs to absorb heat the next day is reduced leading to a
limited contribution to the room-temperature control.
Two systems for the rejection of stored heat were studied
at the Brno University of Technology in the past:
• a natural or mechanical ventilation of the indoor
space;
• a circuit of cooled air or water integrated with the
structure containing a PCM.
Only the PCMs with suitable melting- and solidifi-
cation-temperature ranges can be used with each of the
systems. As can be seen in Table 2 calcium chloride
hexahydrate is suitable only for the naturally ventilated
spaces. The indoor temperature in the residential
buildings and in the offices must be maintained between
20 °C and 26 °C. But a serious disadvantage of this
material is its tendency to supercool during the solidi-
fication process. This kind of PCM cannot be used
without a modification that reduces the supercooling
effect. There are no results from the solidification
process just because of the supercooling. On the other
hand, Parafol 18-97, RT 27 and Thermusol HD26 could
be used for the systems with ventilation of the interior.
PCMs absorb cooling loads and release energy in the
temperature range between 21 °C and 28 °C. This fact
allows for cooling down the indoor environment only to
20 °C to reject the absorbed heat (the regeneration of
PCMs). But these systems cannot commonly guarantee
thermal comfort in the rooms during very hot summer
days.
On the other hand, the temperature of cooled water or
air in a separate circuit integrated with the building
structures containing PCMs can be kept bellow 20 °C
without a negative impact on the thermal comfort of the
occupants. This fact allows for the use of the PCMs with
a lower solidification range (e.g., SP 22A5 from the
tested group).
As can be seen in Figure 2 the samples of Parafol
16-97 and 18-97 have a very narrow range of melting
and solidification temperatures, about 0.3 °C and 0.4 °C.
This is an advantage for the short-term storage systems
that represent building structures.
M. OSTRÝ et al.: CHARACTERIZATION OF SELECTED PHASE-CHANGE MATERIALS ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 185–188 187
Figure 3: Results of DSC for SP 25A8
Slika 3: Rezultati DSC za SP 25A8
Figure 2: Results of DSC for Parafol 16-97 and Parafol 18-97
Slika 2: Rezultati DSC za Parafol 16-97 in Parafol 18-97
Compared to the paraffin-based PCMs, the PCMs
that are a mixture of salt hydrates and organic com-
pounds, tested in our experiment (Figure 3), have a wide
range of melting temperatures and a rather narrow
temperature range of solidification. This could be an
advantage for the systems with a cooled-water loop,
because the thermal energy can be slowly stored in a
PCM during the day and quickly discharged at night by
the cooled-water circuit.
All the materials tested in our experiments are
suitable for building application from the point of view
of thermal decomposition. The thermal decomposition of
all the tested PCMs begins above the expected operation
temperatures. The result for the sample composition of
salt hydrates and organic compounds is shown in Figure
4.
The difficulties may occur with the use of salt-
hydrate-based PCMs because of the changes in the water
content. This effect was observed during the TGA at low
temperatures. Therefore, salt-hydrate-based PCMs must
be very tightly sealed in the containers.
4 CONCLUSION
A series of laboratory experiments was carried out to
assess the suitability of the selected phase-change
materials for the use in built environments. The purpose
of a PCM-based building-integrated thermal storage is to
contribute to the thermal stability or temperature control
in buildings. All the tested materials were found suitable
for this purpose from the point of view of thermal
decomposition. That was due to rather low operating
temperatures (mostly lower than 30 °C in buildings). All
the tested materials exhibited the melting ranges that are
suitable for building applications. However, the suitabi-
lity of PCMs for integration with building structures
from the point view of melting ranges and enthalpies of
fusion depends, particularly, on the type of integrating
and the type of rejecting stored heat.
Acknowledgement
This work was supported by the Czech Grant Agency
under the project No. P104/12/1838 "Utilization of latent
heat storage in phase change materials to reduce primary
energy consumption in buildings".
5 REFERENCES
1 A. Sharma, V. V. Tyagi, C. R. Chen, D. Buddhi, Review on thermal
energy storage with phase change materials and applications, Renew-
able and Sustainable Energy Reviews, 13 (2009) 2, 318–345
2 I. Dinçer, M. A. Rosen, Thermal Energy Storage: Systems and
Applications, Chichester: John Wiley & Sons, Ltd., 2002, 579
3 E. Günther, S. Hiebler, H. Mehling, R. Redlich, International Journal
of Thermophysics, 30 (2009) 4, 1572–9547
4 C. Y. Zhao, G. H. Zhang, Review on microencapsulated phase
change materials (MEPCMS): Fabrication, characterization and
applications, Renewable and Sustainable Energy Reviews, 15 (2011)
8, 3813–3832
5 C. Castellón, E. Günther, H. Mehling, S. Hiebler, L. F. Cabeza,
Determination of the enthalpy of PCM as a function of temperature
using a heat-flux DSC – A study of different measurement proce-
dures and their accuracy, International Journal of Energy Research,
32 (2008) 13, 1258–1265
6 F. Kuznik, D. David, K. Johannes, J. J. Roux, A review on phase
change materials integrated in building walls, Renewable and
Sustainable Energy Reviews, 15 (2011) 1, 379–391
7 J. Skramlik, M. Novotny, K. Suhajda, Modeling of diffusion in
porous medium, International Conference on Numerical Analysis
and Applied Mathematics ICNAAM 2011, Halkidiki, American
Institute of Physics, 2011
M. OSTRÝ et al.: CHARACTERIZATION OF SELECTED PHASE-CHANGE MATERIALS ...
188 Materiali in tehnologije / Materials and technology 47 (2013) 2, 185–188
Figure 4: Example of results of a TGA analysis
Slika 4: Primer rezultatov TGA-analize
R. KOTTNER et al.: IMPROVEMENT OF THE DAMPING PROPERTIES OF CARBON-FIBRE-REINFORCED ...
IMPROVEMENT OF THE DAMPING PROPERTIES OF
CARBON-FIBRE-REINFORCED LAMINATED PLASTICS
USING DAMPING LAYERS
IZBOLJ[ANJE DU[ENJA Z OGLJIKOVIMI VLAKNI OJA^ANE
LAMINIRANE PLASTIKE Z UPORABO PLASTI ZA DU[ENJE
Radek Kottner1, Josef Vacík2, Robert Zem~ík3
1University of West Bohemia, European Centre of Excellence, New Technologies for Information Society, Univerzitní 22, 30614 Plzeò,
Czech Republic
2University of West Bohemia, Department of Machine Design, Univerzitní 22, 30614 Plzeò, Czech Republic
3University of West Bohemia, Department of Mechanics, Univerzitní 22, 30614 Plzeò, Czech Republic
kottner@kme.zcu.cz
Prejem rokopisa – received: 2012-09-03; sprejem za objavo – accepted for publication: 2012-10-08
A suitable hybrid composite consisting of carbon-fibre-reinforced plastic and damping layers was investigated in terms of
damping and natural frequencies using experiments and numerical simulations. The frequency response and the transient
response of cantilever beams were analysed. The damping layers made from rubber or from a cork-rubber composite material
were used in the investigated hybrid structure. A laser-measurement device and an accelerometer were used for the measurement
of the responses. Pareto optimization was performed using three-dimensional numerical simulations with the aim to maximize
the fundamental natural frequency and the damping ratio.
Keywords: hybrid, composite, carbon-fibre-reinforced plastic, rubber, damping
Sestavljeni kompozit iz plastike, oja~ane z ogljikovimi vlakni in s plastmi za du{enje naravnih frekvenc, je bil preiskovan
eksperimentalno in z numeri~no simulacijo. Analiziran je bil frekven~ni odgovor in prehodni odziv konzolnega nosilca. V
preiskovani hibridni strukturi je bila plast za du{enje izdelana iz gume ali sestavljena iz gume in plute. Za meritve odgovora je
bil uporabljen laserski merilnik in merilnik pospe{ka. Z uporabo tridimenzionalne numeri~ne simulacije je bila izvr{ena
Paretova optimizacija za maksimiranje osnovne naravne frekvence in koli~nika du{enja.
Klju~ne besede: hibrid, kompozit, z ogljikovimi vlakni oja~ana plastika, guma, du{enje
1 INTRODUCTION
Recently, the conventional metallic structures have
been replaced with composite structures thanks to their
superior dynamic characteristics. For example, a higher
specific stiffness of carbon-fibre-reinforced plastics
(CFRPs) allows higher natural frequencies of the structu-
res. Furthermore, the damping characteristics of CFRPs
have higher values and, moreover, can be improved using
an integration of the damping layers1.
The aim of this work was to develop a suitable hybrid
structure in terms of damping and natural frequencies.
The investigated samples were the cantilever beams,
whose damping layers were made from rubber or from
the ACM87 cork-rubber composite material. The damp-
ing ratio for their fundamental natural frequency was
analyzed.
2 THEORY
The equation of the motion of a damped system is2:
Mq Cq Kq f  ( )+ + = t (1)
where q[m], q[m s–1], q[m s–2] are the vectors of the ge-
neralized coordinates and their 1st and 2nd differenti-
ations with respect to time t; f(t)[N] is the vector of the
generalized time-dependent applied force; M[kg] is the
mass matrix of the system; C[N s m–1] is the damping
matrix and K[N m–1] is the stiffness matrix. The motion
of the discrete linear systems with a single degree of
freedom can be described as:
mq cq kq F  ( )+ + = t (2)
In the case of free oscillations, Equation (2) can be
consequently rewritten as:
 q q q+ + =2 00 
 (3)
Materiali in tehnologije / Materials and technology 47 (2013) 2, 189–193 189
UDK 519.61/.64:66.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)189(2013)
Figure 1: Response of an underdamped harmonic oscillator
Slika 1: Odgovor harmoni~nega nedu{enja
where the damping ratio  and the undamped natural
frequency 0 are defined as:
 
c
mk2
(4)
 
k
m
(5)
In the case of an underdamped system (0 <  < 1) the
solution of Equation (3) representing the displacement of
the system can be found in the following form:
q t Ce tt    
− +sin( ) (6)
where C[m] is the amplitude, [rad s–1] is the damped
natural frequency of the system and 0[rad] is the phase
shift.
The damped natural frequency can be expressed as:
  
1
2
2− =
π
π
T
f (7)
where T[s] is the period of the waveform (Figure 1) and
f is the damped natural frequency in Hertz.
Based on Equation (6), the exponential attenuation
rate is then defined as:
b = 0 (8)
The damping ratio can be determined using the loga-
rithmic decrement , which is defined as the natural
logarithm of any two peaks:
 
	


 
= = = = =
−
−1 2
1
0
n
q
q
e T bT
n
Tln ln (9)
where q0 is the greater of the two amplitudes and qn is
the amplitude of a peak n periods away. The damping
ratio is then found from the logarithmic decrement:


	  
=
−4
(10)
According to Rayleigh damping, the damping matrix
C is given by:
C = M + K (11)
where  and  are the Rayleigh constants. Assuming 
= 0 and the damping ratio  of both the CFRP and the
damping materials is constant, the Rayleigh constant
can be expressed as:



= (12)
When the difference between the undamped natural
frequency 0 and the damped natural frequency  is
negligible, the Rayleigh constants of the hybrid compo-
site components can be determined as:


i
i
h
=
0 ,
(13)
where i is the damping ratio of the hybrid composite
component (CFRP or the damping material) and 0,h is
the undamped natural frequency of the hybrid composite
determined from the modal analysis.
3 EXPERIMENTS
Two types of CFRP and two types of damping layers
were used in the experiments. A linear behaviour of all
the materials was assumed. The mechanical properties of
the materials are listed in Tables 1 to 4.3 A modal analy-
sis was used for an identification of the elastic-material
properties4.
Two cantilever flat bars consisting of unidirectional
913C-HTS CFRP and rubber were investigated. Sample
A had a thickness of 7.8 mm (Figure 2a), sample B had
a thickness of 12.9 mm (Figure 2b). The thickness of
913C-HTS CFRP was 2.7 mm and the thickness of
rubber was 2.4 mm. The width of the bars was 19.8 mm,
the length was 450 mm and the length of the clamping
was 38 mm. The harmonic response after the initial
deflection was investigated using an optoNCDT
R. KOTTNER et al.: IMPROVEMENT OF THE DAMPING PROPERTIES OF CARBON-FIBRE-REINFORCED ...
190 Materiali in tehnologije / Materials and technology 47 (2013) 2, 189–193
Figure 2: Cantilever flat bar consisting of 913C-HTS CFRP and
rubber: a) sample A, b) sample B
Slika 2: Konzolna plo{~ata palica iz 913C-HTS CFRP in gume: a)
vzorec A, b) vzorec B
laser-measurement device or a Brüel & Kjaer 4507
accelerometer.
Table 1: Mechanical properties of 913C-HTS CFRP
Tabela 1: Mehanske lastnosti 913C-HTS CFRP
Longitudinal modulus E1 GPa 104
Transverse modulus E2 GPa 5.5
Shear modulus G12 GPa 2.4
Poisson’s ratio 12 – 0.34
Density  kg/m3 1.500
Damping ratio  – 0.002
Table 2: Mechanical properties of K63712 CFRP
Tabela 2: Mehanske lastnosti K63712 CFRP
Longitudinal modulus E1 GPa 280
Transverse modulus E2 GPa 3.5
Shear modulus G12 GPa 1.7
Poisson’s ratio 12 – 0.38
Density  kg/m3 1.470
Damping ratio  – 0.003
Table 3: Mechanical properties of rubber
Tabela 3: Mehanske lastnosti gume
Young’s modulus E MPa 10
Poisson’s ratio  – 0.49
Density  kg/m3 1 170
Damping ratio  – 0.072
Table 4: Mechanical properties of the ACM87 composite
Tabela 4: Mehanske lastnosti kompozita ACM87
Young’s modulus E MPa 2.5
Poisson’s ratio  – 0.3
Density  kg/m3 740
Damping ratio  – 0.112
Table 5: Results of the experiments and models
Tabela 5: Rezultati eksperimentov in modelov
Damped natural
frequency f /Hz
Damping ratio

Exp. Model Exp. Model
Sample A 35.5 37.8 0.034 0.031
Sample B 44.3 43.8 0.043 0.041
Sample C – clamped 44.0 43.7 0.039 0.004
Sample C – free 548 549 0.005 0.004
R. KOTTNER et al.: IMPROVEMENT OF THE DAMPING PROPERTIES OF CARBON-FIBRE-REINFORCED ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 189–193 191
Figure 5: Model of a hybrid square tube
Slika 5: Model hibridne kvadratne cevi
Figure 4: Boundary conditions of a model of a cantilever flat bar
Slika 4: Robni pogoji modela konzolne plo{~ate palice
Figure 3: Cantilever square tube consisting of K63712 CFRP and the
ACM87 cork-rubber composite material: sample C
Slika 3: Konzolna kvadratna cev iz K63712 CFRP in ACM87 kom-
pozitnega materiala pluta-guma; vzorec C
Further, the response of the cantilever square tube
consisting of the wound K63712 CFRP and the ACM87
composite with a thickness of 6.5 mm (Figure 3) was
investigated (Sample C). The thickness of K63712 CFRP
was 2.4 resp. 2.1 mm and the thickness of the ACM87
composite was 2 mm. The width of the square tube was
103 mm, the tube length was 1 490 mm and the length of
the clamping was 135 mm. However, the clamping was
not sufficiently rigid. Therefore, the damping ratio was
investigated also in the case of a free square tube (a tube
hung on a rope).
The results are listed in Table 5.
4 NUMERICAL SIMULATIONS
Three-dimensional finite-element models were
created in MSC.Marc using eight-node brick elements
(with an assumed strain option) as shown in Figures 4
and 5. The CFRP materials were considered orthotropic
and homogenous. The damping materials were consi-
dered isotropic and homogenous. The boundary condi-
tions of the models of the cantilever beams are obvious
from Figure 4. The undamped natural frequencies were
obtained with a modal analysis. After an evaluation of
the Rayleigh constant using Equation (12), the damped
natural frequencies and the damping ratio were obtained
with a transient analysis. The Newmark time integration
scheme was used.
The created models were validated by comparing the
experimental and numerical results listed in Table 5. The
difference between the results for samples A and B was
less than 9 %. In the case of the clamped sample C, the
clamping was not sufficiently rigid as already mentioned
above. Due to this fact, the damping ratio had a signifi-
cantly higher value, which was confirmed with the
experiment with the free sample C. The difference
between the results involving the free sample C was less
than 20 %.
A Pareto optimization was performed for the canti-
lever flat bar with the aim to maximize the fundamental
natural frequency and the damping ratio. The analyzed
bar was a symmetrical hybrid laminate with 14 layers
consisting of 913C-HTS CFRP and rubber. The thick-
ness of the layers was 1 mm and the fibre orientation of
the CFRP layers was identical to the beam axis. The
width of the bar was 20 mm, the length 450 mm and the
length of the clamping was 30 mm. The Pareto front is
shown in Figure 6, a detail of the front in Figure 7. It is
obvious from both figures that the situation is very
complex; therefore, the problem should be more pre-
cisely constrained.
5 CONCLUSION
The performed experiments and numerical simu-
lations showed that the investigated natural frequencies
and damping ratios of the hybrid composite structures
R. KOTTNER et al.: IMPROVEMENT OF THE DAMPING PROPERTIES OF CARBON-FIBRE-REINFORCED ...
192 Materiali in tehnologije / Materials and technology 47 (2013) 2, 189–193
Figure 7: Detail of the Pareto front
Slika 7: Detajl Paretove linije
Figure 6: Pareto optimization results, "C" means CFRP, "D" means
the damping layer
Slika 6: Rezultati Paretove optimizacije: "C" pomeni CFRP, "D" po-
meni plast za du{enje
strongly depend on the placement of the damping layers
in the hybrid cross-sections. The accurate place of the
damping layers must be investigated for each application
in dependence on the material of the layers, the geometry
and the requested ratio between the fundamental natural
frequency and the damping ratio.
Acknowledgement
This work was supported by the European Regional
Development Fund (ERDF), within the project "NTIS –
New Technologies for Information Society", European
Centre of Excellence, CZ.1.05/1.1.00/02.0090, and by
the research project no. P101/11/0288.
6 REFERENCES
1 H. Y. Hwang, H. G. Lee, D. G. Lee, Clamping Effects on the Dyna-
mic Characteristics of Composite Machine Tool Structures, Com-
posite Structures, 66 (2004), 399–409
2 V. Zeman, Z. Hlavá~, Kmitání mechanických soustav, Z^U Plzeò,
2004
3 J. Vacík, V. La{ová, R. Kottner, J. Káòa, Experimental Determination
of Damping Characteristics of Hybrid Composite Structure, In:
Experimental Stress Analysis, Velké Losiny, Czech Republic, 2010
4 R. Zem~ík, R. Kottner, V. La{, T. Plundrich, Identification of Mate-
rial Properties of Quasi-unidirectional Carbon-epoxy Composite
Using Modal Analysis, Mater. Tehnol., 43 (2009) 5, 257–260
R. KOTTNER et al.: IMPROVEMENT OF THE DAMPING PROPERTIES OF CARBON-FIBRE-REINFORCED ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 189–193 193

S. KASTELIC et al.: AA413.0 AND AA1050 JOINED WITH FRICTION-STIR WELDING
AA413.0 AND AA1050 JOINED WITH FRICTION-STIR
WELDING
SPAJANJE ZLITINE AA413.0 IN AA1050 Z GNETENJEM
Sebastjan Kastelic1,3, Janez Tu{ek2, Damjan Klob~ar2, Jo`ef Medved3, Primo` Mrvar3
1Institut for Foundry and Heat Treatment, Litostrojska cesta 60, 1000 Ljubljana, Slovenia
2University of Ljubljana, Faculty of Mechanical Engineering, A{ker~eva 6, 1000 Ljubljana, Slovenia
3University of Ljubljana, Faculty of Natural Sciences and Engineering, A{ker~eva 12, 1000 Ljubljana, Slovenia
kastelic.sebastjan@gmail.com
Prejem rokopisa – received: 2012-09-03; sprejem za objavo – accepted for publication: 2012-10-09
Friction-stir-welding (FSW) technology has been growing since it was patented in 1991 at TWI. Since then the majority of
research and industrial applications for joining aluminium alloys were made on wrought aluminium alloys. Lately several
investigations have been done in FSW of dissimilar alloys. FSW also has a big potential in the casting industry – especially in
high-pressure die casting (HPDC). In this article an investigation of a FSW dissimilar joint made from a casting aluminium
alloy (AA413.0) and technically pure aluminium (AA1050) was done. This kind of joint can be used to make an assembled
casting, joined with FSW with the aim to have a casting with different material properties or to join HPDC with FSW to
assemble a casting with inner cavities.
In this article the temperature distribution of the FSW joint of a cast aluminum alloy and technically pure aluminum is
investigated. In the experimental work several FSW parameters were tested: the tool speed, the tool rotation and the position of
the tool regarding the joint center. During joining the temperature was measured with a thermocouple and the temperature
distribution in steady state was calculated with the FEM program Sysweld. The microstructure and mechanical properties of the
joint were investigated.
Keywords: friction-stir welding, AA413.0, AA1050, finite-element method
Spajanje z gnetenjem je tehnologija, ki se intenzivno razvija od patentiranja leta 1991 na Britanskem in{titutu za varjenje. V
industrijski praksi je najbolj raz{irjena uporaba FSW-tehnologije za spajanje gnetnih aluminijevih zlitin. V zadnjem ~asu se
raziskave osredinjajo tudi na spajanje dveh razli~nih zlitin. Spajanje z gnetenjem ima velik potencial tudi v livarski industriji, {e
posebej pri tla~nem litju. Ta ~lanek opisuje izvedeno eksperimentalno spajanje dveh razli~nih zlitin. Spojeni sta bili gnetna
zlitina AA1050 in livarska zlitina AA413.0. Taki spoji imajo svojo uporabno vrednost pri sestavljenih ulitkih, kjer lahko s
spajanjem razli~nih zlitin z razli~nimi fizikalnimi lastnostmi dose`emo nehomogene lastnosti sklopa, kjer je to potrebno. S
spajanjem tla~nih ulitih delov lahko izdelamo sklope z notranjimi votlimi deli. Lastnosti tako izdelanih sklopov so primerljive z
lastnostmi tla~nih ulitkov.
V ~lanku so navedene temperature, izmerjene med spajanjem dveh razli~nih zlitin. Za dolo~itev optimalnih parametrov spajanja
je bilo le-to izvedeno pri razli~nih vrtljajih orodja in pri razli~nih hitrostih spajanja. Spoji so bili mehansko analizirani.
Stacionarno temperaturno polje med spajanjem dveh razli~nih zlitin je bilo izra~unano s programom Sysweld z metodo kon~nih
elementov.
Klju~ne besede: spajanje z gnetenjem, AA413.0, AA1050, metoda kon~nih elementov
1 INTODUCTION
Aluminum die casting alloys are made with a rapid
injection of a molten metal into metal molds under high
pressure. Such an alloy has a dense and fine grain
surface, resulting in excellent wear and fatigue pro-
perties. Approximately 85 % of aluminum die casting
alloys are based on Al-Si-Cu. These alloys provide a
good combination of the cost, strength, and corrosion
resistance, together with high fluidities that are required
for easy casting. In recent years, aluminum die casting
alloys have been widely used in the automotive,
electronics, machine and building industries because
they are light and recyclable. However, these castings
have their limitations arising from the casting-process
limitations. This problem can be solved by joining
several cast parts into one complex product. The welding
of aluminum and its alloys has always represented a
great challenge for designers and technologists. Fusion
welding of aluminum die casting alloys is difficult due to
the formation of welding defects, such as blowholes, and
the welding deformation as a result of a high coefficient
of thermal expansion of aluminum alloys. As welding
defects result in decreased mechanical properties, this
problem must be solved for the use in practical appli-
cations.1,2
From this point of view, friction-stir welding (FSW)
was developed as a new joining process by The Welding
Institute (TWI) in 1991.3 The basic concept of FSW is
remarkably simple. A non-consumable rotating tool with
a specially designed pin and shoulder is inserted into the
matching edges of the sheets or plates to be joined and
traversed along the line of the joint (Figure 1). The tool
serves two primary functions, the heating of a workpiece,
and the movement of a material to produce the joint. The
heating is accomplished by the friction between the tool
and the workpiece, creating a plastic deformation of the
workpiece. The localized heating softens the material
around the pin and a combination of the tool rotation and
translation leads to a movement of the material from the
Materiali in tehnologije / Materials and technology 47 (2013) 2, 195–198 195
UDK 621.792.3:669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)195(2013)
front of the pin to the back of the pin. As a result of this
process a joint is produced in solid state. Because of the
various geometrical features of the tool, the material
movement around the pin can be quite complex. During
a FSW process, the material undergoes an intense plastic
deformation at an elevated temperature, resulting in a
generation of fine and equiaxed recrystallized grains.
The fine microstructure of the friction-stir welds pro-
duces good mechanical properties.4,5
FSW is effective in joining a number of different
materials. Until now it has been aluminum alloys that
FSW has been most successfully applied to. The reason
for this is a combination of the process simplicity and a
wide use of aluminum alloys in the mayor industries.
FSW can be used for joining aluminum alloys that are
difficult to fusion weld. FSW dominates in the fabri-
cation of aluminum components and panels. Now even
friction-stir spot welding (FSSW) is being intensively
studied.6,7
Only a limited number of FSW-joint studies have
been done on cast aluminum alloys, although they have
been intensively studied.4,8
2 EXPERIMENTAL WORK
In this work an investigation of the temperature field
of a FSW joint made from two different aluminum alloys
has been done. The first alloy is a common cast alloy
AA413.0 (AlSi12Cu), the other is made from pure
aluminum AA1050 (Al 99.5 %). Such a joint is very
interesting because it is made of alloys with different
properties. The chemical compositions of AA1050 and
AA413.0 are presented in Tables 19 and 29.
In the experimental work two plates were joined. The
joint was made on a milling machine Prvomajska ALG
2008. For this experiment the optimal feeding speed was
235 mm/min and the rotational speed of the tool was 475
min–1. The tool was inclined at an angle of 3° as pre-
sented in Figure 2. The tool presented in Figure 3 was
made of hot-work tool steel H13. The dimensions of the
used tool can be seen in Figure 4. The AA1050 alloy
was on the retreating side and the AA413.0 alloy was on
the advancing side. These parameters were selected on
the basis of several tests involving the selected tool, plate
thickness and milling-machine gear ratio.
The dimensions of the plates were 380 mm × 60 mm
× 6 mm. To each plate a K-type thermocouple was
mounted. Thermocouples were connected with the
National Instrument CompactDAQ NI 9213 analog
digital converter which was connected to the Labview
software. The sampling rate was 10 Hz.
During FSW the model temperature distribution was
numerically calculated with the Sysweld finite-element
modeling software. The FSW module in Sysweld
S. KASTELIC et al.: AA413.0 AND AA1050 JOINED WITH FRICTION-STIR WELDING
196 Materiali in tehnologije / Materials and technology 47 (2013) 2, 195–198
Table 1: Chemical composition of AA1050 and AA413.0 in mass fractions, w/%9
Tabela 1: Kemijska sestava zlitine AA1050 in AA413.0 v masnih dele`ih, w/%9
Al Si Fe Cu Mn Mg Zn V
AA1050 99.5 max 0.25 max 0.4 max 0.05 max 0.05 max 0.05 max 0.05 max 0.05
AA413.0 rest 11–13 1.3 1.0 0.35 0.1 0.5
Table 2: AA1050 and AA413.0 properties9
Tabela 2: Fizikalne lastnosti zlitine AA1050 in AA413.09
Liquidus temp. Density Tensile strength Yield strength Specific heat Thermalconductivity
Electrical
resistivity
°C g/cm3 MPa MPa J/(kg K) W/(m K) n m
AA1050 657 2.705 76 28 900 234 27.9
AA413.0 575–585 2.657 290 130 963 121 55.6
Figure 2: Tool tilt
Slika 2: Nagib orodja
Figure 1: Schematic drawing of friction-stir welding4
Slika 1: Shematski prikaz spajanja z gnetenjem4
enables a calculation of the steady-state temperature
field based on the custom FSW joining parameters for
joining the parts of the same alloy10. For this experiment
the model was upgraded so that the temperature field
was calculated on the basis of joining two different
alloys. The investigated model had 19277 nodes and the
CPU time was two hours.
3 RESULTS AND DISCUSSION
The joint plate is shown in Figure 5. Some of the
material was extruded during the joining because the
plates were not in perfect alignment. The alignment of
the plates had a big influence on the joint quality during
the testing of the FSW joining parameters; in addition,
the slot on the contact surface between the plates must be
minimum.
The temperatures measured with the thermocouples
and numerically calculated are presented in Figure 6.
The maximum temperature measured on the advancing
side was 250 °C and on the retreading side of the joint it
was 240 °C. Similar temperatures at the thermocouple
places were calculated with the FEM software. The
maximum calculated temperature in the joint was 470
°C. This temperature and calculated stationary tempera-
ture field are presented in Figure 7. In the upper part of
S. KASTELIC et al.: AA413.0 AND AA1050 JOINED WITH FRICTION-STIR WELDING
Materiali in tehnologije / Materials and technology 47 (2013) 2, 195–198 197
Figure 7: Temperature distribution calculated with the FSW module
Slika 7: Razporeditev temperature, izra~unane s FSW-modulom
Figure 4: Tool dimensions
Slika 4: Dimenzije orodja
Figure 5: Plates joint with FSW (upper plate of AA413.0 alloy, lower
plate of AA1050 alloy)
Slika 5: Plo{~i, spojeni s FSW-spojem (zgoraj zlitina AA413.0, spodaj
zlitina AA1050)
Figure 3: Tool used in the experiment
Slika 3: Orodje, uporabljeno za eksperiment
Figure 6: Measured and calculated temperatures in the joint
Slika 6: Izmerjene in izra~unane temperature v spoju
Figure 8: Tensile tests of the alloys and the FSW joint
Slika 8: Natezni preizkusi zlitin in FSW-spoja
Figure 7 the temperature field of the joint cross-section
can be seen. These measurements show that FSW is a
process joining different parts without melting the
material. It can be seen that the heat input into the joint
is very low. Due to a proper support and good clamping
of the plates during FSW, no deformation was found
after unclamping.
For a mechanical investigation of the joint, tensile
tests were done. The results of the tensile tests of the
joined alloys and the FSW joint can be seen in Figure 8.
A contraction and defect of the joint appeared on the
AA1050 alloy. The ultimate tensile stress of the joint
was 74 MPa and it was the same as the ultimate tensile
stress of the AA1050 alloy. Through the cross-section of
the joint the Vickers hardness was measured. The hard-
ness in the AA1050 alloy was 35 HV, 80 HV in the joint
and 100 HV in the AA413.0 alloy. Good mechanical
properties of the joint can also be seen from the micro-
structure investigation of the joint presented in Figure 9.
According to standard ASTM E112 the mean intercept
distance in the area marked with 1 was 92.37 μm, in area
2 the mean intercept distance was 44.86 μm, in area 3 the
distance was 8.15 μm and in area 4 the distance was 5.56
μm. From the microstructure of the cross-section a good
mixing of the used alloys can be seen. The stirring of the
pin causes the grain size in the joint to be more than ten
times smaller than the grain size in the base AA1050
alloy.
4 CONCLUSION
FSW has a big potential in joining a casting with
other castings and/or extruded or rolled parts. In this
article it was shown that the heat input for joining two
different alloys with FSW is lower than for the joints
made with other conventional welding processes. With
the experiment we exactly determined the temperature in
the stirring zone and calculated the established tempe-
rature field during the FSW process. A lower heat input
leads to a small deformation of the workpiece and good
mechanical properties of the joint.
The mechanical properties of the joint are better than
those of the weaker alloy (AA1050) that was friction-stir
welded. The good mechanical properties of the joint can
be confirmed with the mean intercept distance between
the grains in the joint. In these experiments the optimum
FSW parameters for joining the AA413.0 and AA1050
alloys were used. Alloy AA413.0 must be on the
advancing side, the tool rotational speed should be 475
min–1, the joining speed should be 235 mm/min and the
tool angle is 3°.
5 REFERENCES
1 M. Ericsson, R. Sandstrom, International Journal of Fatigue, 25
(2003), 1379–1387
2 Y. G. Kim, H. Fujii, T. Tsumura, T. Komazaki, K. Nakata, Materials
Letters, 60 (2006), 3830–3837
3 W. M. Thomas, Friction Stir Butt Welding International Patent
Application, No. PCT/GB92 Patent Application No. 9125978.8,
1991
4 R. S. Mishra, Z. Y. Ma, Materials Science and Engineering R, 50
(2005), 1–78
5 S. R. Mishra, M. W. Mahoney, Friction Stir Welding and Processing,
ASM International, 2007, 7–35
6 S. Hirasawa, H. Badarinarayan, K. Okamoto, T. Tomimura, T.
Kawanami, Journal of Materials Processing Technology, 210 (2010),
1455–1463
7 M. K. Kulekci, U. Esme, O. Er, Mater. Tehnol., 45 (2011) 5, 395–399
8 Y. G. Kim, H. Fujii, T. Tsumura, T. Komazaki, K. Nakata, Materials
Letters, 60 (2006), 3830–3837
9 ASM International: Volume 2, Properties and selection: Nonferous
alloys and special-purpose materials, Metals Park, Ohio, 1990
10 E. Feulvarch, Y. Gooroochurn, F. Boitout, 3D Modelling of Thermo-
fluid Flow in Friction Stir Welding. In: Proceedings of the 7th inter-
national conference on trends in welding research, Callaway Gardens
resort, Pine Mountain, Georgia, USA, 2005
S. KASTELIC et al.: AA413.0 AND AA1050 JOINED WITH FRICTION-STIR WELDING
198 Materiali in tehnologije / Materials and technology 47 (2013) 2, 195–198
Figure 9: FSW cross-section in the polarolized light and grain size
Slika 9: Prerez FSW-spoja v polarizirani svetlobi in velikost zrn
M. SADÍLEK et al.: INCREASING TOOL LIFE DURING TURNING WITH A VARIABLE DEPTH OF CUT
INCREASING TOOL LIFE DURING TURNING WITH A
VARIABLE DEPTH OF CUT
POVE^ANJE ZDR@LJIVOSTI ORODJA PRI STRU@ENJU Z
VARIABILNO GLOBINO REZA
Marek Sadílek, Robert ^ep, Zuzana Sadílková, Jan Valí~ek, Lenka Petøkovská
V[B-Technical University of Ostrava, Faculty of Mechanical Engineering, 17. listopadu 15/2172, 708 33 Ostrava, Czech Republic
marek.sadilek@vsb.cz
Prejem rokopisa – received: 2012-09-05; sprejem za objavo – accepted for publication: 2012-09-27
The article deals with the improvement of cutting-tool durability by using CAD/CAM systems. It proposes new roughing
turning cycles where a variable depth of cut is applied. The experimental part verifies theoretical prerequisites when a flange is
being machined with a sintered-carbide cutting tool. It compares the turning where the standard roughing cycle is used and the
turning where the proposed roughing cycle with a variable depth of cut is applied.
Keywords: variable depth of cut, durability, CAD-CAM systems, turning
^lanek obravnava izbolj{anje zdr`ljivosti orodij za odrezavanje z uporabo sistema CAD/CAM. Predlagani so novi cikli grobega
stru`enja, kjer se uporablja spremenljivo globino rezanja. V eksperimentalnem delu so potrjeni teoreti~ni prvi pogoji pri
stru`enju prirobnice z rezilnim orodjem iz sintranih karbidov. Primerjano je stru`enje, kjer je uporabljen standarden cikel
odrezavanja, s tistim stru`enjem, kjer je uporabljen predlagan cikel odrezavanja s spremenljivo globino rezanja.
Klju~ne besede: spremenljiva globina rezanja, zdr`ljivost, sistem CAD-CAM, stru`enje
1 SUPPOSED PROCESS OF TOOL WEAR
Tool wear depends on numerous factors, for example:
workpiece material, tool material and geometry, cutting
parameters (cutting speed, feed, cutting depth), used
process liquid, cutting machine and many others1.
It is proved that cutting the depth does not have the
greatest influence. Since it has its share in the intensity
of tool wear (a growing cutting depth increases the tool
wear as well) it makes sense to deal with cutting depth
and we can consider it as a significant factor of tool
wear.
1.1 Tool wear of a selected type of cutting tools
The most used cutting materials in the area of CNC
machining are sintered carbide and cutting ceramics2,3.
As seen in Figure 1, the tool-wear method depends on
the applied tool materials. The tool wear of the cutting
ceramics is rather linear without any marked stepped
increases. The tool wear will therefore grow with an
increasing depth of cut. It is, thus, not advisable to apply
a variable depth of cut to the ceramics with such
behavior. The types of cutting ceramics, such as the
nitride ceramics, in which a more pronounced notch on
the back is formed during the machining may be used
exceptionally4,5.
Sintered carbide tends to form a pronounced notch on
the face and main back. This notch could be very advan-
tageously used in the roughing cycles with a variable
depth of cut. Efforts will be made to distribute this
pronounced notch over the maximum possible length of
the tool’s cutting edge.
Figure 2 below describes the used marking of the
tool wear according to ISO 3685.6
In applying a variable depth of cut (in the sintered
carbide) the durability-improvement effect is expected to
occur in the cutting edge provided the wear shaped as a
notch on the back is distributed over the longer part of
the cutting edge. These prerequisites were verified with
the experiments during the practical machining of the
concerned part – the flange.
Materiali in tehnologije / Materials and technology 47 (2013) 2, 199–203 199
UDK 621.941:004.896 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)199(2013)
1 – surface of wear on the back A – main back
2 – crater wear A´ – minor back
3 – notch wear on the main back A – face
4 – notch wear on the minor back
5 – notch on the face
Figure 1: Cutting-ceramics wear process and the tool-wear process of
sintered carbides4,5
Slika 1: Proces obrabe rezilne keramike in obrabe orodja iz sintranih
karbidov4,5
Both types of the cutting material were tested in
practice by experimental machining of the selected
turning parts. The theoretical assumption that the cutting
ceramic does not tend to form a notch on the back was
proven and it is essential for this tool-life improvement
method. That is why only the experimental works where
sintered carbide is used are stated below.
2 POSSIBILITIES OF A ROUGHING TOOL
PATH IN CAM SYSTEMS
The basic types of the used machining cycles and
operating stages for the 2-axis turning are as follows:
• straight (rectangle) turning,
• face turning,
• rough turning,
• finish turning,
• profile turning,
• groove turning,
• pocket turning,
• thread turning,
• hole machining,
• parting off (cutting off),
• rest cycle (residual turning),
• hand setting of tool motion.
The conventional roughing cycles in turning, where a
cutting tool performs a constant depth of cut can be
adapted and extended with the cycles when the tool cuts
with a variable depth of cut. The proposed roughing
cycles are as follows:
• rough cycle "decreasing of engagement" (Figure 3),
• roughing by creating a conical surface,
• roughing with the use of nonlinear methods, etc.
Figure 4 depicts a commonly used roughing cycle. A
constant depth of cut is used in this roughing cycle. The
machining process results in the wear that prevails in one
point of the cutting edge only.
During the roughing strategy – the cut decrement –
each chip removal is performed with a different depth of
cut so a different cutting part of the tool is under stress
during each cutting operation. This method of machining
can be time-consuming due to several passes. This is
compensated for by an increased tool life, a lower load-
ing of the machine spindle and a reduced machine noise.
The depth of cut is reduced when the final diameter is
being approached. The maximum wear point is therefore
moved outwards from the cut, prolonging the cutting-
tool durability. This type of roughing cycle is already
contained in the advanced CAM systems (EdgeCAM,
M. SADÍLEK et al.: INCREASING TOOL LIFE DURING TURNING WITH A VARIABLE DEPTH OF CUT
200 Materiali in tehnologije / Materials and technology 47 (2013) 2, 199–203
Figure 5: Roughing cycle creating a conical surface
Slika 5: Cikel odvzemanja, pri katerem nastaja koni~na povr{ina
Figure 4: Roughing cycle with a constant depth of cut
Slika 4: Cikel odvzemanja s konstantno globino reza
Figure 2: Tool wear according to ISO 3685:1993 6
Slika 2: Obraba orodja skladno z ISO 3685:1993 6
Figure 3: Roughing cycle – a decreased cut
Slika 3: Cikel odvzemanja – zmanj{evanje odvzema
turning to profile). An application of this feature, when
longer parts are turned, is advantageous for the elimina-
tion of the ever-decreasing workpiece stiffness.
The cutting, with which a conical surface is formed,
starts with the deepest depth of cut which decreases in
the feeding direction, as shown in Figure 5. The second
cut is programmed to be parallel with the workpiece
axis. This provides for an efficient removal of the conical
surface formed in the previous cut. Thanks to this
strategy, the tool wear moves along the cutting edge from
the maximum to the minimum depth of cut (apmax to
apmin). Figure 6 shows a simulation of the machining (a
programmed tool path) when turning the flanges, for
which the experimental part was done.
The non-linear roughing-cycle method also ensures a
variable depth of cut. For example, a tool path’s wavy
profile (Figure 7) will achieve the same effect as the
previous methods. During both the first and the second
cut, the machined material is on a gradual increase and
decrease and a variable depth of cut is thereby achieved.
It is also possible to shift the machined surface and, in
doing so, change the depth of cut. However, this requires
an advanced CAD/CAM system joined with the CNC
cutting machines.
3 EXPERIMENTAL WORK
Two strategies were used during turning: the
conventional method with a constant depth of cut and the
cone-forming machining.
3.1 Experimental characterization
The verification of theoretical presumptions was
carried out in practice. For this purpose we used the
following cutting machine: Mori Seiki SL – 65 B with
the drive system Fanuc and the spindle power of P = 71
kW. The worpiece material was the austenitic stainless
steel 1.4401 that corresponds to DIN X5CrNiMo17-2-2
with a hardness of 180 HB. This material is mainly used
in chemical industry, apparatus engineering, pulp
industry and food industry. The dimensions of the
worpiece (the flange) were as follows: its external
diameter D = 350 mm, internal diameter d1 = 56 mm,
length L = 73 mm and the internal diameter of the
semi-product before finishing d2 = 157 mm. The cutting
tool was an internal radius turning tool (the Sandvik
company) with the cutting inserts: CNMG 12 04 12 –
MR 2025 with the CVD coating. This tool is suitable for
longitudinal medium roughing and roughing. During the
cutting the cutting fluid was used. The cutting conditions
differed only in the cut size and depth shape, as shown in
Table1.
The different depths of cut (stock removals) in the
internal roughing cycle are shown in Figure 8.
Here a gradual removal of the material from the first
tool path to the last one (nth) is shown. The total number
of tool paths (n) depends on the size of the workpiece
and the technological possibilities of the insert as well.
A representative sample of the tool wear has been
selected from all of the performed experiments focusing
M. SADÍLEK et al.: INCREASING TOOL LIFE DURING TURNING WITH A VARIABLE DEPTH OF CUT
Materiali in tehnologije / Materials and technology 47 (2013) 2, 199–203 201
Figure 8: Tool paths: a) a constant depth of cut, b) a variable depth of
cut (i <1,n> = number of tool paths; n = total number of tool paths; L,
D, d1 = dimensions of the workpiece; L, D, d2 = dimensions of the
semi-product before finishing)
Slika 8: Poti orodja: a) konstantna globina reza, b) spremenljiva
globina reza (i <1,n> = {tevilo poti orodja, n = celotno {tevilo poti
orodja, L, D, d1 = dimenzije obdelovanca, L, D, d2 = dimenzije nedo-
kon~anega polproizvoda)
Figure 7: Roughing cycle – the nonlinear method
Slika 7: Cikel odvzemanja – nelinearna metoda
Figure 6: Roughing-cycle simulation when a conical surface is made
during an internal turning
Slika 6: Simulacija cikla odvzemanja, pri katerem nastaja koni~na
povr{ina pri notranjem stru`enju
on the wear in the form of a notch on the face and on the
back.
With regard to the measurement during the pro-
duction, the tool wear was measured with a microscope
with an installed digital camera. To read the wear values
the Micrometrics SE Premium software, version 7.2, was
used.
The upper part of Figure 9 shows the tool wear on
the face. The lower part of the same figure shows the
tool wear on the back. The strategy with a constant depth
of cut reports, for the same period of time (machining
time t = 18.6 min), a greater tool wear than the strategy
with a variable depth of cut.
Figure 9 reveals a visible notch on the face (with the
length xcons of 1.03 mm) and on the back formed while
using the constant depth of cut. The notch wear on the
face is greater than in the case of the strategy with a
variable ap. There is also a visible notch on the back, VBN
= 0.25 mm. The strategy using a variable depth of cut
causes no notch on the back, as shown in Figure 9, the
bottom b part).
With the variable depth of cut, the notches are distri-
buted over the tool’s longer cutting edge corresponding
to the variable depth of cut. The notches shift in relation
to the changing depth of cut (from 3 mm to 5 mm depth
of cut), xvariable = 1.49 mm.
The dependence of the tool wear on time was
recorded and entered into the "Dependence of VBN tool
wear on time" chart, as shown in Figure 10. The cutting
process was interrupted in order to measure the tool wear
after the tool path (between the cuts) so as not to disturb
the cutting process by starting a cut and getting out of a
cut and the measuring interval was kept to last for about
2 min.
This chart clearly implies that the roughing method
with a variable depth of cut causes a lower wear VBN in
the same time and under the same cutting conditions.
The wear has a slighter inclination in the second area of
the chart (in this area the tool wear increases uniformly).
The tool wear criterion was set to be VBN = 0.25 mm.
The durability of the cutting edge in the roughing
cycle with a constant depth of cut was, on average, 18
min. In the newly designed cycle with a variable depth of
cut, the durability of the cutting edge was 26 min, i.e.,
the durability increased by 44 %.
4 DISCUSSION AND CONCLUSIONS
The proposed manufacturing technology of the flange
and shaft components ensures a reduced tool wear, i.e.,
an increased turning-tool durability and life. There is a
more favorable distribution of wear on the replaceable
tool insert when employing the proposed turning techno-
logy.
An application of the new roughing cycle resulted in
a decrease in the spindle load by 10 %. This reduction
was monitored directly on the cutting machine – it was
displayed on the indicator of the spindle load.
This change caused a reduction in the energy demand
of the machine tool. The roughing cycle with a variable
depth of cut was applied in the company of JohnCrane
a.s. It gave excellent results in the form of an increased
durability of the cutting edge by 44 % in the maintained
machining time.
The increased durability of the tool significantly
reduces the total costs for the cutting tools. These costs
are also reduced by less frequent downtimes when a
worn tool is replaced. However, the disadvantage of this
M. SADÍLEK et al.: INCREASING TOOL LIFE DURING TURNING WITH A VARIABLE DEPTH OF CUT
202 Materiali in tehnologije / Materials and technology 47 (2013) 2, 199–203
Figure 10: Dependence of the VBN tool wear on time under the rough-
ing strategy – a constant depth of cut and a variable depth of cut
Slika 10: Odvisnost obrabe VBN orodja od ~asa odrezavanja – kon-
stantna globina reza in spremenljiva globina reza
Figure 9: Comparison of the tool-face and back wear: a) a constant
depth of cut, b) a variable depth of cut
Slika 9: Primerjava obrabe ~ela in za~elja orodja: a) konstantna glo-
bina reza, b) spremenljiva globina reza
Table 1: Cutting conditions: the constant-depth-of-cut strategy and the
conical-surface strategy
Tabela 1: Pogoji rezanja: strategija s konstantno globino rezanja in
strategija s koni~no povr{ino
Cutting conditions
rough turning
constant
depth of cut
variable depth of
cut
cutting speed vc m min–1 180 180
feed f mm r–1 0.3 0.3
depth of cut ap mm 4 apmin = 3 apmin = 5
technology is a more complicated tool-path program-
ming for the roughing cycle with a variable depth of cut.
The suggested roughing cycles may not have a
positive influence on the machining process. It is
necessary to carefully consider all the aspects associated
with the proposed cycles. These can include a reduced
rigidity of the machine tool caused by both the tool’s
simultaneous movement in the directions of the two axes
and the requirements of the software (CAM system) and
the CNC programmer.
Acknowledgement
We would like to thank the company of JohnCrane
a.s. for allowing us to perform our experiments. This
paper was supported by the Czech Science Foundation,
Student Grant Competition (SP2011/23), VSB-Technical
University of Ostrava.
5 REFERENCES
1 A. Antic, B. P. Petrovic, M. Zeljkovic, B. Kosec, J. Hodolic, The
influence of tool wear on the chip forming mechanism and tool
vibrations, Mater. Tehnol., 46 (2012) 3, 279–285
2 M. Sadílek, R. ^ep, I. Budak, M. Sokovic, Aspects of Using Tool
Axis Inclination Angle, Strojni{ki vestnik/Journal of Mechanical
Engineering, 57 (2011) 9, 681–688
3 M. Neslu{an, I. Mrkvica, R. ^ep, D. Kozak, R. Konderla, Defor-
mation After Heat Treatment and Their Influence on Cutting Process,
Tehni~ki vjesnik/Technical Gazette, 18 (2011) 4, 601–608
4 R. ^ep, A. Janásek, B. Martinický, M. Sadílek, Cutting tool life tests
of ceramic inserts for car engine sleeves, Tehni~ki vjesnik/Technical
Gazette, 18 (2011) 2, 203–209
5 M. Forejt, A. Humár, M. Pí{ka, L. Janí~ek, Experimental methods –
syllabus, University of Technology, Faculty of Mechanical Engineer-
ing, Brno, 2003, 83
6 ISO 3685:1993, Tool-life testing with single-point turning tools,
International Organization for Standardization, Geneva, 1993, 48
M. SADÍLEK et al.: INCREASING TOOL LIFE DURING TURNING WITH A VARIABLE DEPTH OF CUT
Materiali in tehnologije / Materials and technology 47 (2013) 2, 199–203 203

A. RAUTER et al.: RAMAN INVESTIGATION OF SOL-GEL ANTICORROSION COATINGS ON ELECTRONIC BOARDS
RAMAN INVESTIGATION OF SOL-GEL
ANTICORROSION COATINGS ON ELECTRONIC
BOARDS
RAMANSKE RAZISKAVE SOL-GEL PROTIKOROZIJSKIH
PREVLEK NA ELEKTRONSKIH VEZJIH
Aleksander Rauter1, Matja` Ko`elj1, Lidija Slemenik Per{e1, Angela [urca Vuk1, Boris Orel1,
Baºak Bengû2, Onder Sunetci2
1National Institute of Chemistry, Hajdrihova 19, 1000 Ljubljana, Slovenia
2Arcelik R&D, 34950 Cayirova Campus – Tuzla/Istanbul, Turkey
angela.surca.vuk@ki.si
Prejem rokopisa – received: 2012-09-10; sprejem za objavo – accepted for publication: 2012-10-12
Due to its non-destructive character and high spatial resolution, confocal Raman spectroscopy was found to be a good technique
for detecting spray-deposited sol-gel coatings on electronic boards (EBs). EBs are demanding substrates to cover because of
their non-flat surfaces, containing various metals, alloys, pins and other elements. Nanocomposite coatings were made on the
basis of bis end-capped organic-inorganic hybrid precursors, bis-(3-(3-(3-triethoxysilyl)propyl)thioureido)propyl terminated
polydimethylsiloxane (PDMSTU) and bis-[3-(triethoxysilyl)propyl]tetrasulphide (BTESPT), imparting a hydrophobic character
to the produced anticorrosion coatings. Reactions of hydrolysis and condensation were initiated with an addition of the 0.1 M
HCl catalyst. Nanoparticles were introduced in the sols as trisilanol-heptaisooctyl-polyhedral oligomeric silsesquioxanes (up to
5 nm). The obtained sols were spray deposited on the test substrate, i.e., the aluminium alloy AA 2024, and also on EB. The
coatings on AA 2024 were used for potentiodynamic electrochemical and surface (SEM, AFM) characterisation of the prepared
coatings. Extensive Raman spectroscopy measurements of the uncovered and covered EBs revealed that the sol-gel coatings on
various elements and pins of EBs are clearly visible in the Raman spectra and that this is an appropriate technique for detecting
the coatings on such demanding substrates.
Keywords: Raman spectroscopy, sol-gel, electronic boards, anticorrosion coatings, nanoparticles
Glede na nedestruktivni zna~aj in veliko prostorsko lo~ljivost ramanske spektroskopije smo to tehniko {tudirali s stali{~a
mo`nosti detekcije napr{enih sol-gel prevlek na elektronska vezja. Elektronska vezja so zahtevne podlage, kar zadeva njihovo
prekrivanje, saj njihova povr{ina ni ravna, na njej se nahajajo razli~ne kovine, zlitine, konice in drugi elementi. Nanokompozitne
prevleke smo pripravili na osnovi dvostransko funkcionaliziranih organsko-anorganskih prekurzorjev bis-(3-(3-(3-trieto-
ksisilil)propil)tioureido)propil zaklju~enega poli(dimetilsiloksana) (PDMSTU) in bis-[3-(trietoksisilil)propil]tetrasulfida
(BTESPT), ki v pripravljene protikorozijske prevleke vneseta hidrofoben zna~aj. Reakcije hidrolize in kondenzacije smo za~eli
z dodatkom katalizatorja 0,1 M HCl. Nanodelce smo v sole vklju~ili kot trisilanol-heptaizooktil-poliedri~ne oligomerne
silseskvioksane (do 5 nm). Tako pripravljene sole smo nanesli na preizkusne podlage, in sicer na aluminijevo zlitino AA 2024 in
tudi elektronska vezja. Prevleke na AA 2024 smo uporabili za potenciodinami~ne elektrokemijske meritve in povr{insko (SEM,
AFM) karakterizacijo. Ob{irne ramanske meritve neza{~itenih in za{~itenih elektronskih vezij pa so pokazale, da lahko sol-gel
prevleke na razli~nih elementih in konicah elektronskih vezij jasno ugotovimo v ramanskih spektrih in da je zato ta tehnika
primerna za detekcijo prevlek na zahtevnih podlagah.
Klju~ne besede: ramanska spektroskopija, sol-gel, elektronska vezja, protikorozijske prevleke, nanodelci
1 INTRODUCTION
Raman spectroscopy has been recognised as a very
useful tool for the investigation of organic and inorganic
materials in various fields, for example in art and
archaeology1, for the investigation of electrode materials
for lithium electrodes2 as well as the materials produced
from organic-inorganic hybrids3,4. Raman spectra show
the energy shift of the excitation light (laser) as a product
of an inelastic scattering of the molecules in a sample,
resulting in obtaining the information about their
chemical structure, i.e., Raman spectra represent the
fingerprints of the molecules. Moreover, this vibrational
technique can be useful for proving the correctness of
technological processes3, for instance, the efficiency of a
deposition process for the coatings on various substrates,
since it does not require background measurements and
is a non-destructive technique. Such complex substrates
include, for example, electronic boards (EBs) composed
of various materials (metals, alloys, soldering alloys,
plastic materials, etc.). Raman spectroscopy can clearly
detect sprayed protective coatings on various individual
elements of EBs.
Numerous enterprises have realised that the enor-
mous amounts of electronic waste can be reduced with
the application of anticorrosion coatings on various
electronic components. However, the anticorrosion coat-
ings for EBs must respond to a unique set of corrosion-
related problems, since different elements (materials) are
positioned in close proximity. This gives rise to a great
possibility of galvanic corrosion, the corrosion of contact
surfaces and joints or a growth of dendritic silver. In
addition, the surfaces of EBs are not flat, so it is difficult
to cover them homogeneously with a coating. Nowadays,
Materiali in tehnologije / Materials and technology 47 (2013) 2, 205–210 205
UDK 543.424.2:621.793:620.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)205(2013)
conformal paint coatings (parylene, urethane, acrylic,
epoxy, etc.) are mostly used for the protection of EBs,
but such coatings are expensive and do not usually offer
satisfactory adhesion and protection (<10 years)5. Many
conformal coatings, in particular the widely used acrylic
coatings, often delaminate from the corners of the leads
and EBs, and develop cracks, in which water may be
entrapped, dissolving contaminants. This can result in
dendritic growth between the components of the leads,
leading to shorts, excessive power consumption and,
therefore, EB malfunctions5.
Nanocomposite barrier coatings prepared from
organic-inorganic hybrid precursors, consisting of nano-
particles (SiO2, Al2O3, ZrO2, ...) or polyhedral oligomeric
silsesquioxanes (POSS) are believed to be promising
materials for successfully replacing conformal coatings.
The preparation of good sol-gel nanocomposite coatings
starts with an appropriate selection of precursors, nano-
particles, inhibitors and additives6. By introducing the
appropriate groups to the composition we can achieve, in
addition to the resistance to corrosion, that the sol-gel
coatings also provide high oxidation, abrasion and
water-resistant properties. In this work, bis end-capped
ethoxysilyl-functionalised precursor bis-(3-(3-(3-trietho-
xysilyl)propyl)thioureido)propyl terminated poly(dime-
thylsiloxane) (PDMSTU in Figure 1) was synthesised as
the organic-inorganic precursor for sol-gel nanocom-
posite coatings7. This organic-inorganic hybrid precursor
also comprises a poly(dimethylsiloxane) chain, which
imparts its hydrophobic character to the deposited
anticorrosion coatings. The nanoparticles were included
as polyhedral oligomeric silsesquioxane trisilanol-hepta-
isooctyl-POSS (TS-IOc7-POSS in Figure 1), which, due
to the open-cage structure, can bind in the sol-gel matrix
via a formation of siloxane bonds. Another bis
end-capped precursor, i.e., bis-[3-(triethoxysilyl)pro-
pyl]tetrasulphide (BTESPT in Figure 1), was included in
the composition of the sol to increase the siloxane
bonding in the coatings. The efficiency of the BTESPT
coatings for the preparation of anticorrosion coatings for
various alloys (AA 2024-T3 8, AZ31 Mg 9) has already
been demonstrated. In the present paper, we would first
like to show that the PDMSTU-based nanocomposite
coatings can be applied on EBs with the spray-deposition
technique. In addition, the suitability of Raman spec-
troscopy for detecting deposited coatings on all the
different elements of an EB substrate is shown.
2 EXPERIMENTAL WORK
The organic-inorganic hybrid precursor bis-(3-(3-
(3-triethoxysilyl)propyl)thioureido)propyl terminated
polydimethylsiloxane (PDMSTU) was synthesised from
aminopropyl terminated poly(dimethylsiloxane) 1000
(PDMS 1000)7. 23 g of PDMS 1000 was dissolved in 40
ml of tetrahydrofurane and 12.5 g of (3-isotiociana-
topropyl)triethoxysilane was then added dropwise. The
solution was refluxed for 2 days and the progress of the
reaction was followed using FT-IR spectroscopy. After
the conclusion of the reaction, tetrahydrofurane was
removed in vacuo and a highly viscous yellow PDMSTU
product was obtained.
The sol for the spray deposition on the EB substrate
was prepared in ethanol by mixing PDMSTU : BTESPT
: TS-IOc7-POSS in a molar ratio of 1 : 2 : 0.5. As a
catalyst of the sol-gel reactions, 0.1 M hydrochloric acid
was applied in a molar ratio of PDMSTU : 0.1 M HCl =
1 : 18, based on the number of ethoxy groups in the
PDMSTU and BTESPT precursor molecules. The sol
was stirred for 4 days before spraying on the EB
substrate. The time of the application was determined by
following the hydrolysis and condensation reactions
using FT-IR measurements (not shown). The sol gelled
in ten days. For a comparison between morphological
properties (SEM, AFM) and electrochemical
measurements, the PDMSTU-based coating was
deposited on the aluminium alloy AA 2024, also using
the spray-deposition technique. Prior to spraying the sol
on AA 2024, the substrate was polished and cleaned in
hexane, acetone, methanol and distilled water. The
deposited coating was thermally treated at 150 °C for
half an hour prior to the treatment.
SEM micrographs were recorded on a FE-SEM
Supra 35 VP electron scanning microscope.
An electrochemical potentiodynamic measurement of
the PDMSTU-based coating deposited on the AA 2024
substrate was recorded on a PGSTAT 302N potentiostat-
galvanostat. The coating was mounted as the working
electrode in a K0235 flat cell (Princeton Applied
Research) filled with 0.5 M NaCl. The reference
electrode was a saturated Ag/AgCl electrode and the
counter electrode was the Pt grid. The samples were held
at an open circuit potential for 1800 s prior to scanning
the potential with a scan rate of 0.5 mV/s from –0.9 V to
–0.2 V.
A. RAUTER et al.: RAMAN INVESTIGATION OF SOL-GEL ANTICORROSION COATINGS ON ELECTRONIC BOARDS
206 Materiali in tehnologije / Materials and technology 47 (2013) 2, 205–210
Figure 1: Structures of the precursors for the preparation of sol-gel
nanocomposite anticorrosion coatings for EBs: A) bis-(3-(3-(3-
triethoxysilyl)propyl)thioureido)propyl terminated poly(dimethyl-
siloxane) (PDMSTU), B) trisilanol-heptaisooctyl-POSS (TS-IOc7-
POSS) and C) bis-[3-(triethoxysilyl)propyl]tetrasulphide (BTESPT)
Slika 1: Strukture prekurzorjev za pripravo sol-gel nanokompozitnih
protikorozijskih prevlek za elektronska vezja: A) bis-(3-(3-(3-
trietoksisilil)propil)tioureido)propil zaklju~en poli(dimetilsiloksan)
(PDMSTU), B) trisilanol-heptaizooktil-POSS (TS-IOc7-POSS) in C)
bis-[3-(trietoksisilil)propil]tetrasulfid (BTESPT)
Raman spectra of the uncovered and covered EB
were collected using a confocal Raman spectrometer
WITec alpha 300, combined with the AFM and SNOM
(Scanning Near-Field Optical Microscopy) techniques.
The excitation line used was 532 nm.
3 RESULTS AND DISCUSSION
The SEM micrograph revealed a homogeneous
surface of the PDMSTU-based coating, but some defects
on the surface can also be observed (Figure 2). AFM, on
the other hand, confirmed the nanocomposite character
of the coating, with a roughness factor of 110 nm
(Figure 3).
A potentiodynamic measurement of the spray-depo-
sited PDMSTU-based coating on AA 2024 revealed a
cathodic current density that is lower, by about two
decades of magnitude, than that of the pure AA 2024
substrate (Figure 4). Moreover, the anodic current den-
sity was lower by about one to two orders of magnitude
(Figure 4). These values are comparable to those of the
anticorrosion coatings prepared by dip-coating on the
AA 2024 substrates from 1 % sols of bis-(3-(3-(3-
triethoxysilyl)propyl)ureido)propyl terminated poly-
dimethylsiloxane (PDMSU)10, i.e., a similar precursor
with the urea instead of thiourea groups between the
poly(dimethylsiloxane) and ethoxysilylpropyl parts of
the molecule (Figure 1). The PDMSU coatings, on the
other hand, showed a better performance when prepared
from more concentrated, 4 %, sols, i.e., the sols with an
increasing thickness confirming their physical barrier
character. Moreover, our PDMSTU-based coatings
revealed superior properties compared to the pure
BTESPT coatings made from water/ethanol solutions8.
Specifically, the hydrophobic BTESPT coatings revealed
a decrease of about one order of magnitude in the
cathodic current density and of about 1–2 orders of
magnitude in the anodic current density. A decrease in
the cathodic current density of one order of magnitude in
relation to the response of the pure AA 2024 substrate
was also observed for the ethoxysilyl-functionalised
POSS compounds, enabling a formation of compact
tri-dimensional networks of regular POSS cubes via
sol-gel reactions11. However, the addition of per-
fluoropropyl groups to the structure of these POSS
compounds increased the reduction in the cathodic
current density, even to above two orders of magnitude12,
which is superior to that of our PDMSTU-based
coatings. The contact-angle measurements showed that
poly(dimethylsiloxane) chains introduced some of the
hydrophobic character to the PDMSTU-based coatings;
the contact angle for water was 97° for the coatings on
AA 2024 in the initial state and the surface-energy value
was 32.1 mJ/m2. The value was higher than the sur-
face-energy values for the coatings prepared from
PDMSU (29.7 mJ/m2) 10, ethoxysilyl-functionalised POSS
(28.7 mJ/m2) 11 and perfluoropropyl ethoxysilyl-func-
A. RAUTER et al.: RAMAN INVESTIGATION OF SOL-GEL ANTICORROSION COATINGS ON ELECTRONIC BOARDS
Materiali in tehnologije / Materials and technology 47 (2013) 2, 205–210 207
Figure 4: Potentiodynamic measurement of the spray-deposited
PDMSTU-based coating on AA 2024
Slika 4: Potenciodinami~ne meritve prevleke, pripravljene na osnovi
PDMSTU in napr{ene na podlago AA 2024
Figure 2: SEM micrograph of the spray-deposited PDMSTU-based
coating on AA 2024
Slika 2: SEM-posnetek prevleke, pripravljene na osnovi PDMSTU in
napr{ene na podlago AA 2024
Figure 3: AFM image of the spray-deposited PDMSTU-based coating
on AA 2024
Slika 3: AFM-posnetek prevleke, pripravljene na osnovi PDMSTU in
napr{ene na podlago AA 2024
tionalised POSS (12.4 mJ/m2) 12. Since a good corrosion
inhibition can be partly ascribed to the low surface value
of the anticorrosion coatings, the coatings with the
perfluoropropyl groups have been found to be the best in
this respect. For comparison, the contact angle for water
of the AA 2024 substrate was 86° and the surface-energy
value was 37.1 mJ/m2, respectively.
The most critical issue in the preparation of the
anticorrosion coatings for EBs, apart from the deve-
lopment of the coatings’ composition, is the coverage of
all the elements (except the specific connectors that are
masked during the spray deposition of a coating) on the
non-flat EB surfaces. The coverage of pins often remains
questionable, since the coating solutions may tend to
flow down the slopes. When the sols exhibit proper
rheological properties, this effect of sagging is prevented
and all the elements are efficiently covered by the
coating. Such defects in coverage can certainly be
recognised during various end industrial application
tests, but these are time consuming and costly. The
proposed Raman spectroscopy is a more straightforward
and quicker way of determining a possible non-coverage
in the coatings. The PDMSTU-based coating sprayed on
EB was therefore placed under the laser beam of the
Raman spectrometer (Figure 5) and investigated at
different points and elements (Figure 6). In Figure 6,
only the elements that are described in this text (pins,
SMD resistor and SMD capacitor) are shown, but other
positions on EB were also studied.
Figure 7 shows the Raman spectra recorded on the
pins of the uncoated and coated EBs. The spectra of the
uncoated pins revealed two peaks, at 1.385 cm–1 and
1.360 cm–1, which can be assigned to the presence of
carbon. Namely, the spectrum of the ash sample revealed
similar bands that are marked in the literature as G- and
D-lines2. The so-called G-band corresponds to the strong
C-C stretching E2g2 mode of the hexagonal graphite
structure. The D-line (A1g) that appeared from 1.350
cm–1 to 1.360 cm–1, on the other hand, is the consequence
of a turbostratic disorder, i.e., a disorder along the c axis
due to a weak interlayer bonding2. When the EB
substrate was covered with the PDMSTU-based coating,
the characteristic bands of the coating appeared in the
Raman spectrum (Figure 7). For comparison, the same
bands also appeared in the Raman spectrum of the
spray-deposited PDMSTU-based coating on the AA
A. RAUTER et al.: RAMAN INVESTIGATION OF SOL-GEL ANTICORROSION COATINGS ON ELECTRONIC BOARDS
208 Materiali in tehnologije / Materials and technology 47 (2013) 2, 205–210
Figure 6: Positions of the pins and other elements that were measured
on the uncoated and coated EB using Raman spectroscopy
Slika 6: Pozicije konic in drugih elementov, na katerih smo izmerili
ramanske spektre na neza{~itenem in za{~itenem elektronskem vezju
Figure 7: Raman spectra of the uncoated and coated pins on the EB
substrate. The PDMSTU-based coating on the AA 2024 substrate is
given for comparison, as well the characteristic Raman spectrum of an
independent ash sample.
Slika 7: Ramanski spektri neza{~itenih in za{~itenih konic na elek-
tronskem vezju. Za primerjavo podajamo ramanski spekter prevleke
na osnovi PDMSTU, ki je napr{ena na podlago AA 2024. Podajamo
tudi zna~ilen ramanski spekter vzorca saj.
Figure 5: Measurement of the Raman spectra on the uncoated and
coated EB substrate
Slika 5: Merjenje ramanskih spektrov neza{~itenega in za{~itenega
elektronskega vezja
2024 substrate. Similarly, in the case of the SMD resistor
and SMD capacitor, it was also found that the bands of
the PDMSTU-based coating appeared in the Raman
spectrum (Figure 8). In addition to the modes of the
coating, carbon bands were also observed on the SMD
resistor. On the other hand, the uncoated SMD capacitor
revealed some bands by itself and they remained in the
Raman spectrum of the covered SMD capacitor, in
addition to the modes of the PDMSTU-based coating
(Figure 8). The depicted spectra measured on the
covered EB clearly revealed that the coating was
successfully deposited covering all the different elements
and parts of the investigated substrate. Other elements
and pins were also investigated, but the results have not
been shown due to a lack of space.
The Raman spectra measured on the coated EB
substrate (Figures 7 and 8) clearly revealed that Raman
spectroscopy is extremely well suited to the investigation
of complex objects that do not allow a sample prepa-
ration. This is also due to the improvements in instru-
ment configurations, for example, the use of confocal
micro-Raman spectrometers that allow effective spatial
resolution and Raman imaging as well as achieving a
better contrast1,2. When non-typical samples, such as EBs
are in question, Raman spectroscopy is important due to
its non-destructive character, certainly when carefully set
excitation conditions are used. In addition, the detection
of coatings on the metallic parts of EBs is easy, since
almost all the metals show no Raman bands. However,
Raman spectroscopy is not a useful quantitative tech-
nique and it would be difficult to determine the amount
of coating on various elements of EBs, since the sensiti-
vity of the Raman response depends on the polarizability
of the molecules and the characteristics of the measured
spot. Another important drawback of Raman spectro-
scopy is the fluorescence caused by certain organic or
other materials, i.e., in the case of the EB substrate, it
was found on plastic areas. Nevertheless, all the afore-
mentioned facts clearly show that Raman spectroscopy is
a suitable technique for detecting the protective sol-gel
coatings on EBs, not only after a deposition of the
coatings, but also after their exposure to accelerated
corrosion tests. The excellent spatial resolution of the
confocal Raman spectrometers allows a detection of the
corroded parts of the coatings on EBs and thus a
determination of the most vulnerable parts of EB
substrates.
4 CONCLUSIONS
Potentiodynamic electrochemical characterisation of
the coatings prepared from organic-inorganic precursors
bis-(3-(3-(3-triethoxysilyl)propyl)thioureido)propyl
terminated poly(dimethylsiloxane) (PDMSTU) and
bis-[3-(triethoxysilyl)propyl]tetrasulphide (BTESPT)
revealed that the coatings possess the properties similar
to the other coatings made of poly(dimethylsiloxane) or
tetrasulphide-based precursors. Specifically, the cathodic
current density of spray-deposited coatings was lower,
by up to two orders of magnitude, than the cathodic
current density of the pure AA 2024 substrate. The
prepared coatings were also spray deposited on an
electronic board (EB) and Raman spectroscopy was
shown to be an appropriate technique for detecting the
sol-gel protective coating on this substrate due to its
non-destructive character and a high spatial resolution.
On the basis of the present results, Raman spectroscopy
will be used in future to determine the sites, at which
corrosion has started after an exposure of the coated EB
substrates to accelerated corrosion tests, including
Raman imaging of larger surfaces. In order to assure a
complete coverage of the elements, the sols will be
studied using a detailed rheological characterisation. The
sol with the optimal rheological properties will be
chosen for further application.
Acknowledgement
This research was funded by the Slovenian Research
Agency (Programme P1-0030) and the Ministry of
Education, Science, Culture and Sport (MNT-ERA.NET
project Bonaco, http://www.bonaco-project.com). A.
A. RAUTER et al.: RAMAN INVESTIGATION OF SOL-GEL ANTICORROSION COATINGS ON ELECTRONIC BOARDS
Materiali in tehnologije / Materials and technology 47 (2013) 2, 205–210 209
Figure 8: Raman spectra of the uncoated and coated SMD-resistor
and SMD-capacitor elements on the EB substrate. The PDMSTU-
based coating on the AA 2024 substrate is given for comparison.
Slika 8: Ramanski spektri neza{~itenega in za{~itenega SMD-upora in
SMD-kondenzatorja na elektronskem vezju. Za primerjavo podajamo
ramanski spekter prevleke na osnovi PDMSTU, ki je napr{ena na
podlago AA 2024.
Rauter thanks the Slovenian Research Agency for the Ph.
D. grant.
5 REFERENCES
1 P. Vandenabeele, H. G. M. Edwards, L. Moens, Chem. Rev., 107
(2006), 675–686
2 R. Baddour-Hadjean, J. Pierre Pereira-Ramos, Chem. Rev., 110
(2010), 1278–1319
3 M. Gnyba, M. Keränen, M. Kozanecki, R. Bogdanowicz, B. B. Kos-
mowski, P. Wroczyñski, Opto-Elecronics Rev., 10 (2002), 137–143
4 I. Jerman, A. [urca Vuk, M. Ko`elj, F. [vegl, B. Orel, Prog. Org.
Coat., 72 (2011), 334–342
5 McCullough, J. L. Wayt, J. N. Butch, US patent 6,127,038, 2000
6 D. Wang, G. P. Bierwagen, Prog. Org. Coat., 64 (2009), 327–338
7 M. Ko`elj, Synthesis of substituted trialkoxysilanes and their
application for the preparation of materials via sol-gel procedures,
Dissertation, University of Ljubljana, Ljubljana
8 D. Zhu, W. J. van Ooij, Electrochim. Acta, 49 (2004), 1113–1125
9 M. F. Montemor, M. G. S. Ferreira, Electrochim. Acta, 52 (2007),
7486–7495
10 M. Fir, B. Orel, A. [urca Vuk, A. Vil~nik, R. Je{e, V. Franceti~,
Langmuir, 23 (2007), 5505–5514
11 I. Jerman, A. [urca Vuk, M. Ko`elj, B. Orel, J. Kova~, Langmuir, 24
(2008), 5029–5037
12 I. Jerman, B. Orel, A. [urca Vuk, M. Ko`elj, J. Kova~, Thin Solid
Films, 518 (2010), 2710–2721
A. RAUTER et al.: RAMAN INVESTIGATION OF SOL-GEL ANTICORROSION COATINGS ON ELECTRONIC BOARDS
210 Materiali in tehnologije / Materials and technology 47 (2013) 2, 205–210
P. TERNIK et al.: NUMERICAL STUDY OF RAYLEIGH-BÉNARD NATURAL-CONVECTION HEAT-TRANSFER ...
NUMERICAL STUDY OF RAYLEIGH-BÉNARD
NATURAL-CONVECTION HEAT-TRANSFER
CHARACTERISTICS OF WATER-BASED Au
NANOFLUIDS
NUMERI^NA ANALIZA PRENOSA TOPLOTE NANOTEKO^IN
VODA-Au V RAZMERAH REYLEIGH-BÉNARDOVE NARAVNE
KONVEKCIJE
Primo` Ternik1, Rebeka Rudolf2,3, Zoran @uni~4
1Private Researcher, Bresterni{ka ulica 163, 2354 Bresternica, Slovenia
2University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia
3Zlatarna Celje, d. d., Kersnikova ul. 19, 3000 Celje, Slovenia
4AVL-AST, Trg Leona [tuklja 5, 2000 Maribor, Slovenia
pternik@pt-rtd.eu
Prejem rokopisa – received: 2012-10-01; sprejem za objavo – accepted for publication: 2012-10-15
The present work deals with the natural convection in a square cavity filled with a water-based Au nanofluid. The cavity is
heated from the lower and cooled from the adjacent wall, while the other two walls are adiabatic. The governing differential
equations have been solved with the standard finite volume method and the hydrodynamic and thermal fields have been coupled
using the Boussinesq approximation. The main objective of this study is to investigate the influence of the nanoparticles’ volume
fraction on the heat-transfer characteristics of Au nanofluids at a given base-fluid (i.e., water) Rayleigh number Rabf .
Accurate results are presented over a wide range of the base-fluid Rayleigh numbers (102  Rabf  105) and the volume fraction
of Au nanoparticles (0 %    10 %). It is shown that adding nanoparticles to the base fluid delays the onset of convection.
Contrary to what is argued by many authors, we show, with numerical simulations, that the use of nanofluids can reduce the heat
transfer instead of increasing it.
Keywords: Rayleigh-Bénard natural convection, water-Au nanofluid, heat transfer, numerical modelling
V prispeveku obravnavamo naravno konvekcijo v kvadratni kotanji, napolnjeni z nanoteko~ino voda-Au. Kotanja je bila greta s
spodnje in hlajena s prile`ne zgornje stene, preostali dve steni sta bili adiabatni. Vodilne diferencialne ena~be smo re{evali s
standardno metodo kon~nih prostornin. Hidrodinami~no in temperaturno polje sta bila sklopljena z uporabo Boussinesqove
aproksimacije. Glavni cilj prispevka je raziskati vpliv prostorninskega dele`a nanodelcev na zna~ilnosti prenosa toplote
Au-nanoteko~ine pri podani vrednosti Rayleighjevega {tevila nosilne teko~ine (vode) Rabf .
Natan~ni rezultati so predstavljeni za {iroko obmo~je vrednosti Rayleighjevega {tevila nosilne teko~ine (102  Rabf  105) in
prostorninskega dele`a Au-nanodelcev (0 %    10 %). Pokazali smo, da dodajanje nanodelcev v nosilno teko~ino zakasni
za~etek naravne konvekcije. V nasprotju s trditvami mnogih avtorjev smo z numeri~nimi simulacijami pokazali, da lahko
uporaba nanodelcev prenos toplote zmanj{a in ne pove~a.
Klju~ne besede: Rayleigh-Bénardova naravna konvekcija, nanoteko~ina voda-Au, prenos toplote, numeri~no modeliranje
1 INTRODUCTION
Buoyancy-induced flow together with the associated
heat transfer is an important phenomenon found in many
engineering applications (e.g., selective laser melting
process1, cooling of electronic devices2). An enhance-
ment of heat transfer in such systems is crucial from the
energy-saving point of view. In recent years, nanosized
particles dispersed in a base fluid, known as nanofluid,
has been used and researched extensively to enhance the
heat transfer. The presence of nanoparticles shows an
unquestionable heat-transfer enhancement in forced
convection applications3. However, with respect to the
buoyancy-driven flow, there is still a dispute on the effect
of nanoparticles on the heat-transfer enhancement.
Several researchers have been focused on the
numerical modelling of buoyancy-induced flows. Recent
numerical studies by Ternik et al.4, Ternik and Rudolf5,
Oztop et al.6 and Abu-Nada and Oztop7 illustrated that
the suspended nanoparticles substantially increase the
heat-transfer rate for any given Rayleigh number. In
addition, they showed that the heat-transfer rate in
water-based nanofluids increases with an increasing
volume fraction of Al2O3, Cu, TiO2 and Au nanoparti-
cles.
On the other hand, an apparently paradoxical beha-
viour of the heat-transfer deterioration was observed in
many experimental studies8-10. For example, Putra et al.8
reported that a presence of Al2O3 nanoparticles in a base
fluid reduces the natural convective heat transfer.
However, they did not clearly explain why the natural
convective heat transfer is decreased with an increase in
the volume fraction of nanoparticles.
The above review of the existing literature shows that
the problem of natural convection in a bottom-heated
horizontal cavity filled with a nanofluid is an issue still
Materiali in tehnologije / Materials and technology 47 (2013) 2, 211–215 211
UDK 519.61/.64:620.3:536.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)211(2013)
far from being completely solved. Framed in this general
background, the purpose of the present study is to exa-
mine the effect of adding Au nanoparticles to the base
fluid at the conduction and convection heat-transfer rates
in a square cavity heated from below (Rayleigh–Bénard
configuration) over a range of base-fluid Rayleigh
numbers 102  Rabf  105 and volume fractions 0 % 
  10 %.
2 NUMERICAL MODELLING
The standard finite-volume method, successfully
used in many recent studies,11–13 is used to solve the
coupled conservation equations of mass, momentum and
energy. In this framework, a second-order central diffe-
rencing scheme is used for the diffusive terms and a
second-order upwind scheme for the convective terms.
Coupling of the pressure and velocity is achieved using
the SIMPLE algorithm. The convergence criteria were
set to 10–9 for all the relative (scaled) residuals.
2.1 Governing equations
For the present study, a steady-state flow of an
incompressible water-based Au nanofluid is considered.
It is assumed that both the fluid phase and nanoparticles
are in thermal equilibrium. Except for the density, the
properties of the nanoparticles and fluid (presented in
Table 1) are taken to be constant. The Boussinesq appro-
ximation is invoked for the nanofluid properties to relate
density changes to temperature changes, and to couple
the temperature field with the velocity field.
The governing equations (mass, momentum and
energy conservation) of such a flow are:4,5
∂
∂
 i
ix
= 0 (1)
 




nf j
i
j j
nf
i
j
i
nf
x x x
p
x
g T
∂
∂
∂
∂
∂
∂
∂
∂
−
⎛
⎝
⎜
⎞
⎠
⎟ =
= − + −( ( T
x xj
nf
j
i
C )+
⎛
⎝
⎜
⎞
⎠
⎟
∂
∂
∂
∂


(2)
(  c
T
x x
c
T
xp nf j j j
nf
j
∂
∂
∂
∂
∂
∂
=
⎛
⎝
⎜
⎞
⎠
⎟ (3)
where the cold-wall temperature TC is taken to be the
reference temperature for evaluating the buoyancy term
()nfg(T – TC) in the momentum conservation equation.
Relationships between the properties of the nanofluid
(nf) and those of the base fluid (bf) and pure solid (s) are
given with the following empirical models4,5:
• Dynamic viscosity:


nf
bf
=
−( ) .1 2 5
• Density:
   nf bf s= − +( )1
• Thermal expansion:
   nf bf s= − +( )1
• Heat capacitance:
( ( )( (      c cp nf p bf s= − +1
• Thermal conductivity:
k k
k k k k
k k k knf bf
s bf bf s
s bf bf s
=
+ − −
+ + −
2 2
2


( )
( )
2.2 Geometry and boundary conditions
The simulation domain is shown schematically in
Figure 1. The two horizontal walls of a square enclosure
are kept at different constant temperatures (TH > TC),
whereas the other boundaries are considered to be
adiabatic. Both velocity components (i.e., vx and vy) are
identically zero on each boundary because of the no-slip
condition and the impenetrability of the rigid boundaries.
In the present study, the heat-transfer characteristics
are presented in terms of the mean Nusselt number:
Nu
L
Nu x x
L
= ∫
1
0
( )d (4)
P. TERNIK et al.: NUMERICAL STUDY OF RAYLEIGH-BÉNARD NATURAL-CONVECTION HEAT-TRANSFER ...
212 Materiali in tehnologije / Materials and technology 47 (2013) 2, 211–215
Table 1: Thermo-physical properties of the Au nanofluid4,5
Tabela 1: Toplotno-fizikalne lastnosti Au-nanoteko~ine4,5
 (Pa s)  (kg/m3) cp (J/kg K) k (W/m K)  (1/K)
Pure water 1.003 × 10–3 997.1 4179 0.613 2.1 × 10–4
Au / 19320 128.8 314.4 1.416× 10–7
Figure 1: Schematic diagrams of the simulation domain
Slika 1: Shematski prikaz obmo~ja simulacije
and the ratio of the nanofluid heat-transfer rate to the
base-fluid one:
Q
Q
k Nu
k Nu
h
k
nf
bf
nf nf
bf bf
nf
bf
= = (5)
where hnf and hbf are the convection heat-transfer coeffi-
cients of the nanofluid and the base fluid.
In order to investigate the influence of volume frac-
tion  on the heat-transfer characteristics, the Rayleigh
(Ranf) and the Prandtl numbers (Prnf) for the nanofluids
are expressed as follows:
Ra
k c
k c
Ranf
nf bf p nf bf
bf nf p bf nf
bf
n
=
   
   
(
(
Pr f
nf p nf bf
bf p bf nf
bf
c k
c k
=


Pr
(6)
Using equation (6) we show that Ranf < Rabf (Figure
2a) and Prnf < Prbf (Figure 2b) for all the values of
.The ratio of the water-Au-nanofluid Rayleigh and
Prandtl numbers to the base-fluid Rayleigh and Prandtl
numbers decreases with the increasing volume fraction
of Au nanoparticles.
2.3 Grid-dependency study
The grid independence of the results has been
established on the basis of a detailed analysis of three
different uniform meshes: M1(50 × 50), M2(100 × 100)
and M3(200 × 200. For the general primitive variable 
the grid-converged (i.e., extrapolated to the zero element
size) value according to Richardson extrapolation is
given as11,12, ext = M3 – (M2 – M3)/(rp – 1) where M3 is
obtained on the basis of the finest grid and M2 is the
solution based on the next level of the coarse grid, r = 2
is the ratio between the coarse- and the fine-grid spacing
and p = 2 is the order of accuracy.
The numerical error e = I(M2 – 
ext )/
extI for the
mean Nusselt number Nu is presented in Table 2. It can
be seen that the differences in the grid refinements are
exceedingly small and the agreement between mesh M2
and the extrapolated value is extremely good (the discre-
tisation error is well below 0.2 %). Based on this, the
simulations in the remainder of the paper were conduc-
ted on mesh M2 that provided a reasonable compromise
between high accuracy and computational efficiency.
Table 2: Effect of a mesh refinement upon the mean Nusselt number
( = 0.10, Rabf = 105)
Tabela 2: Vpliv zgo{~evanja mre`e na srednjo vrednost Nusseltovega
{tevila ( = 0,10, Rabf = 105)
Mesh M1 Mesh M2 Mesh M3 Nuext e
3.325 3.304 3.299 3.298 0.184 %
2.4 Benchmark comparison
In addition to the aforementioned grid-dependency
study, the simulation results have also been compared
with the recent results of Turan et al.14 for the Rayleigh-
Bénard natural convection in a square cavity. The com-
parisons between the present-simulation results and the
corresponding benchmark values (summarised in Table
3) are very good and entirely consistent with our grid-
dependency studies.
Table 3: Comparison of the present results for Nu with the benchmark
results
Tabela 3: Primerjava pri~ujo~ih rezultatov za Nu z referen~nimi rezul-
tati
Ra = 103 Ra = 104 Ra = 105
Pr = 1 Pr = 10 Pr = 1 Pr = 10 Pr = 1 Pr = 10
Present study 1.000 1.000 2.164 2.190 3.941 3.875
Turan et al.14 1.000 1.000 2.162 2.188 3.934 3.868
3 RESULTS AND DISCUSSION
Figure 3 presents the variation of the mean Nusselt
number (equation 4) along the hot wall for different
values of Rabf and Ranf. For Rabf < 2586, there is no
convection in the nanofluid or the base fluid, and the
heat transfer occurs due to pure conduction, so the mean
Nusselt number equals 1 and is independent of the base-
P. TERNIK et al.: NUMERICAL STUDY OF RAYLEIGH-BÉNARD NATURAL-CONVECTION HEAT-TRANSFER ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 211–215 213
Figure 2: Variation of the dimensionless numbers of the water-Au
nanofluid with the volume fraction of Au nanoparticles: a) Rayleigh
number and b) Prandtl number
Slika 2: Spreminjanje brezdimenzijskih {tevil nanoteko~ine voda-Au s
prostorninskim dele`em Au-nanodelcev: a) Rayleighjevo {tevilo in b)
Prandtlovo {tevilo
fluid Rayleigh number (Figure 3a). As the base-fluid
Rayleigh number increases, the nanofluid remains in the
conductive regime, while convection appears in the base
fluid. The point of transition (i.e., the value of Rabf) from
conduction to convection depends on the volume fraction
of Au nanoparticles. The higher is the value of , the
more delayed is the onset of convection (Figure 3a).
When the nanofluid is in the convective heat-transfer
regime, the mean Nusselt number is a monotonic
increasing function of Rabf.
On the other hand, it is interesting to notice that the
transition from conduction to convection occurs at the
same value of the nanofluid Rayleigh number, i.e., Ranf
 2586 (Figure 3b). Furthermore, the value of the mean
Nusselt number at a given Ranf is practically independent
of the nanoparticles’ volume fraction. This finding is a
reflection of the nanofluid Prandtl number values con-
sidered in the present study. Its value varies (decreases
with the increasing) from Prnf ( = 0 %) = 6.84 to Prnf (
= 10 %) = 2.26 and for this range of the Prandtl-number
values (Pr > 1) the relative balance between viscous and
buoyancy forces is modified, so that the heat transport in
the thermal boundary layer gets only marginally affec-
ted14. This marginal modification is reflected in a weak
Prandtl-number (and therefore nanoparticles’ volume
fraction) dependence of the mean Nusselt number.
Figure 4 shows the effect of the base-fluid Rayleigh
number on the ratio of the heat-transfer rate for the
water-based Au nanofluid for different values of the
volume fraction. In the range Rabf < 2586 the heat
transfer occurs by pure conduction, so the ratio of heat
transfer is equal to the ratio of thermal conductivities and
is constant and independent of Rabf. For Ranf < 2586 and
Rabf > 2586 the nanofluid remains in the conductive
regime, while convection appears in the base fluid. The
heat transfer is more important in the base fluid than in
the nanofluid and the ratio Qnf/Qbf is on a decrease until
Ranf  2586. From this point onwards (i.e., the transition
from conduction to convection) the ratio Qnf/Qbf is on an
increase and its value becomes higher than 1, but
remains lower than the ratio that is obtained when both
the nanofluid and the base fluid are in the conductive
regime. When the ratio Qnf/Qbf > 1, the heat-transfer rate
in the nanofluid becomes higher than that in the base
fluid.
Finally, in Figure 4, we observe that the heat transfer
can decrease or increase depending on the value of the
base-fluid Rayleigh number. For example, for a water-
based Au nanofluid and for  = 10 %, the ratio of the
heat-transfer rate becomes higher than 1 when the base-
fluid Rayleigh number reaches the value of around 9500,
so we obtain an enhancement of the heat transfer only
after this value (Rabf > 9500). Therefore, adding nano-
particles increases the heat transfer only for a given
value of the temperature difference.
4 CONCLUSIONS
In the present study, the steady laminar natural con-
vection of water-based Au nanofluids in a square enclo-
P. TERNIK et al.: NUMERICAL STUDY OF RAYLEIGH-BÉNARD NATURAL-CONVECTION HEAT-TRANSFER ...
214 Materiali in tehnologije / Materials and technology 47 (2013) 2, 211–215
Figure 4: Effect of the base-fluid Rayleigh number on the ratio of heat
transfer
Slika 4: Vpliv Rayleighjevega {tevila nosilne teko~ine na razmerje
prenosa toplote
Figure 3: Variation of the mean Nusselt number along the hot wall
with the: a) base-fluid Rayleigh number and b) nanofluid Rayleigh
number
Slika 3: Spreminjanje srednjega Nusseltovega {tevila vzdol` tople ste-
ne z: a) Rayleighjevim {tevilom nosilne teko~ine in b) Rayleighjevim
{tevilom nanoteko~ine
sure with differentially heated horizontal walls and with
the bottom wall at a higher temperature has been nume-
rically analysed. The effects of the base-fluid Rayleigh
number (102  Rabf  105) and the solid volume fraction
(0 %    10 %) on heat-transfer characteristics have
been systematically investigated in detail.
The influence of a computational grid refinement on
the present numerical predictions was studied throughout
the examination of the grid convergence at Rabf = 105 and
 = 10 %. By utilizing extremely fine meshes the result-
ing discretisation error for Nu is well below 0.2 %.
The numerical method was validated for the case of
Rayleigh-Bénard natural convection in a square cavity,
for which the results are available in the open literature.
A remarkable agreement of the present results with the
benchmark results of Turan et al.14 yields sufficient
confidence in the present numerical procedure and its
results.
Highly accurate numerical results pointed out some
important points such as:
• In the classical Rayleigh–Bénard configuration, just
after the onset of convection, there is more heat
transfer in the base fluid than in the nanofluid. For a
fixed value of the base-fluid Rayleigh number Rabf,
the nanofluid Rayleigh number Ranf decreases with
the volume fraction of nanoparticles. Thus, the nano-
particles delay the onset of convection.
• In the convective heat-transfer regime the mean
Nusselt number Nu is found to increase with the
increasing values of the base-fluid Rayleigh number
Rabf, but the Nu values obtained for the higher values
of the nanoparticles’ volume fraction  are smaller
than those obtained in the case of the base fluid ( =
10 %) with the same numerical values of Rabf.
• The transition from the conductive to convective
heat-transfer regime occurs at the same value of the
nanofluid Rayleigh number, i.e., Ranf  2586.
• The values of the mean Nusselt number at a given
Ranf are practically independent of the nanoparticles’
volume fraction.
• The heat transfer can decrease or increase depending
on the value of the Rayleigh number. So, an addition
of nanoparticles increases the heat transfer only for
the given values of the temperature difference.
Acknowledgements
The research leading to these results was carried out
within the framework of a research project "Production
technology of Au nano-particles" (L2-4212) and has
received the funding from the Slovenian Research
Agency (ARRS).
5 REFERENCES
1 N. Contuzzi, S. L. Campanelli, A. D. Ludovico, 3D finite element
analysis in the selective laser melting process, International Journal
of Simulation Modelling, 10 (2011), 113–121
2 A. Ijam, R. Saidur, Nanofluid as a coolant for electronic devices
(cooling of electronic devices), Applied Thermal Engineering, 32
(2012), 76–82
3 W. Daungthongsuk, S. Wongwises, A critical review of convective
heat transfer in nanofluids, Renewable & Sustainable Energy
Reviews, 11 (2009), 797–817
4 P. Ternik, R. Rudolf, Z. @uni~, Numerical study of heat transfer
enhancement of homogeneous water-Au nanofluid under natural
convection, Mater. Tehnol., 46 (2012) 3, 257–261
5 P. Ternik, R. Rudolf, Heat transfer enhancement for natural
convection flow of water-based nanofluids in a square enclosure,
International Journal of Simulation Modelling, 11 (2012), 29–39
6 H. F. Oztop, E. Abu-Nada, Y. Varol, K. Al-Salem, Computational
analysis of non-isothermal temperature distribution on natural
convection in nanofluid filled enclosures, Superlattices and
Microstructures, 49 (2011), 453–467
7 E. Abu-Nada, H. F. Oztop, Effects of inclination angle on natural
convection in enclosures filled with Cu–water nanofluid, Interna-
tional Journal of Heat and Fluid Flow, 30 (2009), 669–678
8 N. Putra, W. Roetzel, S. K. Das, Natural convection of nano-fluids,
Heat and Mass Transfer, 39 (2002), 775–784
9 B. H. Chang, A. F. Mills, E. Hernandez, Natural convection of
microparticles suspension in thin enclosures, International Journal of
Heat and Mass Transfer, 51 (2008), 1332–1341
10 C. J. Ho, W. K. Liu, Y. S. Chang, C. C. Lin, Natural convection heat
transfer of alumina-water nanofluid in vertical square enclosures: an
experimental study, International Journal of Thermal Sciences, 49
(2010), 1345–1353
11 I. Bilu{, P. Ternik, Z. @uni~, Further contributions on the flow past a
stationary and confined cylinder: Creeping and slowly moving flow
of Power law fluids, Journal of Fluids and Structures, 27 (2011),
1278–1295
12 P. Ternik, New contributions on laminar flow of inelastic non-New-
tonian fluid in the two-dimensional symmetric expansion: Creeping
and slowly moving conditions, Journal of Non-Newtonian Fluid
Mechanics, 165 (2010), 1400–1411
13 K. T. Rai}, R. Rudolf, P. Ternik, Z. @uni~, V. Lazi}, D. Stamenkovi},
T. Tanaskovi}, I. An`el, CFD analysis of exothermic reactions in
Al-Au multi-layered foils, Mater. Tehnol., 45 (2011) 4, 335–338
14 O. Turan, N Chakraborty, R. J. Poole, Laminar Rayleigh-Bénard
convection of yield stress fluids in a square enclosure, Journal of
Non-Newtonian Fluid Mechanics, 171–172 (2012), 83–96
P. TERNIK et al.: NUMERICAL STUDY OF RAYLEIGH-BÉNARD NATURAL-CONVECTION HEAT-TRANSFER ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 211–215 215

J. BURJA et al.: ALUMOTHERMIC REDUCTION OF ILMENITE IN A STEEL MELT
ALUMOTHERMIC REDUCTION OF ILMENITE IN A
STEEL MELT
ALUMOTERMI^NA REDUKCIJA ILMENITA V JEKLENI TALINI
Jaka Burja1, Franc Tehovnik1, Jakob Lamut2, Matja` Knap2
1 Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2 University of Ljubljana, Faculty of Natural Sciences and Engineering, Department of Materials and Metallurgy, A{ker~eva 12,
1000 Ljubljana, Slovenia
jaka.burja@imt.si
Prejem rokopisa – received: 2012-10-15; sprejem za objavo – accepted for publication: 2013-01-04
Experiments regarding the alumothermic reduction of ilmenite (FeO·TiO2), a mineral that contains iron and titanium oxides,
were carried out. The results of the experiments showed that the alumothermic reduction takes place in steel, but the products
obtained with the reduction suggest that the reaction mechanism is more complicated and includes different metallic phases
other than a simple reduction of pure elemental titanium. Two kinds of experiments were carried out, the alumothermic
reduction of ilmenite in a steel melt and the alloying of the alumothermic mixture into the steel melt. The experiments were
carried out in order to get a view of the phase boundary between the steel and reduced titanium and the metallic phases that
occur during the reduction, before the dissolution of titanium in the steel melt. Metallic phases that contained aluminium, iron
and titanium were gained during the alumothermic reduction of ilmenite. Titanium was successfully alloyed into the steel melt
by introducing the alumothermic mixture into the melt, while the presence of titanium nitrides confirms that the titanium was
reduced in the melt and reacted with the dissolved nitrogen.
Keywords: alloying of titanium, ilmenite, alumothermic reduction, titanium in steel
Izvedeni so bili poskusi alumotermi~ne redukcije ilmenita. Ilmenit je mineral, ki vsebuje titanove in `elezove okside
(FeO·TiO2). Rezultati so potrdili, da v jeklu pote~e alumotermi~na redukcija. Produkti redukcije pa ka`ejo na to, da je reakcijski
mehanizem bolj zapleten, kot pa zgolj nastajanje elementarnega titana, saj nastajajo razli~ne kovinske faze. Izvedeni sta bili dve
vrsti poskusov: alumotermi~na redukcija ilmenita in legiranje alumotermi~ne me{anice v jekleno talino. Dobljen je bil vpogled
v procese na fazni meji med jekleno talino, titanom in nastalimi kovinskimi fazami, preden se titan raztopi v jeklu. Ugotovljeno
je bilo, da se titan raztaplja v jeklu prek nastanka intermetalnih faz. Produkti redukcije so bile kovinske faze, ki so vsebovale
aluminij, titan in `elezo. Prisotnost titanovih nitridov v jeklu, ki smo ga legirali z me{anico, pa je dokaz, da je bil titan legiran v
kovinski obliki, kjer je reagiral z raztopljenim du{ikom.
Klju~ne besede: legiranje titana, ilmenit, alumotermi~na redukcija, titan v jeklu
1 INTRODUCTION
Titanium is an important alloying element in steel
making, among other things, it is used to stabilise
stainless steel by forming titanium carbides and prevent-
ing the formation of chromium carbides. It has also been
observed that additions of titanium significantly reduce
the austenite grain size in the as-cast microstructure of
continually cast steels.1 Titanium’s high chemical affinity
to nitrogen and carbon is what makes it such a valuable
alloying element, but unfortunately it also makes it
difficult to alloy (low yields) and produce.2 The pro-
duction of titanium is complex and therefore expensive.3
In steelmaking titanium is used in the form of ferro-
titanium that contains iron and between 20 to 75 % of
titanium, while its eutectic composition is at the mole
fraction 71.1 % of Ti.4,5 Ferrotitanium is mostly pro-
duced by remelting titanium scrap and iron; the high
prices of titanium consequently mean that the price of
ferrotitanium is also relatively high. Experiments that
concern direct alloying of titanium from the oxide form
may show an alternative way of producing ferrotitanium,
as the minerals like ilmenite that contain titanium oxides
are inexpensive. Alumothermic reduction was chosen
because aluminium is often added to ferrotitanium in
order to increase the yield by reducing the oxidised
titanium in a steel melt.6
The Ellingham diagram (Figure 1) clearly shows that
the Gibbs free energy for the alumothermic reduction of
titanium is negative. The alumothermic reduction of
titanium in the oxide form is as follows:
3TiO2 + 4Al = 3Ti + 2Al2O3
The alumothermic reduction of ilmenite has an even
lower Gibbs free energy because ilmenite contains iron
oxides and its equation is as follows:
FeO·TiO2 + 2Al = Fe + Ti + Al2O3
As we can see from reaction (2) iron is another metal
product besides titanium. The graph for the value of the
Gibbs free energy for equations 1 and 2 is given in
Figure 27.
Titanium forms TiO2 if oxygen is present in the steel
melt, while titanium oxide forms high temperature
phases with other oxide components in the slag. If CaO
is present perovskite CaO·TiO2 forms with its melting
point at around 2000 °C. Titanium oxide particles can
also get trapped in spinel Al2O3·MgO, which can also
Materiali in tehnologije / Materials and technology 47 (2013) 2, 217–222 217
UDK 669.094.23:66.046.51:549.641.23 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)217(2013)
present a problem, because a spinel melts at the tempe-
ratures higher than 1600 °C and the titanium oxide
particles have a melting point at around 1800 °C.8 The
formation of non-metallic inclusions that have a high
melting point is problematic with respect to the clean-
liness of steel and is therefore undesirable. The control
of oxide non-metallic inclusions can be achieved by
lowering the aluminium content and, therefore, the Al2O3
content and by decreasing the MgO content in the
rafination slag.8 The CaO·TiO2 content is lowered by
lowering the basicity of the slag (CaO/SiO2).8
In the production of stainless steel, a strong nitride-
forming element, such as titanium, is often added to
stabilise nitrogen and improve the mechanical properties
of steel via the grain refinement during hot rolling. On
the other hand, titanium nitride formed in liquid steel can
agglomerate and cause a nozzle-clogging problem during
continuous casting, and surface defects in the final pro-
ducts.9 In practice the deposit material that is clogging
the nozzle is titanium oxide and spinel, but research
work has shown that the agglomerates form because
titanium nitride particles get caught in the spinel; these
are in turn oxidised and become titanium oxides. The
particles remain rectangular like nitrides and not globular
like oxides formed in the melt. The presence of oxygen
is probably the result of porosity of the refractory mate-
rial, from which the nozzles are made. The flow of the
metal trough the nozzle creates a low pressure, which, in
turn, promotes the diffusion of oxygen trough the pores
into the melt.10
The experiments were carried out taking into account
the specific nature of titanium and its behaviour as an
alloying element.
2 EXPERIMENTS
The experiments were carried out using steel pipes
filled with an alumothermic mixture of aluminium and
ilmenite. These pipes were then introduced into the steel
melt where the mixture was heated up to approximately
the temperature of the steel melt. The mixture reacted at
such high temperatures. In one set of the experiments the
pipes were retrieved before they fully dissolved and still
contained the mixture, which had been sintered by the
high temperatures. The other set of experiments had been
designed to alloy titanium into the steel and in this case
the pipes were fully dissolved, together with the mixture.
This procedure was carried out in order to recreate the
conditions of alloying with the use of a cored wire.
The alumothermic mixture consisted of the ilmenite
dust and aluminium dust. The molar ratio was 1 : 4 for
ilmenite to aluminium, and the surplus of aluminium was
J. BURJA et al.: ALUMOTHERMIC REDUCTION OF ILMENITE IN A STEEL MELT
218 Materiali in tehnologije / Materials and technology 47 (2013) 2, 217–222
Figure 1: Ellingham-Richardson diagram
Slika 1: Ellingham-Richardsonov diagram
Figure 3: Ilmenite grains
Slika 3: Ilmenitna zrna
Figure 2: Gibbs free energy for equation 27
Slika 2: Gibbsova prosta energija za reakcijo 27
used so that the reduction took place. Theoretically, a
molar ratio of 1 : 2 is needed to reduce ilmenite, as can
be seen from equation 2.
Ilmenite is a mineral that mainly consists of iron and
titanium oxides in the form of FeO·TiO2, but it also
contains impurities such as magnesium oxides and man-
ganese oxides. SEM analyses show that grains of ilme-
nite are far from being homogenous but have a "striped"
appearance that can be seen in Figure 3.
The phase analysis in Figure 4 clearly shows that the
darker stripes are richer in titanium, while the lighter
ones are richer in iron.
An electric induction furnace with a capacity to melt
18 kg of steel was used to melt the steel for the experi-
ments.
A steel melt was prepared for the experiments that
were developed to model the alumothermic reduction of
ilmenite. The steel melt was heated to the temperature of
1500 °C, measured with an optical pyrometer. The che-
mical composition of the steel melt is not important for
the experiment because the steel melt is only used as a
medium to transfer heat to the alumothermic mixture,
not to interact with it chemically. Then the pipe with the
alumothermic mixture was submerged into the melt for
approximately 40 s. The pipe was 25.4 mm in diameter
and had a wall thickness of 2 mm. This method was
chosen because it was essential that the mixture was
quickly heated up to the temperature of the steel melt in
order to prevent the aluminium dust from oxidising and
thus losing its ability to reduce ilmenite.
Another experiment was made in order to determine
whether titanium can be alloyed into the steel melt with
the alumothermic ilmenite reduction. The aim was to
alloy the mass fraction 0.3 % of titanium into 18 kg of
steel. The required amount of ilmenite was 171 g and the
amount of aluminium was 121 g.
First the steel was melted; it had a small quantity of
alloying elements and a deep drawing quality. When the
steel was melted a sample was taken for a chemical
analysis. It was found that it did not contain detectable
levels of titanium (the method of determining the
titanium content was the classical chemical analysis).
Then the melt was deoxidised with aluminium, after that
the steel pipes were filled with the alumothermic mixture
and submerged into the melt until they dissolved, thus,
alloying the steel with titanium when the pyrometer gave
a temperature reading of 1600 °C.
A sample of the alloyed steel was taken after the
alloying was complete; the rest of the steel was cast into
an ingot. Samples of slag were taken as well.
3 RESULTS
The steel pipes that were used for the reduction of
ilmenite were partially melted, but there was a sintered
mass in the pipe. Metallographic samples were made and
a further SEM analysis showed that metallic phases con-
taining aluminium, iron and titanium were present.
Figure 5 shows the products of the alumothermic reac-
tion; metallic phases contain aluminium, iron and, most
importantly, titanium.
Figure 6 shows a part of the sintered mass in the
steel pipe after the reduction; individual intermetallic
phases can be seen. The metallic phases are surrounded
by the oxide products of the reduction.
The chemical compositions of the phases from
Figure 6 are given in Table 1.
In the experiment that investigated the option of
alloying titanium, the mixture had reacted and the
J. BURJA et al.: ALUMOTHERMIC REDUCTION OF ILMENITE IN A STEEL MELT
Materiali in tehnologije / Materials and technology 47 (2013) 2, 217–222 219
Figure 5: Mapping of alumothermic products
Slika 5: Ploskovna porazdelitev elementov redukcijskih produktov
Figure 4: Phase analysis of ilmenite
Slika 4: Fazna analiza ilmenita
product or the alumothermic reduction began to dissolve
into the steel melt. The presence of titanium nitrides in
the microstructure, as seen in Figure 7, confirms that
titanium had indeed been alloyed into the steel and had
reacted with the nitrogen in the steel melt. The chemical
analysis showed that the content of titanium had been
raised up to w = 0.064 %. The yield of titanium can be
calculated with equation 2 and is therefore 21 %.
Yield = 0.064/0.3 × 100 % = 21 %
Figure 8 shows the SEM analysis of the nitrides
confirming that the inclusions in Figure 7 were indeed
titanium nitrides.
J. BURJA et al.: ALUMOTHERMIC REDUCTION OF ILMENITE IN A STEEL MELT
220 Materiali in tehnologije / Materials and technology 47 (2013) 2, 217–222
Figure 8: SEM image of a titanium nitride
Slika 8: SEM-posnetek titanovega nitrida
Figure 6: SEM image of the products of the alumothermic reduction
Slika 6: SEM-posnetek produktov alumotermi~ne redukcije ilmenita
Table 1: Chemical compositions of the phases (in mass fractions w/%)
from Figure 6
Tabela 1: Sestava faz (v masnih dele`ih w/%) s slike 6
1 48.4 % Al 6.0 % Ti 43.4 % Fe 1.3 % Cu 0.9 % Mn
2 52.2 % Al 45.5 % Ti 1.6 % Fe 0.7 % Si
3 33.5 % Al 30.0 % Ti 28.7 % Fe 1.6 % Cu 5.4 % Si 0.8 % Mn
4 44.5 % Al 42.7 % Ti 12.8 % Fe
Figure 9: Mapping of the slag
Slika 9: Ploskovna porazdelitev elementov v `lindri
Figure 7: Titanium nitrides in the steel microstructure
Slika 7: Titanovi nitridi v mikrostrukturi jekla
Next the slag was analysed in order to further widen
the understanding of the processes that took place during
alloying. An interesting discovery was made during the
analysis of the slag: a relatively large content of metallic
phases.
The content of the elements is clearly shown in
Figure 9, where the metallic parts in the slag are mostly
aluminium, but the most important factor is that the
metallic parts contain reduced titanium and that can be
directly linked to a lower yield. There is also a signi-
ficant amount of titanium in the oxide part of the slag as
shown in Figure 9. A part of the alloying mixture had
clearly floated onto the surface and got entrapped in the
slag.
4 DISCUSSION
Alloying titanium into the steel melt by alumother-
micaly reducing ilmenite can be divided into several
stages: heating up the alumothermic mixture, the alumo-
thermic reaction, the formation of alloys and inter-
metallic phases of iron, titanium and aluminium, the
dissolution of the intermetallic phases and, therefore,
titanium into the steel melt, and the reaction between
titanium and nitrogen in the steel melt. The first part of
the experiments showed us that elemental titanium does
not form, instead intermetallic phases are formed and
they consist mostly of aluminium. These results can be
compared with those of N.J. Welham and associates,
considering that the surplus of aluminium they used was
significantly higher.11 Possible future work should
consider that aluminium is not needed just for the
reduction, but for forming the alloys with titanium. The
phase that contained the highest amount of titanium,
45.5 %, the phase number 2 from Table 1, was based on
aluminium. It is clear that the tendency of titanium to
form intermetallic phases with aluminium is higher than
that to form them with iron. It can be speculated, from
the phase diagram Ti–Al (Figure 10), that Al3Ti with
63 % Al and 27 % Ti forms during the alumothermic
reduction.12 But in the cases that have a lower surplus of
aluminium, AlTi should form, as can be seen from the
Ti–Al phase diagram.
The experiments that dealt with the alloying of
titanium into the steel melt show us quite a different
problem than that of getting the right molar ratio of the
ingredients for the reduction: the problem of a lower
density and, therefore, buoyancy. A large part of the
titanium that was reduced in the melt ended up in the
slag due to its floating onto the surface of the melt and
into the slag. The products of the alumothermic reduc-
tion were found in the slag together with aluminium. The
metallic phases, especially aluminium, in the slag indi-
cate that not only did some products of the reduction get
entrapped in the slag, but aluminium and unreduced
ilmenite did as well and the reduction also took place in
the slag. The other part of the titanium was alloyed into
the melt and formed nitride inclusions in the steel melt.
A part of the alumothermic products clearly had time to
dissolve in the melt.
A study of the thermodynamic stability of the inter-
metallic phases of aluminium and titanium and their
effect on the thermodynamics and kinetics of the reduc-
tion should be made and their ability to dissolve in liquid
steel should be observed. Further experiments with
cored-wire injections should be studied. The effect of
such alloying on the number and size of non-metallic
inclusions should be studied as well.
5 CONCLUSIONS
Titanium can be alloyed into the steel melt by alumo-
thermicaly reducing ilmenite.
Intermetallic phases of titanium and aluminium, not
elemental titanium, form during the reduction.
The forming of intermetallic phases of titanium and
aluminium requires an even higher surplus of aluminium.
The low density of the ingredients for the alumo-
thermic reduction and the products further decreases the
yield.
6 REFERENCES
1 M. Ohno, C. Murakami, K. Matsuura, K. Isobe, Effects of Ti
Addition on Austenite Grain Growth during Reheating of As-Cast
0.2 mass% Carbon Steel, ISIJ International, 52 (2012) 10,
1832–1840
2 J. Jong-Oh, K. Wan-Yi, K. Dong-Sik, P. Jong-Jin, Thermodynamics
of Titanium, Nitrogen, and Oxygen in Liquid Alloy Steels, Metals
and Materials International, 14 (2008) 4, 531–537
3 Extraction of Titanium – the Kroll process, available from: http://
wwwchem.uwimona.edu.jm/courses/titanium.html
4 G. Volkert, K. D. Frank, Metallurgie der ferrolegierungen, Berlin
Heidelberg New York, 1972
5 H. Okamoto, Fe-Ti (iron-titanium), Journal of Phase Equilibria, 17
(1996) 4, 369
J. BURJA et al.: ALUMOTHERMIC REDUCTION OF ILMENITE IN A STEEL MELT
Materiali in tehnologije / Materials and technology 47 (2013) 2, 217–222 221
Figure 10: Ti-Al binary phase diagram12
Slika 10: Binarni fazni diagram Ti – Al12
6 B. Korou{i~, J. Triplat, B. Arh, Improvements To The Production
Process For Stainless Steel Alloyed With Titanium, Mater. Tehnol.,
37 (2003) 6, 347–352
7 HTC Chemistry 5, (software)
8 J. Wan Kim, S. Koo Kim, D. Sik Kim, Y. Deuk Lee, P. Keun Yang,
Formation Mechanism of Ca-Si-Al-Mg-Ti-O Inclusions in Type 304
Stainless Steel, ISIJ International, 36 (1996), S140–S143
9 P. Jong-Jin, J. Yong-Soo, H. In-Kook, C. Woo-Yeol, K. Dong-Sik, L.
Yun-Yong, Thermodynamics of TiN Formation in Fe-Cr Melts, ISIJ,
45 (2005) 8, 1106–1111
10 G. Yang, K. Kenichi Sorimachi, Formation of Clogging Materials in
an Immersed Nozzle during Continuous Casting of Titanium
Stabilized Stainless Steel, ISIJ International 33 (1993) 2, 291–297
11 N. J. Welham, Mechanochemical reaction between ilmenite (FeTiO )
and aluminium, Journal of Alloys and Compounds, 270 (1998) 1,
228–236
12 ASM handbook, Vol. 3 – Introduction to Alloy Phase Diagrams,
ASM International, 1992
J. BURJA et al.: ALUMOTHERMIC REDUCTION OF ILMENITE IN A STEEL MELT
222 Materiali in tehnologije / Materials and technology 47 (2013) 2, 217–222
D. SOJER et al.: ANALYSIS OF CORROSION PROPERTIES OF MELT SPUN Nd-Fe-B RIBBONS ...
ANALYSIS OF CORROSION PROPERTIES OF MELT
SPUN Nd-Fe-B RIBBONS COATED BY ALUMINA
COATINGS
ANALIZA KOROZIJSKIH LASTNOSTI HITRO STRJENIH
Nd-Fe-B-TRAKOV, OPLA[^ENIH Z ALUMINIJEVIM OKSIDOM
David Sojer1, Irena [kulj2, Spomenka Kobe1, Janez Kova~1, Paul John McGuiness1
1Institut "Jo`ef Stefan", Jamova cesta 39, 1000 Ljubljana, Slovenia
2Magneti Ljubljana, d. d., Stegne 37, 1000 Ljubljana, Slovenia
davidsojer@yahoo.com
Prejem rokopisa – received: 2012-12-04; sprejem za objavo – accepted for publication: 2013-02-05
We have coated Nd-Fe-B melt spun powders, used for the production of bonded magnets via the sol-gel route by Al2O3.
Topography and chemical composition of as-spun and protected ribbons was compared by Auger electron spectroscopy, X-ray
photoelectron spectroscopy, secondary electron spectroscopy and electron diffraction spectroscopy. To determine the corrosion
properties, we have conducted a Highly accelerated stress test, at 110 °C and 90 % humidity, followed by measuring the weight
change. To confirm the effectiveness of the coated layer, magnetic properties were compared with a vibrating sample
magnetometer. Al2O3 coatings resulted in superior corrosion resistance and magnetic properties and thus expanding the
applicability of bonded magnets to severe atmospheric conditions.
Keywords: Nd-Fe-B, bonded magnets, HAST, Al2O3 coatings
S sol-gel metodo smo opla{~ili hitro strjene Nd-Fe-B-trakove za uporabo pri izdelavi plastomagnetov. Za analizo topografije in
kemi~ne sestave hitro strjenih in opla{~enih trakov smo uporabili Augerjevo elektronsko spektroskopijo, rentgensko
fotoelektronsko spektroskopijo, spektroskopijo sekundarnih elektronov in elektronsko difrakcijsko spektroskopijo. Za dolo~itev
protikorozijske u~inkovitosti opla{~enja smo izvedli pospe{eni stresni preizkus pri temperaturi 110 °C in 90-odstotni vla`nosti,
~emur je sledila analiza masnih izgub. Z vibracijskim magnetometrom smo primerjali magnetne lastnosti prahov pred
opla{~enjem in po njem. Opla{~enje z Al2O3 se je izkazalo kot izvrstna protikorozijska za{~ita materiala, ki je odli~no za{~itila
tudi magnetne lastnosti. Aplikacije plastomagnetov se s tem lahko raz{irijo tudi na podro~ja, ki zahtevajo zahtevnej{e
atmosferske razmere, tj. pri vi{jih temperaturah in visoki vla`nosti.
Klju~ne besede: Nd-Fe-B, plastomagneti, HAST, opla{~enje z Al2O3
1 INTRODUCTION
Nd-Fe-B magnets market continues to grow1. They
possess a high magnetic energy product, a combination
of a high remanence (Br) and sufficient coercivity (HcI), a
much desired property allowing miniaturization and
diversification of application. Automotive industry, hard
drives, or wind turbines are just am exemple of applica-
tion where their use can be found2.
Nd-Fe-B magnets can be produced by sintering or by
bonding. Sintered magnets possess magnetic properties
as high as with a Br up to 1.4 T and an HcI up to 2500
kA/m 3. Machining is used to give the magnets their final
shape. In contrast to sintered magnets, bonded Nd-Fe-B
magnets are produced by blending various polymer
binders, such as epoxy or nylon, where melt-spun
ribbons play the role of raw material4. The blend is cast
moulded or injection moulded into final shape. There-
fore, bonded magnets are chosen when complex shapes
are demanded, with the advantage of being less expen-
sive than sintered magnets. The magnetic properties of
bonded magnets are, however, inferior to those of
sintered magnets. Commercial bonded magnets have a Br
up to 0.7 T and an HcI up to 1400 kA/m.
Nd-Fe-B magnets are very susceptible to corrosion,
mainly due to the high rare-earth content in the grain-
boundary phase5. Susceptibility to corrosion is naturally
increased at elevated temperatures and especially in
humid environments, affected by rare-earth elements.
This restricts not only their range of applications, but can
also damage the magnetic material during the moulding
process6–8.
Adjusting the composition, by adding Co, Ga or TiC
to the melt-spun ribbons9,10 improves the corrosion
resistance, but it also affects the magnetic properties, by
reducing either the HcI or Br – or both of them. However,
even with an optimized composition, additional corro-
sion protection is necessary. Attempts to protect the
surface of ribbons have ben reported, such as direct
surface oxidation during rapid solidification, electro-
plating, chemical vapour deposition and applying thin
films of SiO2, TiO2 and MgO via the sol-gel route11,12.
But, difficulties from the perspective of process control
and mass production appear. Applying a thin film of
SiO2 or TiO2 through a sol-gel route offered very good
corrosion resistance, especially for TiO2. The MgO thin
film was investigated in strongly acidic media, and good
corrosion resistance was reported, but with no corrosion
Materiali in tehnologije / Materials and technology 47 (2013) 2, 223–228 223
UDK 621.78.019.84:621.318.1:537.622 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(2)223(2013)
test in a humid environment. MgO has the same dis-
advantage as rare-earth elements, since both react rapidly
in the presence of humidity.
The sol-gel approach was to some extent put to one
side in the past, since the chemical reagents (precursors)
used to be extremely expensive. But now, sol-gel precur-
sors are increasingly widely used and more commer-
cially available, while the price of Nd-Fe-B raw material
has become extremely expensive and Nd-Fe-B magnets
are in greater demand for use at high temperatures. As a
result, the sol-gel approach is becoming very interesting
from the commercial perspective.
In this paper we present an investigation of sol-gel-
applied coatings of Al2O3, which were applied directly to
the melt-spun ribbons. This method has several advan-
tages over the coating of final products or other coatings,
reported by other authors. Firstly, the ribbons are pro-
tected from oxidation during transport; secondly, such a
coating offers better protection to the final magnets
because each piece of melt-spun ribbon is coated; and
thirdly, it tends not to reduce the overall Br of the final
magnet, which is a problem associated with coating the
outside of an already-bonded magnet. Such a coating can
provide a much better corrosion-resistant protection than
the standard protective film, usually applied around the
final-shaped bonded magnet. A protective layer can be
scratched off the magnet by force during transport or
while installing the magnet into the final application.
The same surface scratching does not damage the
bonded magnets where the coating has been applied to
each individual powder particle, since all the powder
particles are coated.
We have investigated the influence of applied coat-
ings on the corrosion resistance of powders in high
humidity, using the Highly accelerated stress test
(HAST) 13. The presence of the coating layers was
determined by electron diffraction spectrometry (EDS),
and X-ray photoelectron spectroscopy (XPS). The
appearance of the surface of the samples was analysed
by scanning electron microscopy (SEM) before and after
the HAST using secondary electron imaging (SEI) as
well as backscattered-electron imaging (BEI). The mag-
netic properties of the samples before and after exposure
to the HAST were analyzed with a vibrating-sample
magnetometer (VSM).
2 EXPERIMENTAL
MQB commercial melt-spun ribbons, put through a
sieving analysis, in the mesh between 100 μm and 250
μm were selected for the experimental work. We chose
this size range because this is the standard size for the
production of bonded Nd-Fe-B magnets. We analysed
the powders using a SEM JEOL 7200. The samples for
the SEM analysis were prepared by mixing the powder
of selected size with acrylic resin, which was followed
by the standard metallographic procedure of grinding
and polishing. The result was a metallographic sample
that contained a large number of powder particles, where
we could observe the cross-section of the particles.
This was followed by the sol-gel process; the details
of the sol-gel process are presented in Figure 1. The
powders were first degreased and cleaned in isopropanol
(IPA) and acetone, then dried at 60 °C, 1 h. 0.5 g of
Al2O3 precursor, aluminium isopropoxide was added to
100 mL of IPA, respectively. 10 g of MQB powder in
were added to 100 mL of IPA and stirred. The mixture of
precursor and IPA was then slowly added to a mixture of
MQB powders and IPA, while stirring to ensure an even
distribution of precursor, powder and IPA. No water was
added to promote the chemical reaction. The mixture of
powder, aluminium isopropoxide and IPA was stirred for
10 min, and then the IPA was removed from the pow-
ders. The powders were then dried at 60 °C for 15 min,
D. SOJER et al.: ANALYSIS OF CORROSION PROPERTIES OF MELT SPUN Nd-Fe-B RIBBONS ...
224 Materiali in tehnologije / Materials and technology 47 (2013) 2, 223–228
Figure 1: Scheme of the sol-gel process
Slika 1: Shema sol-gel procesa
followed by additional drying at 120 °C for 30 min. The
coated powders were then put in a glass tube under an
argon atmosphere and heated to 450 °C for 7 min to
densify the Al2O3 coating. Three coatings were applied
by repeating the above-described process two more
times. This treatment was applied because we wanted to
thicken the Al2O3 coating and to ensure that all the
particles were covered with Al2O3, leading to full
coverage of the surface.
The presence of the coating layer was confirmed by
XPS spectrometer produced by Physical Electronics Inc.,
model TFA XPS. The Al monochromatized source of
X-ray light with the power of 200 W was used. The
energy of the X-ray beam was 1486.6 eV. The analysis
area was 0.4 mm in diameter. The energy scale was
aligned to the carbon C 1s spectrum at 284.8 eV. Main
peaks in this spectrum are O 1s, C 1s, Al 2p, Al 3s and
Fe 2p. Unfortunately the Nd 3d peak at binding energy
of 980 eV was not possible to identify in this spectrum
due to strong overlap with oxygen KLL peak at this
energy. Also the boron peak B 1s at energy of 190 eV is
not visible due to low sensitivity of B in XPS spectro-
scopy.
We also took advantage of the XPS and used it to
perform a profile analysis on the Al2O3 coated particles
in order to determine the thickness of the Al2O3 coating
layer. The sputtering rate was estimated to be about
2 nm/min on the Ni layer of known thickness. Relative
sensitivity factors provided by instrument producer were
used to calculate elemental concentrations. The analysis
depth of the XPS method is about 3–5 nm.
HAST followed the XPS analysis to determine the
corrosion behaviour of the uncoated and coated particles.
The HAST experiment was performed in an industry-
standard HAST chamber, type Kambi~. The conditions
of the corrosion test were 110 °C, 90 % humidity, and
the duration of the corrosion test was set to 192 h (8 d).
The powders were carefully weighed to (2 ± 0.01) g
before the HAST test, with an accuracy of ±0.1 mg, and
placed in separate Al2O3 vessels. During the corrosion
test, the powders were removed from the HAST chamber
every 48 h, with the purpose of following the corrosion
rate. After each removal the vessels containing the
powders were dried at 120 °C, for exactly 30 min, then
weighed. This step was repeated until the final 192 h.
Weight change was calculated as a weight increase
expressed in a mass fraction (w/%)using the equation:
(mass (x h) – mass(0 h)) / mass (0 h)
with x being the number of hours inside the HAST
chamber.
We also used the SEM to observe the surface topo-
logy of the MQB powders to compare the effect of tem-
perature and humidity on coated and uncoated particles.
VSM measurements were conducted on a Lakeshore
7307 VSM, which was used to compare the magnetic
properties of the as-spun and the sol-gel-coated ribbons
with the ribbons subjected to the HAST test.
3 RESULTS AND DISCUSSION
The SEM analysis of the selected powders showed
that the actual size distribution is larger. The reason for
that can be seen in Figure 2. The powder particles are far
from being spherical; rather they are thin, elongated
particles of irregular shape.
Following the SEM analysis the particles were
cleaned. In this way we were hoping to remove all the
dust, grease, and at least partially also the Nd-oxide layer
from the surface. Cleaning the surface was also
necessary as part of the preparation for coating via the
sol-gel route.
EDS analysis of the sol-gel-processed powders
showed the presence of aluminium. The results are given
in Table 1. Figure 3 presents a SEI image of the coated
powders, with spots indicating the points of EDS
analyses. The oxygen content cannot be determined
using EDS, for the reasons already mentioned. Also, it is
D. SOJER et al.: ANALYSIS OF CORROSION PROPERTIES OF MELT SPUN Nd-Fe-B RIBBONS ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 223–228 225
Figure 3: SEI image of Al2O3 sol-gel-processed powder, with marked
spots for EDS analysis
Slika 3: SEI-posnetek prahu, opla{~enega z Al2O3 po sol-gel metodi, z
ozna~enimi to~kami, kjer so bile izvedene EDS-analize
Figure 2: SEI image of MQB as-spun ribbons
Slika 2: SEI-posnetek MQB hitro strjenih trakov
clear that the aluminium concentration varies a lot,
probably for the same reasons as mentioned in the
previous section for SiO2 coated particles. Nevertheless,
aluminium was present at all the measured points, from
which we concluded that the particles were fully
covered.
Table1: Results from the EDS analysis of Al2O3 sol-gel-processed
powder
Tabela 1: Rezultati EDS-analize prahu, opla{~enega z Al2O3 po
sol-gel metodi
Element spot 1 (x/%) spot 2 (x/%)
Al 14.8 3.53
Fe 73.77 82.43
Nd 11.43 14.04
To prove the presence of an Al-oxide structure on the
particle surfaces we used XPS. Since XPS can be used
for depth-profile analyses we tried to determine not only
the composition of the top layer, but also to determine
how thick the oxide layer is. The results of the depth
profile by XPS are presented in Figure 4. It is clear that
Al and oxygen are present in high concentrations
throughout the measured depth profile. There is also
some Nd present in the top layer. In contrast, iron is not
present in the top layer; its concentration rises slowly
towards the end of the measured profile.
From this we could draw two conclusions. First,
Al-oxide of some sort is present throughout the
measured profile. The most probable composition of this
oxide is Al2O3, since it is the product of the sol-gel
reaction. Second, the measured compositions are rather
stable throughout the measured profile. This is attributed
to the diameter of the sputtering beam, which was
estimated to be 0.4 mm and the possible tilt angle of the
measured particle to the sputtering beam. First is beyond
the size range of the measured powder (0.1–0.25 mm),
meaning that the sputtering might have been conducted
on neighbouring powder particles, resulting in a very
stable, and misleading, concentration depth profile. The
tilt angle is difficult to control, since we are dealing with
rather small particles, which align randomly. So, we
could have been sputtering a particle, which was not
exactly perpendicular to the sputtering beam, causing a
seemingly thicker layer of elements than they actually
are. On the other hand, it could mean this is the actual
depth profile, meaning the oxide layer is thicker than the
measured profile.
3.1 HAST
Uncoated, and Al2O3-coated particles were put
through the HAST, as described previously. The mass
change in the mass fractions (%) is given in Figure 5.
All the particles gained weight rapidly during the first 48
hours. However, the mass increase of the uncoated
particles was much higher than that of all the coated
particles. The Al2O3 coated only once provided good
corrosion protection, but Al2O3 was clearly more effec-
tive. After the first 48 h the rate of oxidation of the Al2O3
coated particles was still approximately three times
slower than that of the uncoated particles.
It is unknown to us, why the weight increase rate in
the first 48 h is much higher than in the following time.
It could be possible that moist oxidized the Nd-Fe-B
flakes in case of uncoated particles, while in the case of
sol-gel coated particles some moist diffuses into the
gelated coating, due to H2O deprivation inside the coated
layer. Second option is that some water diffused through
the coating layers under the effect of high pressure and
temperature. It would be impossible to evaporate this
water during drying. But it is difficult to provide infor-
mation, whether this water formed to create oxides. If it
did, we should observe some spalling behaviour of the
coating layer on the SEI images on Figure 5. But no
spalling was found.
The SEM images shown in Figure 6 reveal the diffe-
rence between the uncoated and coated particles. The
uncoated particles are clearly corroded, while the Al2O3
D. SOJER et al.: ANALYSIS OF CORROSION PROPERTIES OF MELT SPUN Nd-Fe-B RIBBONS ...
226 Materiali in tehnologije / Materials and technology 47 (2013) 2, 223–228
Figure 5: Effectiveness of the Al2O3 coating during the HAST
Slika 5: U~inkovitost Al2O3 opla{~enja po HAST-u
Figure 4: XPS depth profile on Al2O3 coated powder
Slika 4: Profilna XPS-analiza prahu, opla{~enega z Al2O3
coated particles look unaffected by the conditions of the
test. The Al2O3 particles that were coated three times, in
particular, do not look any different than prior to the
HAST experiment. Also, we could not spot any spalling
behaviour.
3.2 VSM measurements
We measured uncoated and Al2O3 coated MQB
powders before and after the HAST. The results pre-
sented as a B – H curve, normalized to sample mass can
be seen in Figures 7 and 8. For a clearer view, Figure 7
contains only the results of the powders prior to the
HAST, while Figure 8 represents powders after the
HAST test, and the MQB powder prior to HAST test, as
a comparison. The HcI is plotted on the x-axis, while the
Br is plotted on the y-axis.
As can be seen in Figure 8, the MQB uncoated and
Al2O3 coated samples’ magnetic properties are practi-
cally identical.
As for the HcI of the uncoated powder, it dropped by
20 % after the HAST. This drop in HcI can be attributed
to corrosion, which changes the microstructure of the
affected particle. Al2O3-coated powder lost magnetic
none of the magnetic properties after the HAST, showing
the coating layer to be extremely effective.
4 CONCLUSIONS
1. We have successfully coated Nd-Fe-B powders with
Al2O3 using the sol-gel process.
2. VSM measurements proved the effectiveness of
Al2O3 coating, since magnetic properties remained as
prior to the HAST test.
3. The sol-gel coating technique of Al2O3 proved to be a
reliable and non-complicated process for enhancing
the corrosion properties of the Nd-Fe-B powders,
suitable for mass production.
4. This method is unique compared to other techniques,
since it protects all the individual magnetic powder
particles, making the corrosion protection much more
reliable. Thus, the magnetic raw material is protected
D. SOJER et al.: ANALYSIS OF CORROSION PROPERTIES OF MELT SPUN Nd-Fe-B RIBBONS ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 223–228 227
Figure 8: VSM measurements of the uncoated powder and coated
powders after the HAST
Slika 8: Primerjava magnetnih lastnosti neopla{~enih in opla{~enih
prahov po HAST-u
Figure 6: SEI images of particles after HAST: a) uncoated, b) Al2O3 –
three times coated
Slika 6: SEI-posnetki pra{nih del~kov po HAST-u: a) neopla{~eni, b)
opla{~eni z Al2O3 (3-krat)
Figure 7: VSM comparison of magnetic properties between powders
prior to the HAST
Slika 7: Primerjava magnetnih lastnosti neopla{~enih in opla{~enih
prahov pred HAST-om
from atmospheric conditions throughout the produc-
tion process, starting with the transport of the
powder, the moulding, installing into the final
application, and use during the final application.
5. With the price of used metal alkoxide precursor fall-
ing rapidly, this process offers a commercially
affordable corrosion-protection route for Nd-Fe-B
melt spun ribbons used in bonded magnets.
6. We have demonstrated that the application of bonded
magnets can be expanded to high-temperature and
high-humidity environments, ranging up to 110 °C
and 90 % humidity.
Acknowledgments
This research was partially financed by the European
Union.
5 REFERENCES
1 L. Yang, Development of NdFeB Magnet Industry in New Century,
Journal of Iron and Steel Research, 13 (2006), 1–11
2 L. Yang, The inexorable rise of China’s NdFeB magnet industry,
Metal Powder Report, 63 (2009), 8–10
3 B. M. Ma, J. W. Herchenroeder, B. Smith, M. Suda, D. N. Brown, Z.
Chen, Recent development in bonded NdFeB magnets, Journal of
Magnetism and Magnetic Materials, 239 (2002), 418–423
4 H. Masaaki, Overview and outlook of bonded magnets in Japan,
Journal of Alloys and Compounds, 222 (1995), 8–12
5 G. Yan, P. J. McGuiness, J. P. G. Farr, I. R. Harris, Environmental
degradation of NdFeB magnets, Journal of Alloys and Compounds,
478 (2009), 188–192
6 J. Xiao, J. Otaigbe, Polymer-bonded magnets III. Effect of surface
modification and particle size on the improved oxidation and
corrosion resistance of magnetic rare earth fillers, Journal of Alloys
and Compounds, 309 (2000), 100–106
7 Q. Chen, J. Asuncion, J. Landi, B. M. Ma, The effect of the coupling
agent on the packing density and corrosion behavior of NdFeB and
SmCo bonded magnets, Journal of Applied Physics, 85 (1999),
5684–5686
8 D. N. Brown, Z. Chen, P. Guschl, P. Campbell, Developments with
melt spun RE–Fe–B powder for bonded magnets, Journal of Magne-
tism and Magnetic Materials, 303 (2006), 371–374
9 A. Gebert, M. Rada, A. Kirchner, J. Lyubina, O. Gutfleisch, L.
Schultz, Corrosion Behavior of NdFeB-Based Nanocrystalline
Permanent Magnets, Journal of Metastable and Nanocrystalline
Materials, 631 (2005), 631–634
10 A. A. El-Moneim, A. Geberta, M. Uhlemann, O. Gutfleisch, L.
Schultz, The influence of Co and Ga additions on the corrosion
behavior of nanocrystalline NdFeB magnets, Corrosion Science, 44
(2002), 1857–1874
11 S. N. B. Hodgsoa, C. G. Hoggarth, H. A. Davies, R. A. Buckley,
Protection of NdFeB magnets by ultra-thin sol gel derived films,
Journal of Materials Processing Technology, 92–93 (1999), 518–524
12 Q. Li, S. Y. Zhang, J. P. Wang, H. Gao, Process analysis of MgO film
on NdFeB magnet by sol–gel method, Surface Engineering, 25
(2009), 589–593
13 N. Sinnadurai, The correct model for and use of HAST, International
symposium on microelectronics, Boston MA, 4339 (2000), 733–736
D. SOJER et al.: ANALYSIS OF CORROSION PROPERTIES OF MELT SPUN Nd-Fe-B RIBBONS ...
228 Materiali in tehnologije / Materials and technology 47 (2013) 2, 223–228
N. [TREKELJ et al.: INFLUENCE OF THE MICROSTRUCTURE ON MACHINING A CENTRAL HOUSING ...
INFLUENCE OF THE MICROSTRUCTURE ON
MACHINING A CENTRAL HOUSING MADE OF
PEARLITE GREY CAST IRON
VPLIV MIKROSTRUKTURE NA OBDELOVALNOST
CENTRALNEGA OHI[JA IZ PERLITNE SIVE LITINE
Neva [trekelj, Milanka Nuni}, Iztok Nagli~, Bo{tjan Markoli
Faculty of Natural Sciences and Engineering, University of Ljubljana, A{ker~eva 12, 1000 Ljubljana, Slovenia
neva.strekelj@omm.ntf.uni-lj.si
Prejem rokopisa – received: 2012-07-31; sprejem za objavo – accepted for publication: 2012-10-03
This article presents the cause(s) of a relatively increased wear and failure of cutting tools during the final treatment of a central
housing made of pearlitic (grey) cast iron with lamellar graphite. Castings with a proper (designated as šgood’) and an
inadequate (designated as šbad’) microstructures were investigated. Chemical analyses showed a higher concentration of
carbide-forming elements in the bad casting, particularly at the edges. Vickers hardness was also measured and the results
indicated a higher hardness of bad castings. A microstructural analysis showed that the good casting had the targeted
microstructure of the pearlitic matrix and the type A graphite with the size of 4–6. In addition to the narrow, initially austenitic
zone that extended only by about 200 μm into the bad casting, steadite was observed, which adversely affected the properties.
Moreover, the shape of the graphite and its distribution were uneven, which was also reflected in a low machinability of the bad
casting. The inner regions of the castings included graphite of a suitable shape and size, while the edges showed that the
solidification of the alloy started by following a stable system with a solidification of the primary austenite and continued
according to a metastable one. In the bad casting this area occurred just below the surface, while in the good casting it stretched
into the interior. In the bad casting the type B graphite was mainly developed. The influence of the quality of the cutting tools
was not investigated.
Keywords: microstructure, grey cast iron, machining
Namen ~lanka je bil ugotoviti vzrok relativno pove~ane obrabe in lomljenja no`ev pri obdelavi centralnega ohi{ja iz perlitne sive
litine z lamelnim grafitom. Analiza je bila opravljena na ulitku s primerno (ozna~en kot šdober’ ulitek) in neprimerno (ozna~en
kot šslab’ ulitek) mikrostrukturo. Preverjena je bila kemijska analiza. Rezultati so pokazali vi{jo koncentracijo karbidotvornih
elementov v slabem ulitku, zlasti na robovih. Izmerjena je bila mikrotrdota po Vickersu, pri ~emer so nastopile te`ave zajemanja
grafita v posamezen vtisek, vendar pa se je izkazalo, da je slab{i ulitek tr{i od dobrega. Analiza mikrostrukture je pokazala, da
ima dober ulitek predpisano mikrostrukturo iz perlitne osnove in grafita tipa A, velikosti 4–6. V slabem ulitku je poleg ozke
izhodne avstenitne cone, ki se razteza v notranjost le pribli`no 200 μm, opa`ena {e prisotnost steadita, ki neugodno vpliva na
lastnosti sive litine. Poleg tega je oblika grafita in njegova porazdelitev neenakomerna, kar se tudi izra`a kot slab{a
obdelovalnost ulitka. V notranjosti ulitkov sta bili velikost in oblika grafita primerni, medtem ko je bilo na robovih opaziti, da se
je strjevanje zlitine za~elo po stabilnem sistemu z izlo~anjem primarnega avstenita, nadaljevalo pa po metastabilnem. Pri slabem
ulitku je to obmo~je segalo tik pod povr{ino, pri dobrem pa bolj v notranjost. V slabem ulitku se je grafit razvijal ve~inoma z
obliko B. V okviru raziskav vpliv kvalitete rezilnega orodja ni bil obravnavan.
Klju~ne besede: mikrostruktura, siva litina, kon~na strojna obdelava
1 INTRODUCTION
A central housing is a casting made of grey cast iron
with lamellar graphite in the pearlitic matrix. It is the
essential supporting component for a turbocharger,
which is under heavy loads so the accuracy and quality
of the final machining are essential. The final mechanical
machining was carried out using common practice in
casting from two different foundries.
The grey cast iron with lamellar graphite is an
iron-based alloy with certain amounts of carbon and
silicon.1 It can also contain manganese, phosphorus and
sulphur. The toughness and tensile strength of such a
material are usually lower because of the graphite
lamellae intersecting the metallic matrix, which may also
cause a notch effect. Mechanical properties depend
heavily on the quantity, size, shape and distribution of
the graphite particles.2–4 The notch effect is more pro-
nounced if these are larger and vice versa.
Moreover, alloying elements can affect the machi-
nability of grey cast iron with lamellar graphite, as grey
cast iron with fine lamellae (which is due to the additions
of modifying alloying elements) is very hard and has
high strength3. In practice it is found that the greatest
difficulty in the final machining of grey cast iron is the
presence of the so-called hard spots.
Therefore, a comparison of the two castings of the
central housing from two different foundries was
performed in order to establish the effect of the
microstructure on machinability. The reasons for the
poor machinability of the castings can also be the cutting
tools themselves, but this was not studied in this
investigation.
The castings supplied by the first foundry were made
with the Croning cast procedure, while in the second
case the castings were cast into a bentonite-clay-mixture
mould.
Materiali in tehnologije / Materials and technology 47 (2013) 2, 229–234 229
UDK 669.131.6 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 47(2)229(2013)
For a reliable identification of the constitution of grey
cast-iron castings and revealing the cause(s) for the poor
machinability, it is important to calculate the values of
certain indicators of microstructural features. In our case
we calculated the following ones: the graphitization fac-
tor (k), the amount of eutectic graphite (AEG), the degree
of saturation (SC) and the carbon equivalent (CE). For
calculation purposes the data on chemical compositions
were used as supplied by the foundries.
1.1 Calculation of the graphitization factor k (degree
of graphitization)
In addition to the slow cooling rates, the cast iron
should have a chemical composition ensuring a suffi-
ciently high tendency to form graphite or a tendency
towards graphitization. This criterion is met if the
graphitization factor k is properly adjusted in relation to
the cooling rate or wall thickness of the castings as
expressed in the equation (1)5,6:
k = −
+
⎛
⎝
⎜ ⎞
⎠
⎟4
3
1
5
3
Si
C Si
(1)
1.2 Calculation of the amount of eutectic graphite
(AEG)
AEG = Ctotal – 2,0 + 0,1(Si + P) (2)
The amount of eutectic graphite (equation (2)6)
greatly influences the properties of grey cast iron (in
addition to the state of the matrix – ferrite and pearlite or
pearlite only7). Because of its physicochemical
properties graphite has a strong, favourable influence in
terms of tribology, namely, it reduces the friction and
acts as a lubricant for the cutting tools.
1.3 Calculation of the degree of saturation (SC)
The influence of individual elements on a eutectic
composition is expressed with the degree of saturation,
which is given as follows (Pfannenschmidt8):
S C = − − +
C
4,23 0,312Si 0,330P 0,066Mn
(3)
1.4 Calculation of the carbon equivalent (CE)
Since grey cast iron contains the chemical elements
that promote a formation of graphite instead of cemen-
tite, influencing the amount of carbon developed in the
form of graphite, it is essential to use an equivalent
amount of carbon in a Fe-C system according to the
equation (4)5:
CE = % C + 0,30 % Si + 0,33 % P (4)
2 EXPERIMENTAL WORK
In order to identify the cause(s) for the problems
arising during the final machining of the castings (the
good and the bad one), the following methods were used:
• Light optical microscopy (LOM);
• Chemical analyses via a mass spectroscopy – the
results obtained from the foundries were verified;
• Microhardness measurements via Vickers hardness
and
• Size, type and distribution of the graphite determi-
nation.
According to the given geometry and the final
machining process, critical points on the castings were
selected and samples were cut out using a water-cooled
circular saw (Figure 1). The specimens were metallo-
graphically prepared and were suitable for determining
the shape, size, type and distribution of the graphite
without etching. Furthermore, the samples were etched
(2 % nital), so that the constitution of the matrix was
revealed along with the presence of the other micro-
structural constituents. The samples prepared in this way
were suitable for the microhardness measurements and
an investigation using LOM.
The microhardness measurements using the Vickers
method were carried out on one good and two bad
castings, using a Shimadzu microhardness tester. The
used loads were of 100 g with the loading times of 10 s.
The method of LOM using a ZEISS Axio Imager.
A1m microscope gave an insight into the size, shape and
distribution of the microstructural components.
The procedure for determining the shape, size and
type of graphite followed the EN ISO 945: 1994 instruc-
tions.9
N. [TREKELJ et al.: INFLUENCE OF THE MICROSTRUCTURE ON MACHINING A CENTRAL HOUSING ...
230 Materiali in tehnologije / Materials and technology 47 (2013) 2, 229–234
Figure 1: Samples of the: a) good casting and b) bad casting
Slika 1: Vzorci: a) dobrega ulitka in b) slabega ulitka
3 RESULTS AND DISCUSSION
3.1 Calculations
For a reliable identification of the constitution of grey
cast-iron castings and determination of the cause(s) for a
difficult machining of a product, the values of certain
indicators of microstructural features had to be assessed.
In both cases the graphitization factor was greater
than 1, i.e., 1.713 and 1.528, which confirmed that the
cast iron in our case was grey cast iron. In the good
casting the graphitization factor was slightly higher than
in the bad one. The values suggested that lamellar
graphite should be one of the constituents of a micro-
structure. It was obvious that the bad casting had a
smaller amount of eutectic graphite (1.566) than the
good one (1.712), which, of course, caused poorer
machinability of the bad casting. The values of the
degree of saturation for both castings clearly indicated
that microstructural differences in both castings were to
be expected. In both castings the alloy was hypoeutectic,
since the Sc was less than one in both cases (0.970 and
0.920). When the undercooling of the melt is taken into
account (common in the foundry operations) the solidi-
fication starts with primary crystallization of both
graphite and austenite. With a lower value of the SC the
amount of austenite in the microstructure was higher,
thus, when the cooling was fast a higher quantity of
pearlite was obtained, which increased the wear of the
cutting tools. The values of the carbon equivalent were
4.165 in the good casting and 3.980 in the bad casting
indicating that the bad casting is more hypoeutectic than
the good one. These calculations alone do not provide
sufficient grounds for predicting the casting behaviour
during the final machining. Therefore, the measurements
of microhardness, chemical analyses and a microstruc-
tural characterization of the castings were necessary.
3.2 Microhardness
Microhardness measurements for one good and two
bad castings were performed. Figure 2 shows the points
of the microhardness measurements on individual
castings. Points 1–6 (Figure 2a) on the good casting
represent the edges where the process of the machining
was carried out. Points 7 and 8 were in the inner region
of the good casting. It should be noted that the imprints
of the microhardness measurements very often included
graphite. This is not in accordance with the principles of
reliable hardness measurements, therefore the results are
somewhat compromised. But for the evaluation of the
casting machinability the hardness of the whole casting
is important and, thus, the results clearly indicated why
the machining of a bad casting was more difficult and the
cutting-tool wear greater.
The microhardness on the edges of the good casting
(points 1–5, Figure 2a) was 160–260 HV, while for the
bad casting (bad 1) it was higher, as high as 350 HV
(point 7, Figure 2b). Similar values were obtained with
the microhardness measurements for the other bad
casting (not in Figure 2), where the microhardness of the
matrix was 335 HV. The results of the hardness
measurements are graphically presented in Figure 3.
Apparently the highest microhardness was on the
edges of all the castings, but for the bad castings the
values were significantly higher. The difference between
the maximum hardness of the bad and the good casting
was almost 100 HV, which clearly indicated a higher
probability of an increased wear of the cutting tools.
For the clarification of the reasons for the increased
tool wear the microhardness measurements were not
sufficient. Therefore, the data on the chemical compo-
sitions of the castings from the foundries were reviewed.
3.3 Chemical composition of the castings
The results of the microchemical analyses of the
good and bad castings were compared considering the
foundry data and are summarized in Table 1. Chemical
analyses revealed small differences between the bad and
the good castings in the contents of carbon, silicon and
N. [TREKELJ et al.: INFLUENCE OF THE MICROSTRUCTURE ON MACHINING A CENTRAL HOUSING ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 229–234 231
Figure 3: Microhardness measurements where the points of measu-
rements correspond to the numbers in Figure 2 for the two bad and
one good casting
Slika 3: Mesta meritev mikrotrdote, ki odgovarjajo {tevilkam na sliki
2 v dveh slabih in enem dobrem ulitku
Figure 2: Areas of microhardness measurements: a) good casting, b)
bad casting
Slika 2: Podro~ja meritev mikrotrdote: a) dobrega ulitka, b) slabega
ulitka
manganese (Table 1). The good casting had higher
amounts of carbon and silicon, while the content of
manganese (as well as chrome, sulphur and phosphorus)
was lower. This explained, to a great extent, the diffe-
rence in the machining between the good and the bad
castings.
Table 1: Results of the microchemical analyses carried out in the
foundries in mass fractions (w/%)
Tabela 1: Rezultati mikrokemijskih analiz iz livarn v masnih dele`ih
(w/%)
Element Good casting (w/%) Bad casting (w/%)
C 3.49 3.36
Si 2.13 1.96
Mn 0.51 0.79
Cr 0.06 0.07
P 0.45 0.11
S 0.05 0.06
Lower contents of silicon and carbon with increased
levels of manganese meant a higher propensity to white
solidification (in the bad casting) and, thus, a higher
amount of pearlite. Chromium can also contribute to this
by stabilizing and forming carbides, thus increasing the
casting hardness. Furthermore, phosphorus formed hard
microstructural constituents, for example Fe3P, which
appeared within the ternary eutectic (Fe+Fe3C+Fe3P)
called steadite. Ultimately, the casting parameters and
casting geometry can influence the homogeneity of a
microstructure.
3.4 Metallographic analysis
In this part of the investigation the form and the
distribution of the graphite was analysed first.
Based on the images of the good casting (Figures 4a,
b) graphite was mainly of form A and size 4–6 and its
skeleton showed a better contiguity than the one in the
bad casting (Figures 5a, b). In the bad casting the
graphite skeleton was frequently interrupted by
austenitic areas, which were also more extensive. This
caused a larger dynamic load for the cutting tools
because of a great number of the transitions from the
areas with a better machinability into the areas with a
poorer machinability.
The constitution of the castings after the etching was
determined also by using LOM. An image of the edge of
the good casting (Figure 6a) showed a ferrite-graphite
area, which had extended into the casting interior, where
the matrix was slowly becoming pearlitic. It was obvious
that the solidification of the casting was conducted
separately with the primary solidification of the austenite
on the edge of the casting and eutectic crystallization of
the graphitic eutectic. Then the consecutive evolution of
the microstructure followed the constitution of the
Fe-Fe3C system and the microstructure of the casting
N. [TREKELJ et al.: INFLUENCE OF THE MICROSTRUCTURE ON MACHINING A CENTRAL HOUSING ...
232 Materiali in tehnologije / Materials and technology 47 (2013) 2, 229–234
Figure 4: LOM images of the distribution and shape of graphite A,
4–6 in size, in the good casting
Slika 4: Posnetka porazdelitve in oblike grafita A, velikosti 4–6 v
dobrem ulitku
Figure 5: LOM images of the distribution and shape of graphite,
showing larger graphite lamellae and a poorer contiguity of the gra-
phite skeleton in the bad casting
Slika 5: Posnetka porazdelitve in oblike grafita, ki ka`eta ve~je gra-
fitne lamele in slab{o kontiguiteto grafitnega skeleta v slabem ulitku
a)
b)
a)
b)
interior consisted of fine lamellar pearlite, interrupted by
graphite lamellae (A5) (Figure 6b).
Furthermore, the ferrite-graphite zone situated mostly
on the edge of the casting extended into the casting
interior to a depth of 1200–1500 μm. This was exactly
the depth to which the cutting tools reached during the
final machining. According to the hardness measure-
ments of the matrix on the edge of the casting and the
presence of the graphite skeleton, it is clear that the good
casting exhibited a better machinability and caused a
lower wear of the cutting tools.
Metallographic analyses of the bad casting showed a
pearlite matrix and graphite (A–B and 4–5), but also the
presence of steadite (Figure 7). In addition to steadite,
the form of the graphite was less favourable in terms of
the final machinability. The graphite skeleton also has a
lower contiguity.
When comparing the depths of the ferrite zones in the
good (Figure 8a) and bad castings (Figure 8b), it was
obvious that the latter had a less profound ferrite zone.
The depth of this area was only 200 μm in the case of
the bad casting, while in the good one it was
1200–1500 μm. Thus, during the final machining of
the bad casting, the cutting tool quickly reached the hard
pearlite matrix, which also contained extremely hard and
brittle particles of steadite. Because of the coarser and
more unevenly distributed graphite in the casting, the
wear of the cutting tools was even higher.
N. [TREKELJ et al.: INFLUENCE OF THE MICROSTRUCTURE ON MACHINING A CENTRAL HOUSING ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 229–234 233
Figure 8: LOM images of: a) the good-casting edge, where the ferrite
zone was 1200–1500 μm thick and b) the bad-casting edge, where
the ferrite zone was only 200 μm thick
Slika 8: Posnetka mikrostrukture: a) roba dobrega ulitka, feritno
obmo~je sega 1200–1500 μm globoko v vzorec in b) roba slabega
ulitka, feritno obmo~je sega le 200 μm globoko v vzorec
Figure 7: LOM image of the bad-casting edge, showing the pearlite
matrix, graphite A-B, 4-5 in size, and steadite
Slika 7: Posnetek mikrostrukture roba slabega ulitka, ki ka`e perlitno
osnovo, grafit A-B, velikosti 4–5, in steadit
Figure 6: LOM images of the good casting: a) on the edge, showing a
ferrite-graphite area and b) in the centre, showing fine lamellar
pearlite and graphite lamellae
Slika 6: Posnetka mikrostrukture dobrega ulitka: a) na robu, ki ka`e
feritno-grafitno obmo~je, in b) v sredini, ki ka`e lamelarni perlit in
lamele grafita
4 CONCLUSIONS
In this article the possible causes for a poorer final
machinability of grey cast-iron castings with lamellar
graphite are discussed and the potential reasons for the
increased wear and failures of the cutting tools were
presented. The conclusions can be summarized as
follows:
• Established calculations confirmed that the inve-
stigated material was hypoeutectic grey cast iron with
lamellar graphite. The calculations were not abso-
lutely accurate and provided only the information
about which type of grey cast iron was most likely to
be formed.
• The hardness of grey cast iron with lamellar graphite
was measured via Vickers microhardness, because it
enables the measurements of the hardness of indivi-
dual microconstituents and the matrix. The hardness
of bad castings was higher than that of the good
casting.
• Chemical analyses showed higher contents of manga-
nese and phosphorus in the bad casting and higher
silicon levels in the good casting.
• The casting edges of the polished samples showed a
dendritic morphology, which was a result of the
solidification sequence with primary austenite. Then
it transformed into ferrite and graphite and/or into
pearlite. In the casting interior we identified the size
and shape A of graphite. In the bad casting a larger
quantity of type-B graphite with fine lamellar gra-
phite and uneven distribution of graphite was
observed.
• The depths of the transformed primary austenite in
the bad and good castings were compared and clearly
the depth in the good casting was greater.
• Microstructural analyses showed a presence of gra-
phite, pearlite and, on the edges of the casting, also
the undercooled D graphite and ferrite. The bad cast-
ing also contained a significant amount of steadite,
which was one of the main reasons for the difficult
final machining of the central housing.
• The main reason for a poor machinability and greater
wear of the cutting tools lies in the constitution of the
microstructure. The nucleation of austenitic primary
crystals and of the graphite from the eutectic is not
optimal or it is even harmful. It is obvious that the
transformed primary austenitic particles are relatively
large. Undercooled forms of graphite were also ob-
served.
• The influence of the cutting-tool quality should be
taken into account when the final and definitive
reason for the breakage and wear of cutting knives is
determined.
5 REFERENCES
1 Y. Boran, International Iron & Steel Symposium, Karabük, Turkey,
2012
2 T. Alp, A. A. Wazzan, F. Yilmaz, Arabian Journal for Science and
Engineering, 30 (2005), 163–175
3 H. Nakae, H. Shin, Materials Transactions, 42 (2001), 1428–1434
4 D. J. Celentano, M. A. Cruchaga, B. J. Schulz, International Journal
of Cast Metals Research, 18 (2005), 237–247
5 D. M. Stefanescu, ASM Handbook, volume 15, Casting, ASM Inter-
national, 1998
6 M. Nuni}, Influence of microstructure on machining of central
housing made of pearlite CI, Diploma work, Ljubljana, 2011 (in
Slovenian)
7 P. Mrvar, M. Petri~, J. Medved, Key Engineering Materials, 457
(2010), 163–168
8 Dru{tvo livarjev LR Slovenije, @epni livarski priro~nik, Ljubljana,
Zalo`ba Litostroj, 1960 (in Slovenian)
9 Available from World Wide Web: http://www.atilim.edu.tr/~ktur/
mate401/Dosyalar/26-ELKEM_poster-graphite%20structures%20
in%20cast%20irons.pdf
N. [TREKELJ et al.: INFLUENCE OF THE MICROSTRUCTURE ON MACHINING A CENTRAL HOUSING ...
234 Materiali in tehnologije / Materials and technology 47 (2013) 2, 229–234
M. ALIZADEH et al.: THE WET-CHEMICAL SYNTHESIS OF FUNCTIONALIZED Zn1-xOMnx QUANTUM DOTS ...
THE WET-CHEMICAL SYNTHESIS OF
FUNCTIONALIZED Zn1–xOMnx QUANTUM DOTS
UTILIZABLE IN OPTICAL BIOSENSORS
MOKRA KEMIJSKA SINTEZA FUNKCIONALIZIRANIH
KVANTNIH DELCEV Zn1–xOMnx, UPORABNIH V OPTI^NIH
BIOSENZORJIH
Mohamad Alizadeh1, Reza Salimi1,2, Hassan Sameie1,2, Ali Asghar Sarabi1,
Ali Asghar Sabbagh Alvani2, Mohammad Reza Tahriri3
1Amirkabir University of Technology, Faculty of Polymer Engineering & Color Tech., 424 Hafez Ave., 15875-4413 Tehran, Iran
2Amirkabir University of Technology, Color and Polymer Research Center (CPRC), 424 Hafez Ave., 15875-4413 Tehran, Iran
3Amirkabir University of Technology, Biomaterials Group, Faculty of Biomedical Engineering, P.O. Box, 15875-4413 Tehran, Iran
sabbagh_alvani@aut.ac.ir
Prejem rokopisa – received: 2012-08-09; sprejem za objavo – accepted for publication: 2012-10-03
ZnO quantum dots (QDs) were successfully synthesized via the precipitation technique, and the effect of the Mn2+ ion
concentration as dopant on the optical properties was studied. In order to control the particle size by limiting the growth of the
particles after the nucleation and to provide a side group on the surface, which can be further conjugated to bio-cells,
polyethylene glycol (PEG) was used as a bio-compatible capping agent. XRD analyses revealed a single-phase ZnO wurtzite
crystal structure. A TEM micrograph illustrates that the final QDs were about 15 nm in diameter and had a spherical shape.
Also, the organic groups on the surface of nano-particles (NPs) were characterized with Fourier transform infrared spectroscopy
(FTIR). It is clear from the photo-luminescence spectra that doping Mn2+ ions in the host lattice led to an appearance of a new
visible emission band in the range of 410–450 nm because of the 4T1  6A1 transition. The results show that the final QDs have a
potential for biochemical optical sensing.
Keywords: quantum dots, surface modification, biosensor, photoluminescence
Kvantni delci (QD) ZnO so bili uspe{no sintetizirani s tehniko izlo~anja in preu~evan je bil u~inek koncentracije ionov Mn2+ kot
dopantov na opti~ne lastnosti. Polietilen glikol (PEG) kot omejitveno sredstvo je bil uporabljen za kontrolo velikosti delcev z
omejevanjem rasti delcev po nukleaciji in za zagotovitev posebne skupine na povr{ini, ki jo je mogo~e naprej zdru`evati z
biocelicami. XRD-analize so odkrile enofazni ZnO s kristalno strukturo wurtzita. Posnetek s TEM poka`e, da so kon~ni QDs
okrogle oblike s premerom okrog 15 nm. Organske skupine na povr{ini nanodelcev (NPs) so bile dolo~ene s
FTIR-spektroskopijo. Iz fotoluminiscen~nega spektra izhaja, da dopiranje ionov Mn2+ v gostujo~o mre`o povzro~i prehod 4T1 
6A1 in zato pojav novega vidnega emisijskega pasu v obmo~ju 410–450 nm. Rezultati ka`ejo, da imajo kon~ni QDs potencial
biokemi~nega opti~nega zaznavanja.
Klju~ne besede: kvantni delci, modifikacija povr{ine, biosensor, fotoluminiscenca
1 INTRODUCTION
The synthesis of semiconductor quantum dots (QDs)
has received a great deal of interest because of their
potential for biomedical applications in imaging, drug
targeting and delivery. Compared with the conventional
organic fluorophores, QDs have significant advantages,
for example, a narrow and size-tunable emission spec-
trum, broadband excitation, a high resistance against
photo-bleaching and a good chemical stability1,2. Due to
their size-dependent properties and dimensional simila-
rities to biomolecules, these bio-conjugate QDs are well
suited as contrast agents for the in-vivo magnetic
resonance imaging (MRI) as carriers for drug delivery,
and as structural scaffolds for tissue engineering3. ZnO
nano-particles are environment friendly and have a wide
band-gap of 3.37 eV and a rather large exciton binding
energy, which makes the exciton state stable even at
room temperature. Therefore, wide band-gap semicon-
ductor nanocrystals can be doped with transition metal
ions4,5. An important section of synthesising NPs is
surface modification. Most surface modifications of NPs
for bio-imaging applications are based on chemisorption
since it offers a stronger and more robust bond and a
more stable surface ligand, compared with physisorption.
A successful conjugation of biomolecules to NPs
depends on a proper surface modification. Poly ethylene
glycol (PEG) has been commonly conjugated to various
drugs, liposomes, polymeric micelles and nanoparticles
to prolong their blood-circulation time by reducing the
nonspecific adsorption of proteins via a steric stabiliza-
tion effect6,7. In this paper, we report on a wet chemical
synthesis of Zn1–xOMnx QDs and functionalizing them
with PEG to achieve the appropriate particle size and
good optical properties used in biological labelling. The
main advantages of this process are the ability to control
the ZnO purity and the use of biocompatible capping
agents that provide triple functions: (1) to control the
particle size by limiting the growth of the particles after
the nucleation, (2) to provide a side group on a surface
Materiali in tehnologije / Materials and technology 47 (2013) 2, 235–237 235
UDK 620.3 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 47(2)235(2013)
that can be further conjugated to a bio-cell, (3) to
eliminate surface defects that would affect the optical
properties of ZnO nanocrystals.
2 EXPERIMENTAL WORK
Manganese acetate tetrahydrate (Mn(Ac)2·4H2O),
zinc acetate dihydrate (Zn(Ac)2·2.5H2O), polyethylene
glycol (Mw = 400 g/mol), extra pure water, ethanol and
hydroxide sodium were used as the starting materials.
We have developed a room-temperature technique for
synthesising semiconductor nanocrystals that employs a
capping agent to control both the size and the shape. The
manganese precursor (Mn(OAc)2·4H2O) dissolved in
1 ml of water and then 100 ml of ethanol was added to
the solution. After stirring the solution for 2 h,
Zn(OAc)2·2.5H2O was added; then the temperature
raised to 50 °C and the solution was stirred again. After
that the obtained solution was quenched in ice, and the
required PEG was added as a capping agent. The mixture
was again stirred for 3 h and then hydrolyzed by adding
NaOH under ultrasound for 2 h. The solvents were then
removed with the rotation vaporization. Eventually, the
resulting mixture was washed with water leading to a
precipitation of Zn1–xMnxO nanocrystals. The resulting
powders were analyzed with X-ray diffraction (XRD;
Bruker AXS: D8 Advance), transmission electron micro-
scopy (TEM; a Hitachi H-800 electron microscope) and
fluorescence spectrophotometer (Perkin-Elmer LS-55,
the exciting source: a near-ultraviolet-Xenon lamp) and
Fourier transform infrared spectroscopy (Equninox-55).
3 RESULTS AND DISCUSSION
ZnO colloids were prepared with the precipitation
from the solution using Zn(CH3CO2)2 and NaOH. The
overall reaction of the synthesis of ZnO nanoparticles
from Zn(II) acetate can be written as:
Zn(CH3CO2)2 + 2NaOH 
 ZnO + 2Na(CH3CO2)2 + H2O (1)
The XRD pattern of the ZnO nanocrystals synthe-
sized via precipitation is shown in Figure 1. The main
phase can be indexed to the pure phase of zinc oxide
according to JCPDS 80-0075. Also, the wurtzite struc-
ture of the ZnO lattice of a nanocrystalline sample can be
confirmed.
From the FTIR spectrum presented in Figure 2a, it is
obvious that PEG properly caps the surface of the QDs.
The main peaks can be related to different groups as
follows: 1) the peaks around 3670 cm–1 and 3407 cm–1
are due to the O-H stretching; 2) the peaks around (2974,
2916 and 1241) cm–1 are attributed to the C-H stretching
from the -CH3 alkanes, the C-H stretching from the
-CH2- alkanes and the R-O-R stretching (the conjugate
ether), respectively. In addition, the TEM image of the
ZnO:Mn QDs caped with PEG is illustrated in Figure
2b. It is clearly observed that nanoparticles are spherical
in shape and their average size is approximately 10 nm.
The PL spectrum of the Mn-doped ZnO d-dots with
different doping concentrations (from mol fraction 0–7
%) are shown in Figure 3(a–d). It is worth mentioning
that by varying the Mn doping concentration, the relative
PL intensities of the dual-colour emissions can be well
manipulated. The UV emission (356 nm) is characte-
ristics of the ZnO host and usually influenced by the
M. ALIZADEH et al.: THE WET-CHEMICAL SYNTHESIS OF FUNCTIONALIZED Zn1-xOMnx QUANTUM DOTS ...
236 Materiali in tehnologije / Materials and technology 47 (2013) 2, 235–237
Figure 2: a) FTIR spectra of the PEG-capped ZnO nanoparticles and
b) a TEM image of 5 % Mn-doped ZnO nano-particles
Slika 2: a) FTIR-spekter s PEG omejevanih ZnO-nanodelcev in b)
TEM-posnetek s 5 % Mn dopiranih ZnO-nanodelcev
Figure 1: XRD pattern of ZnO prepared with precipitation
Slika 1: Rentgenski pra{kovni posnetek ZnO, pridobljen z izlo~anjem
doping situation. Also, the blue emission band, centred
at about 426 nm, is attributed to the transition between
the 4T1  6A1 energy levels of the Mn2+ 3d states. A
schematic representation of the emission levels of the
Mn-doped ZnO d-dots is briefly illustrated in Figure 3e.
When an intrinsic q-dot is excited by the photons with
the energy higher than its band gap, an exciton (an
electron-hole pair) will be generated. A direct recom-
bination of an electron-hole pair, typically being
quantum-confined in the case of nanocrystals, gives the
well-known band-edge emission, or exciton emission.
However, the emission of a d-dot is fundamentally
different. After the exciton is generated by the absorption
of the host semiconductor nanocrystal, the energy of a
photogenerated electron and a hole pair will be
transferred into the electronic levels of the Mn2+ ions.
The recombination in a Mn2+ ion center leads to the
characteristic dopant emission from the Mn2+ ion,
namely the 4T1 to 6A1 transition.
4 CONCLUSION
In this research, ZnO:Mn QDs have been success-
fully synthesized with a precipitation synthesis and then
the surface of nanoparticles was covered with hydroxyl
groups. In the meantime, the effect of the dopant con-
centration on the optical characteristics was investigated
in detail. According to experimental results, the func-
tionalized QDs can be used as a core of optical bio-
sensors after conjugating with the desired anti-bodies.
Acknowledgment
The authors would like to acknowledge the Iranian
Nano Technology Initiative Council for supporting the
research.
5 REFERENCES
1 X. Gao, W. Chan, S. Nie, Quantum-dot nanocrystals for ultra-
sensitive biological labelling and multi color optical encoding, J.
Biomed. Opt., 7 (2002), 532–537
2 T. Jamieson, R. Bakhshi, D. Petrova, R. Pocock, Biological appli-
cations of quantum dots, Biomater., 28 (2007), 4717–4732
3 A. Smith, H. Duan, A. Mohs, S. Nie, Bioconjugated quantum dots
for invivo molecular and cellular imaging, Adv. Drug Deliv. Rev., 60
(2008), 1226–1240
4 Y. Yang, Y. Jin, H. He, Q. Wang, Y. Tu, H. Lu, Z. Ye, Dopant-
Induced Shape Evolution of Colloidal Nanocrystals: The Case of
Zinc Oxide, J. Am. Chem. Soc., 132 (2010), 13381–13394
5 E. Badaeva, C. M. Isborn, Y. Feng, S. T. Ochsenbein, D. R. Gamelin,
X. Li, Theoretical Characterization of Electronic Transitionsin Co2+
and Mn2+ Doped ZnO Nanocrystals, J. Phys. Chem. C, 113 (2009),
8710–8717
6 H. Mok, K. H. Bae, C. H. Ahn, T. G. Park, PEGylated and MMP-2
Specifically DePEGylated Quantum Dots: Comparative Evaluation
of Cellular Uptake, Langmuir, 25 (2009), 1645–1650
7 R. Gref, Y. Minamitake, M. T. Peracchia, V. Trubetskoy, V. Torchilin,
R. Langer, Science, 263 (1994), 1600–1603
M. ALIZADEH et al.: THE WET-CHEMICAL SYNTHESIS OF FUNCTIONALIZED Zn1-xOMnx QUANTUM DOTS ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 235–237 237
Figure 3: PL emission spectra of the Zn1–xOMnx QDs with different
doping concentrations: a) x = 0 %, b) x = 2 %, c) x = 5 %, d) x = 7 %
excited at 325 nm and e) a schematic representation of the emission
levels of the Mn-doped ZnO d-dots
Slika 3: Spekter PL-emisij Zn1–xOMnx QD z razli~nimi koncentra-
cijami dopiranja: a) x = 0 %, b) x = 2 %, c) x = 5 %, d) x = 7 %
vzbujenih pri 325 nm in e) shemati~en prikaz nivojev emisij z Mn
dopiranega ZnO d-delcev

M. CAGALA et al.: INFLUENCE OF ALUMINIUM-ALLOY REMELTING ON THE STRUCTURE AND ...
INFLUENCE OF ALUMINIUM-ALLOY REMELTING ON
THE STRUCTURE AND MECHANICAL PROPERTIES
VPLIV VE^KRATNEGA PRETALJEVANJA ALUMINIJEVIH ZLITIN
NA STRUKTURO IN MEHANSKE LASTNOSTI
Michal Cagala, Marek Bøuska, Petr Lichý, Jaroslav Beòo, Nikol [pirutová
V[B-Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of Metallurgy and Foundry, 17. Listopadu
15/2172, Ostrava – Poruba, Czech Republic
michal.cagala@vsb.cz
prejem rokopisa – received: 2012-08-31; sprejem za objavo – accepted for publication: 2012-10-02
The aim of this work was to assess the repeated-remelting influence upon the mechanical properties, thermomechanical
properties, chemical composition and structure changes of the selected material. An Al-Cu-type aluminum alloy was chosen on
the basis of the ever increasing experiments with non-ferrous metals in the industry. The technical nomenclature of the selected
alloy is RR.350 according to the German standard ALUFOND 60. The RR.350 alloy is known for its poor foundry properties
which deteriorate due to remelting and affect mechanical properties and the cast-material structure. This negative influence upon
the structure and usable properties of a re-melted alloy is further confirmed in the submitted paper. The samples for
tensile-strength determination were cast into a metal mould. The gating system and the riser, which served as a charge for the
second melt, were removed from the casting. In this way we re-melted the material four times. The samples were machined and
ruptured within the temperature range between 20 °C and 350 °C. A sample for metallography and hardness determination
(HBS) was taken from the cast material. It can be seen in the tensile-strength diagram that the mechanical properties of the first
melt are higher, by 11 % at the temperature of 20 °C, than the properties of the third melt. This difference is evident up to 100
°C. At the temperatures of above 100 °C the cast-material strength characteristics are the same. This tendency shows itself on all
the materials tested so far. The hardness and microhardness evaluations show that the material reaches the highest values with
the fourth melt. This phenomenon is attributed to the repeated reoxidation and exclusion of oxide membranes. Further, the
material structure properties and chemical-composition change were evaluated. The results of this study confirmed a negative
influence of alloy remelting upon the material properties and structure.
Keywords: aluminium alloys, metallographic analysis, microstructures, thermomechanical properties, remelting, mechanical
properties
Namen tega dela je bil oceniti ponavljanje pretaljevanja na mehanske lastnosti, termomehanske lastnosti, kemijsko sestavo in
spremembe mikrostrukture izbranega materiala. Izbrana je bila vrsta aluminijeve zlitine Al-Cu na podlagi nara{~anja preizkusov
na ne`eleznih kovinah v industriji. Tehni~na oznaka izbrane zlitine, skladno z nem{kim standardom ALUFOND 60, je RR.350.
Zlitina RR.350 je poznana zaradi slabih livarskih lastnosti, ki se s pretaljevanjem poslab{ujejo in vplivajo na spremembe
mehanskih lastnosti in mikrostrukture v litem stanju. V predstavljenem ~lanku je potrjen negativni vpliv pretaljevanja na
uporabne lastnosti. Vzorci za natezne preizkuse so bili uliti v kovinsko kokilo. Ulivni in napajalni sistem, ki se je uporabljal za
sekundarno napajanje, je bil odstranjen iz ulitka. Tako je bil material {tirikrat pretaljen. Izdelani vzorci so bili poru{eni v
temperaturnem intervalu med 20 °C in 350 °C. Iz ulitega materiala so bili odrezani vzorci za metalografijo in za meritve trdote
(HBS). Iz nateznih diagramov je razvidno, da ima prva talina za okrog 11 % vi{jo natezno trdnost pri 20 °C v primerjavi s tretjo
talino. Ta razlika se opazi do temperature 100 °C. Pri temperaturah nad 100 °C so trdnostne lastnosti materiala v litem stanju
enake. Ta tendenca se je pokazala pri vseh do sedaj preizku{enih materialih. Primerjava trdote in mikrotrdote ka`eta, da je
najvi{jo vrednost dosegel material ~etrte taline. Ta pojav se pripisuje ponovljeni reoksidaciji in odsotnosti oksidnih ko`ic.
Ocenjene so bile tudi zna~ilnosti mikrostrukture in dolo~ene kemijske sestave. Rezultati teh preiskav so potrdili negativen vpliv
ve~kratnega pretaljevanja na mikrostrukturo in lastnosti materiala.
Klju~ne besede: aluminijeve zlitine, metalografska analiza, mikrostrukture, termomehanske lastnosti, pretaljevanje, mehanske
lastnosti
1 INTRODUCTION
At present the significance of non-ferrous metal
alloys as structural materials is growing thanks to their
good mechanical and thermomechanical properties, low
specific weight and heat-treatment possibility enhancing
usable properties of the alloy. By virtue of the afore-
mentioned properties, non-ferrous metal alloys replace
iron-based alloys in various industry branches. Non-
ferrous metals utilization in various industry branches
will have an increasing trend in the future. Considering
the rising prices of the materials, the enterprises engaged
in the processing of alloys and in the manufacture of
castings seek for a production-cost reduction. This trend
has been intensified due to the contemporary economic
crisis. One saving way is the utilization of recycled
material or remelting of foundry returns.
Some of the most widely used non-ferrous metal
alloys are aluminum-based alloys. In this paper we have
focused on the Al-Cu alloy RR.350. Due to its favour-
able properties, low weight, minimal dilatation and
capability to resist high temperatures up to 350 °C, the
RR.350 alloy is used for the thermally stressed castings
and the castings subjected to higher pressures1. This
alloy has an extensive usability not only in the automo-
tive industry. The RR.350 alloy is used in heat-treatment
conditions when its usable properties reach the highest
values. This study is focused on an evaluation of the
Materiali in tehnologije / Materials and technology 47 (2013) 2, 239–243 239
UDK 669.715:620.18:620.17 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 47(2)239(2013)
mechanical and thermomechanical properties of the
above-mentioned alloy after multiple remelting in
casting conditions. After the melting and casting of the
samples, a gating system, which then served as a new
charge, was cut off. These steps were performed four
times and then the material changes after remelting were
assessed. The research team focused on the changes in
following properties: tensile strength, hardness, micro-
hardness, alloy chemical composition and alloy structure
after remelting. The research took place in the V[B –
Technical University Ostrava laboratories.
2 MATERIALS AND METHODS
To observe the re-melting of non-ferrous metals, an
Al-Cu alloy with the technical nomenclature of RR.350
was chosen. The chemical composition of the alloy used
for the experiment is shown in Table 1.
Table 1: Chemical composition of the RR.350 alloy (in mass
fractions, w/%)
Tabela 1: Kemijska sestava zlitine RR.350 (v masnih dele`ih, w/%)
Fe Cu Mn Mg Ni Zn
0.368 4.731 0.310 0.034 1.937 0.126
Ti Pb Sn Co Cr Al
0.117 0.011 0.042 0.203 0.007 92.115
The original alloy with the as-delivered chemical
composition was melted in an electric resistance furnace
in a graphite-fireclay crucible and marked as I. melt. For
the casting of the test bars, a metal mould equipped with
a silicon sprayed-on coat preventing the sticking of the
casting to the mould was used. The mould consists of a
gating system with a bottom gate (for the laminar filling
of the mould), a test bar and a riser. The riser is over
dimensioned to eliminate defects, for instance the
shrinkage cavities. After having cast a batch of the test
bars marked as I. melt, the gating system with the riser,
which served as a charge for the next heat, was removed.
In this way we re-melted the original material four times.
The test bars (Figure 1) were machined and subjected to
a tensile test in the temperature range of 20 °C to 350 °C.
The measurement was performed on the INOVA
PRAHA tensile-testing machine. The heating of the
samples was performed in a resistance furnace with an
inert atmosphere. At each temperature, three samples
were ruptured and the numerical values of their diame-
ters were entered into the diagram.
For the hardness measurement, nine samples were
used and each of them was subjected to three incisions.
The sample preparation for the hardness testing (HB)
was performed in a metallographic laboratory. At first,
we cut off a one-centimeter-long cylinder from the cast
sample with the help of an emulsion-cooled saw. Then
the surface polishing on a horizontal water-cooled grin-
der followed using the emery cloths with a granularity of
360 up to 1.200. As a test specimen, a hardened ball of
2.5 mm in diameter was used. The test sample was
placed on a work plate and adjusted so that the indentor
could be pointed to the centre. Now the indentor arm was
moved to the vertical position and the arm movement
onto the sample was switched on. The machine carried
out the loading to the value of 306.5 N, with the holding
time of 10 s, followed by a release of the indentor.
The metallographic analysis was done with the help
of a GX 51 microscope, equipped with light polarization,
with a magnification of 12.5–1000.
The chemical analysis was performed on a GDS-
LECO spectrometer with the aid of the reference-mate-
rial calibration with guaranteed element content.
3 EXPERIMENTAL RESULTS AND DISCUSSION
The obtained results confirm a negative influence
upon the remelted alloy.
The measured results of the tensile tests (Figure 2)
show that the maximum strength of the original material
(I. melt) is reached at the temperatures of 20 °C and
100 °C. During the third remelting, the tensile strength
values at the temperatures of 20 °C and 100 °C
decreased by approximately 24 MPa, which was 11 %.
At the temperatures of over 100 °C, the strength values
M. CAGALA et al.: INFLUENCE OF ALUMINIUM-ALLOY REMELTING ON THE STRUCTURE AND ...
240 Materiali in tehnologije / Materials and technology 47 (2013) 2, 239–243
Figure 2: Dependence of tensile strength on temperature for different
melts
Slika 2: Odvisnost natezne trdnosti od temperature pri razli~nih
talinah
Figure 1: Test bar
Slika 1: Preizkusna palica
of both melts were the same. This tendency was similar
for the most of the tested materials2. The equation of the
values of the tensile strength at elevated temperatures
can be explained with the melting of low-melting
components in the material which could occur by
segregation in the material.
For this paper, the tensile-strength values of I. and III.
melts were selected. II. and IV. melts will be finished and
published in the next papers.
In both diagrams, the hardness values (Figure 3) and
the microhardness values (Figure 4) are the highest for
IV. melt. This tendency can be explained with the
repeated melt oxidation and the formation of oxide
membranes and intermetallic phases. The formation of
oxides and intermetallic phases increases the alloy
hardness and decreases the tensile strength owing to the
unsuitably excluded shapes causing the notch effects in
the material matrix3. Through their properties, aluminum
oxides and intermetallics have a negative effect on the
machining of the casting.
Multiple remelting does not affect only mechanical
properties, but it changes chemical composition and
structure morphology as well.
Comparing I. and II. melt (Tables 1 and 2) we can
see that weakening of the alloy occurred at II. melt on
account of decrease of some elements. The biggest
decrease was observed for Cu (0.9129 %), Mn (0.06921
%), Ni (0.2191 %), Mg (0.0213 %). Compared to this, Al
content was increased (1.318 %).
Table 2: Chemical analysis of II. melt (w/%)
Tabela 2: Kemijska sestava II. taline (w/%)
Fe Cu Mn Mg Ni Zn
0.234 3.8181 0.24079 0.0127 1.7179 0.0828
Ti Pb Sn Co Cr Al
0.172 0.00366 0.0141 0.271 0.000426 93.433
Table 3: Chemical analysis of III. melt (w/%)
Tabela 3: Kemijska sestava III. taline (w/%)
Fe Cu Mn Mg Ni Zn
0.251 3.7737 0.25548 0.0112 1.6621 0.0758
Ti Pb Sn Co Cr Al
0.159 0.00257 0.00947 0.529 0.000488 93.541
Comparing I. and III. melts (Tables 1 and 3) we can
see that a weakening of the alloy also occurred with III.
M. CAGALA et al.: INFLUENCE OF ALUMINIUM-ALLOY REMELTING ON THE STRUCTURE AND ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 239–243 241
Figure 6: I. melt – magnified 1000-times
Slika 6: I. talina – pove~ava 1000-kratna
Figure 4: Values of Micro-Hardness of individual melts
Slika 4: Mikrotrdota materiala posameznih talin
Figure 3: Values of Hardness of individual melts
Slika 3: Trdota materiala posameznih talin
Figure 5: I. melt – magnified 100-times
Slika 5: I. talina – pove~ava 100-kratna
melt due to a decrease in some elements. The biggest
decrease was again observed for Cu (0.9573 %), Mn
(0.05452 %), Ni (0.2749 %) and Mg (0.0228 %). On the
other hand, the Al content was increased (1.426 %).
Due to multiple remelting the material structure
changes its morphology as well (Figures 5 to 12). The
changes in the material structure are shown as grain
coarsening, unevenness of dendritic cells and dendrites4,
higher content of cavities and coarsening of intermetallic
phases.
M. CAGALA et al.: INFLUENCE OF ALUMINIUM-ALLOY REMELTING ON THE STRUCTURE AND ...
242 Materiali in tehnologije / Materials and technology 47 (2013) 2, 239–243
Figure 9: III. melt – magnified 100-times
Slika 9: III. talina – pove~ava 100-kratna
Figure 12: IV. melt – magnified 1000-times
Slika 12: IV. talina – pove~ava 1000-kratna
Figure 8: II. melt – magnified 1000-times
Slika 8: II. talina – pove~ava 1000-kratna
Figure 11: IV. melt – magnified 100-times
Slika 11: IV. talina – pove~ava 100-kratna
Figure 7: II. melt – magnified 100-times
Slika 7: II. talina – pove~ava 100-kratna
Figure 10: III. melt – magnified 1000-times
Slika 10: III. talina – pove~ava 1000-kratna
4 CONCLUSIONS
Properties of the RR.350 alloy after repeated remelt-
ing were observed. For each remelting, the basic material
properties were determined, i.e., tensile strength, hard-
ness, structure and chemical composition. The cast-mate-
rial tensile strength dropped by 11 % at the temperatures
of up to 100 °C. The hardness reached the highest values
in the case of IV. melt. The material structure change
influenced most of the above-mentioned changes. By
multiple remelting, a grain coarsening, unevenness of
dendritic cells, a higher content of cavities and coarsen-
ing of the intermetallic phases occur. The reason for a
decrease in the alloying elements is the re-melt loss in
the resistance furnace. The RR.350-alloy testing offers
other ways of mechanical-property enhancement. An
advisable option is a substitution of the elements
decreased due to multiple remelting and follow-up
inoculations with an AlTi5B1 master alloy5. Further, the
alloy can be very well hardened with the help of a heat
treatment (the strength values reach up to 300 MPa). As
mentioned above, the alloy is susceptible to various
kinds of cavities, therefore degassing before casting
would be advisable. For the manufacture of castings, we
can also recommend filtration to remove the inclusions
which decrease mechanical properties.
This paper was created within the project No.
CZ.1.05/2.1.00/01.0040 "Regional Materials Science and
Technology Centre", under the frame of the operation
programme "Research and Development for Innova-
tions" financed by the Structural Funds and by the state
budget of the Czech Republic.
5 REFERENCES
1 J. Rou~ka, Metalurgie ne`elezných slitin, VUT Brno, 2004
2 R. Koøený, Mo`nosti zvý{ení kvality vysokopevnostních a `áropev-
ných slévárenských slitin hliníku, V[B, Ostrava, 1991
3 F. Pí{ek, L. Jení~ek, P. Ry{, Nauka o materiálu I., Nauka o kovech 3.
svazek: Ne`elezné kovy, Academia, Praha 1973
4 L. Bäckerud, G. Chai, J. Taminen, Solidification Characteristic of
Aluminium Alloys, Foundry Alloys, Volume 2, ASF Skanaluminium,
Stockholm, 1990
5 D. Bolibruchová, E. Tillová, Zlievarenské zliatiny Al-Si, @U v
@ilinì, 2005
M. CAGALA et al.: INFLUENCE OF ALUMINIUM-ALLOY REMELTING ON THE STRUCTURE AND ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 239–243 243

C. S. ÇETINARSLAN, ALI GÜZEY: TENSILE PROPERTIES OF COLD-DRAWN LOW-CARBON STEEL WIRES ...
TENSILE PROPERTIES OF COLD-DRAWN
LOW-CARBON STEEL WIRES UNDER DIFFERENT
PROCESS PARAMETERS
NATEZNE LASTNOSTI HLADNO VLE^ENE MALOOGLJI^NE
JEKLENE @ICE PRI RAZLI^NIH PARAMETRIH PROCESA
Cem S. Çetinarslan1, Ali Güzey2
1Department of Mechanical Engineering, Faculty of Engineering and Architecture, Trakya University, 22180 Edirne, Turkey
2Arsay Wire Production Company-Kirklareli, Turkey
cemc@trakya.edu.tr
Prejem rokopisa – received: 2012-10-01; sprejem za objavo – accepted for publication: 2012-10-23
This study demonstrates the influence of drawing-process parameters such as reduction (deformation) ratio and drawing
velocity on the tensile properties of various low-carbon cold-malleable steel wires. Standard tensile tests were realized on four
types of wires – SAE1006, SAE1008, SAE1015 (Cq15) and SAE10B22 (20MnB4) – at various process parameters. This
experimental study shows how two of the main process parameters, the deformation ratio and drawing velocity, clearly influence
the tensile properties (yield stress, ultimate tensile strength, and elongation at rupture) of steel-wire materials.
Keywords: wire drawing, tensile properties, deformation (reduction) ratio, drawing velocity
Ta {tudija prikazuje vpliv procesnih parametrov pri vle~enju, kot sta odvzem (deformacija) in hitrost vle~enja, na natezne
lastnosti razli~nih malooglji~nih mehkih jeklenih `ic. Izvr{en je bil standardni natezni preizkus na {tirih vrstah `ice SAE1006,
SAE1008, SAE1015 (Cq15) in SAE10B22 (20MnB4) po razli~nih procesnih parametrih. Ta eksperimentalna {tudija je
pokazala, kako dva glavna procesna parametra, stopnja deformacije in hitrost vle~enja, vplivata na natezne lastnosti (napetost
te~enja, natezna trdnost in raztezek pri pretrgu) jeklene `ice.
Klju~ne besede: vle~enje `ice, natezne lastnosti, stopnja deformacije (odvzem), hitrost vle~enja
1 INTRODUCTION
Wire drawing is a metal-reducing process, in which a
wire rod is pulled or drawn through a single die or a con-
tinuous series of dies, thereby reducing its diameter.
Wire drawing is one of the most common plastic-
deformation processes. A wire rod is pulled or drawn
through a die or a series of dies, causing a reduction of
its diameter. In general, drawing is known as a process
performed at room temperature. Drawing of low-car-
bon-content steel wires is generally conducted at room
temperature employing a number of passes or reductions
through several dies. Sometimes it may be performed at
elevated temperatures for large wires to reduce drawing
forces.
Generally, steel wire is made of plain-carbon steel
grades. The steel-wire materials are semi-products suit-
able for cold-drawing processes. Although a steel wire
can be produced from stainless steel and other alloyed
steels, in industry it is mostly produced using plain-
carbon steels. The steel containing up to 1 % C is usually
used for steel-wire production; however, the largest part
of steel-wire production constitutes low-carbon steels
with less than 0.1 % C.1
Ferrous wires are used as semi products for electrical
wiring, ropes (rope wires are usually made of pearlitic
steel and have very high tensile properties), cables, struc-
tural components, springs, nails, spokes, musical instru-
ments, electrodes, paper clips, etc.2
Several studies on wire-drawing processes and some
process parameters that affect the wire-drawing process
have been performed. Toribio and Ovejero have investi-
gated the effect of cumulative cold drawing on the pear-
lite interlamellar spacing in eutectoid steel. Interlamellar
spacing in fully pearlitic steels decreases progressively
during the cold-drawing process and the diminishing rate
is not constant throughout the manufacturing route.3 The
effect of degree of deformations, ranging from 5 % to
30 % reductions, on the mechanical properties of cold-
drawn, mild-steel rods was experimentally investigated
by Alawode and Adeyemi.4 Languillaume et al. have
presented the results of a study concerning the influence
of heavy cold drawing and post-deformation annealings
on the microstructure of such pearlitic steel wires.5 On
the other hand, Vega et al. have studied the effect of the
process variables such as the semi-die angle, the
reduction in area and the friction coefficient on the
drawing-force value. The results of this study indicate
clearly that friction has a significant effect on the
drawing force, which becomes lower due to a decrease in
the area reduction.6 The influence of the main process
parameters (the wire yield stress – S, the cross-sectional
area reduction – Re and the die half angle – ) on the
shape quality and area fraction of the round-to-hexagonal
Materiali in tehnologije / Materials and technology 47 (2013) 2, 245–252 245
UDK 621.778:620.172 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 47(2)245(2013)
composite wire drawing were investigated by Norasetha-
sopon.7 This study shows that Re and S strongly
influence the shape quality, and S slightly influences the
change in the area fraction of the core. The change in the
area fraction of the core, which equals zero, was
obtained with the value of  that increased with the
increasing S. Re and S strongly influence the drawing
stress. Within this order, Re and S directly, strongly and
inversely influenced the optimal die half angle. The pass
schedule of a wire-drawing process designed to prevent a
delamination of a high-strength-steel cord wire was
studied by Lee et al.8 From their findings it is clear that
the applied drawing process reduced the diameter of the
wire from 3.5 mm to 0.95 mm, and that it consisted of
nine passes. On the other hand, another model for
predicting the fatigue strength of two different eutec-
toid-steel wires, one of them being zinc coated, used in
ropeway applications, has been presented by Beretta and
Boniardi.2 Within this method the fatigue process of
wires has been described in terms of propagation of the
surface defects caused by cold drawing.
The aim of our research was to investigate the tensile
properties of various low-carbon cold-malleable steel
wires with respect to drawing velocity and deformation
ratio. These parameters also have an influence on the
final wire quality, the drawing force, the lubrication in
the process, the mechanical properties and the die wear.
2 EXPERIMENTAL PROCEDURE
2.1 Preparation process
Wire-rod (raw) materials were of four different types
of low-carbon steel: SAE1006, SAE1008, SAE1015
(Cq15) and SAE10B22 (20MnB4). The steel chemical
compositions are given in Table 1.9
First, the chemical compositions of the steels were
measured using a SPECTROLAB M7 spectrometric test
device. Then the surface-cleaning process including two
stages, the mechanical and chemical cleaning, was per-
formed. The first step, the mechanical surface cleaning,
was applied to remove the scale layer from the wires and
then the chemical purification was realized. The
chemical cleaning consisted of causticization (for 25 min
in a KMnO5 + NaOH solution at 70 °C), dipping into an
acid bath (for 1 h in a HCl concentrated solution at room
temperature), washing and rinsing, passivation with lime
and, finally, drying (for 1h at 100 °C).
2.2 Wire-drawing process
After these treatments the drawing process was
performed. Figure 1 shows the outlet of a drawing die
with a coil (end product) and a drawing die (matrix) with
a soap box. A wire first passes through the soap box and
then through the die (matrix). The reduction of the
diameter of a metal wire is realized by pulling it through
the die (Figure 2). The working region of a die is
typically and made of W carbide. The die is cooled with
a cooling hose (water) as shown in Figure 1. A series of
dies is used to obtain the required diameter reduction of
the wire. Table 2 shows a series of dies with the
reductions of 5.5 to 2.2, to 1.8 and to 2.1 made in 8 or 9
passes, used to obtain the wire diameters of (4.80, 4.00
and 3.01) mm. The reduction ratio (R/%) was determined
for each diameter decrease as to the equation:
R/% =
D D
D
inlet outlet
inlet
2 2
2 100
−
× (1)
C. S. ÇETINARSLAN, ALI GÜZEY: TENSILE PROPERTIES OF COLD-DRAWN LOW-CARBON STEEL WIRES ...
246 Materiali in tehnologije / Materials and technology 47 (2013) 2, 245–252
Table 1: Chemical compositions of wire-rod (raw) steels
Tabela 1: Kemijska sestava jekla v palicah
Steel Type %C %Si % Mn %P %S %Cu %Cr %Ni %Mo %Al % B
SAE1006 0.06 0.2 0.35 0.04 0.05 0.30 0.15 0.3 0.03 – –
SAE1008 0.08 0.30 0.55 0.03 0.05 0.35 0.3 0.25 0.03 0,02 –
SAE1015 0.14 0.15 0.40 0.02 0.015 0.1 0.08 0.1 0.05 0.03
SAE10B22 0.21 0.15 1.00 0.015 0.015 0.1 0.08 0.1 0.05 0.02 0.002
Table 2: Series of dies for each steel type for the drawing process (reduction of 5.5 to 4.8, to 4.00 and to 3.01)
Tabela 2: Serija orodij za vsako vrsto jekla pri vle~enju (odvzem 5,5 do 4,8, do 4,00 in do 3,01)
Inlet dia. Outlet dia. Pass number 1 2 3 4 5 6 7 8 9
5.50 2.20 8 4.80 4.21 3.72 3.31 2.96 2.67 2.41 2.20 –
5.50 1.80 8 4.67 4.00 3.44 2.98 2.60 2.29 2.02 1.80 –
5.50 2.10 9 4.82 4.25 3.77 3.37 3.02 2.73 2.49 2.28 2.10
Figure 1: a) Outlets of a drawing die with a wire coil and b) a drawing
die (matrix) with a soap box
Slika 1: a) Sestav vle~ne matrice z navijalcem `ice, b) matrica za
vle~enje s posodo za milo
In a multipass drawing process, the temperature rise
during each pass can affect the mechanical properties of
the final product (such as its bending and torsion proper-
ties, and its tensile strength).8
A wire-drawing process was carried out with
different drawing velocities and different total-reduction
ratios of the deformation to determine how the tensile
properties of various low-carbon wires were affected.
The effect of drawing velocity and deformation ratio
was investigated in some references. One of those
focused on the influence of drawing speed on the
properties of multiphase TRIP (transformation induced
plasticity) steel wires10 and the other on the study of the
effect of total-reduction ratio on wire breaks by Cu
fine-wire drawing.11
In general, each pass ratio is between 1.68 and 1.09.
Dfinal-1 = k · Dfinal (k = 1.68–1.09) (2)
2.3 Tensile test
Experiments were carried out on a tensile-test
machine at room temperature (Figures 3 and 4) and
SAE1006, SAE1008, SAE1015 (Cq15) and SAE10B22
(20MnB4) coil wires (end products) were used as test
steels. The wires were submitted to tensile tests to
determine the yield stress, the ultimate tensile strength
and the elongation at rupture. Wire cuts of 250 mm in
length were used as the test specimens. The tensile
strength was determined on a 3 t tensile tester with a ram
(lower jaw) speed of 10 mm/min using various test
parameters and three experiments were carried out and
then averaged for each point in the diagrams.
3 RESULTS AND DISCUSSIONS
3.1 Test results at a constant drawing velocity (3.6 m/s)
and a constant reduction ratio (from 5.50 to
4.81)
Firstly, the wire-rod specimens were tested and then
the coil (drawn) wires were tested at a constant drawing
velocity and a constant reduction ratio from 5.50 to
4.81. Experimental findings on the yield strength,
tensile strength and elongation at rupture were deter-
mined for wire rods (before drawing) and coil (drawn)
wires and are given in Table 3.
The yield-strength values of rod wires and coil wires
were found as expected and, as shown in Figure 4, they
increase in accordance with the increasing C content.
The increase in the C content causes brittleness, making
C. S. ÇETINARSLAN, ALI GÜZEY: TENSILE PROPERTIES OF COLD-DRAWN LOW-CARBON STEEL WIRES ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 245–252 247
Figure 4: Variation in the yield-strength values for a rod and coil steel
wire at a constant velocity and reduction ratio (from 5.50 to 4.81)
Slika 4: Spreminjanje meje plasti~nosti palice in `ice v kolobarju pri
konstantni hitrosti in odvzemu (od 5,50 do 4,81)
Figure 2: Drawing die (matrix) with a tip (pressure type)
Slika 2: Orodje za vle~enje (matrica) s konico (tla~ne vrste)
Figure 5: Variation in the ultimate tensile-strength values for a rod
and coil steel wire at a constant velocity and reduction ratio (from
5.50 to 4.81)
Slika 5: Spreminjanje vrednosti natezne trdnosti palice in `ice v kolo-
barju pri konstantni hitrosti in odvzemu (od 5,50 do 4,81)
Figure 3: Tensile-test machine
Slika 3: Stroj za natezne preizkuse
plastic deformation more difficult. The ultimate tensile-
strength values of the rod materials and coils of the
tested steels are also in line with the increasing C content
(Figure 5). The variation in the elongation at rupture is
shown in Figure 6. The values of the elongation at
rupture for drawn wires decreased with the increasing
plastic deformation for all the tested steels. The
(20MnB4) steel shows a slight decrease in the elongation
due to a higher C content.
The wire specimens used for the tensile tests are
shown in Figures 7 and 8.
3.2 Test results at a constant drawing velocity (3.6 m/s)
and different reduction ratios (from 5.50 to 4.81,
to 4.00 and to 3.01)
The coil materials were tested at a constant drawing
velocity (3.6 m/s) and different reduction ratios (from
5.50 to 4.81, 4.00 and 3.01). The reduction ratio
was determined as depending on the constant inlet
diameter (5.50) and different outlet diameters (4.81,
4.00 and 3.01). Experimental findings are given in
Table 4. The yield strength, tensile strength and elonga-
tion at rupture were determined for the wire rods (before
drawing) and for drawn wires after various reduction
ratios.
C. S. ÇETINARSLAN, ALI GÜZEY: TENSILE PROPERTIES OF COLD-DRAWN LOW-CARBON STEEL WIRES ...
248 Materiali in tehnologije / Materials and technology 47 (2013) 2, 245–252
Table 3: Tensile properties of a rod and coil steel wire at a constant drawing velocity and a constant reduction ratio (from 5.50 to 4.81)
Tabela 3: Natezne lastnosti palice in `ice v kolobarju pri konstantni hitrosti vle~enja in konstantnem odvzemu (od 5,50 do 4,81)
Dia., /mm
Material
type
V
m/s
Yield Strength
Rp0,2/MPa
Ultimate Tensile
Strength
Rm/MPa
Elongation at Rupture
%
Specimen
No
Wire rod
(before
drawing)
Coil
(after
drawing)
Wire rod
(before
drawing)
Coil
(after
drawing)
Wire rod
(before
drawing)
Coil
(after
drawing)
Wire rod
(before
drawing)
Coil
(after
drawing)
1 5.50 4.81 SAE1006 3.6 257 357 357 477 41 29
2 5.50 4.81 SAE1006 3.6 261 359 361 481 44 30
3 5.50 4.81 SAE1006 3.6 259 360 363 484 43 31
1 5.50 4.81 SAE1008 3.6 293 388 413 535 39 32
2 5.50 4.81 SAE1008 3.6 297 388 421 540 38 28
3 5.50 4.81 SAE1008 3.6 301 390 422 542 40 29
1 5.50 4.81 SAE1015 3.6 316 410 441 561 41 28
2 5.50 4.81 SAE1015 3.6 311 411 437 558 40 28
3 5.50 4.81 SAE1015 3.6 324 407 444 564 41 32
1 5.50 4.81 SAE10B22 3.6 408 480 568 689 36 25
2 5.50 4.81 SAE10B22 3.6 405 467 555 675 37 26
3 5.50 4.81 SAE10B22 3.6 400 471 540 666 39 29
Figure 8: SAE10B22 (20MnB4) specimen
Slika 8: Vzorec SAE10B22 (20MnB4)
Figure 7: SAE1008 specimen
Slika 7: Vzorec SAE1008
Figure 6: Variation in the elongation-at-rupture (%) values for a rod
and coil steel wire at a constant velocity and reduction ratio (from
5.50 to 4.81)
Slika 6: Spreminjanje raztezka pri pretrgu (%) za palico in `ice v
kolobarju pri konstrantni hitrosti in odvzemu (od 5,50 do 4,81)
The yield-strength values for all the specimens
increase with the increasing reduction ratio (Figure 9);
and the ultimate tensile strength for all the specimens
shows similar tendencies (Figure 10). Variations in the
elongation at rupture are shown in Figure 11. In general,
the values of the elongation at rupture for coil wires
decrease as the reduction ratio increases. The decrement
is a bit larger for the relatively high C-content steel spe-
cimens, SAE1015 (Cq15) and SAE10B22 (20MnB4), as
the increase in the reduction ratio is more effective for
the steels containing a higher C content with respect to
the strain hardening. In addition, the Mn content is also a
strength-increasing alloying element for the steel
SAE10B22.12
3.3 Tensile-test results at a constant drawing velocity
(3.6 m/s) and different reduction ratios – different
inlet diameters and constant outlet diameters
Firstly, the wire rods were tested. The tests for coil
wires were realized at the constant drawing velocity (3.6
m/s) and different reduction ratios (from 5.50 to 4.81
C. S. ÇETINARSLAN, ALI GÜZEY: TENSILE PROPERTIES OF COLD-DRAWN LOW-CARBON STEEL WIRES ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 245–252 249
Table 4: Tensile properties of a rod and coil steel wire at a constant drawing velocity (3.6 m/s) and different reduction ratios (from 5.50 to
4.81, to 4.00 and to 3.01)
Tabela 4: Natezne lastnosti palice in `ice v kolobarju pri konstantni hitrosti vle~enja (3,6 m/s) in razli~nih odvzemih (od 5,50 do 4,81, do
4,00 in do 3,01)
Dia., /mm
Material
type
V
m/s
Yield Strength
Rp0,2/MPa
Tensile Strength
Rm/MPa
Elongation at Rupture
%
Specimen
No
Wire rod
(before
drawing)
Coil
(after
drawing)
Wire rod
(before
drawing)
Coil
(after
drawing)
Wire rod
(before
drawing)
Coil
(after
drawing)
Wire rod
(before
drawing)
Coil
(after
drawing
1 5.50 4.81 SAE1006 3.6
258
357
359
477
43
36
2 5.50 4.00 SAE1006 3.6 505 605 24
3 5.50 3.01 SAE1006 3,6 623 725 12
1 5.50 4.81 SAE1008 3.6
294
388
407
535
42
34
2 5.50 4.00 SAE1008 3.6 536 643 22
3 5.50 3.01 SAE1008 3.6 664 774 11
1 5.50 4.81 SAE1015 3.6
316
410
438
561
40
33
2 5.50 4.00 SAE1015 3.6 558 676 18
3 5.50 3.01 SAE1015 3.6 684 808 8
1 5.50 4.81 SAE10B22 3.6
406
480
571
689
38
31
2 5.50 4.00 SAE10B22 3.6 649 811 15
3 5.50 3.01 SAE10B22 3.6 772 944 5
Figure 11: Variation in the elongation-at-rupture (%) values for a rod
and coil steel wire at a constant velocity and different reduction ratios
(from 5.50 to 4.81, 4.00 and to 3.01)
Slika 11: Spreminjanje raztezka pri pretrgu (%) palice in `ice v
kolobarju pri konstantni hitrosti in razli~nih odvzemih (od 5,50 do
4,81, do 4,00 in do 3,01)
Figure 10: Variation in the ultimate-tensile-strength values for a rod
and coil steel wire at a constant velocity and different reduction ratios
(from 5.50 to 4.81, to 4.00 and to 3.01)
Slika 10: Spreminjanje naztezne trdnosti palice in `ice v kolobarju pri
konstantni hitrosti in razli~nih odvzemih (od 5,50 do 4,81, do
4,00 in do 3,01)
Figure 9: Variation in the yield-strength values for a rod and coil steel
wire at a constant velocity and different reduction ratios (from 5.50
to 4.81, to 4.00 and to 3.01)
Slika 9: Spreminjanje meje plasti~nosti za palico in `ico v kolobarju
pri konstantni hitrosti in razli~nih odvzemih (od 5,50 do 4,81, do
4,00 in do 3,01)
and from 6.50 to 4.81). These ratios were determined
according to different inlet diameters and constant outlet
diameters. The results are shown in Table 5 and Figures
12, 13 and 14. The yield strength, tensile strength and
elongation at rupture were determined for the wire rods
(before drawing) and coil (drawn) wires with different
ratios.
The yield strength, ultimate tensile strength and elon-
gation at rupture were affected by the reduction ratio for
each material as shown in Section 3.2. As the approxi-
mate reduction ratios (45.7 % for the reduction of 6.5
to 4.81 and 47.1 % for the reduction of 5.50 to 4.00)
were considered, it was understood that the variation in
the inlet diameters was not significant.
3.4 Tensile-test results at a constant reduction ratio
(from 5.50 to 4.81) and with different drawing
velocities (3.6 m/s and 2.4 m/s)
The coil-wire tests were realized at a constant
reduction ratio from 5.50 to 4.81 and different
drawing velocities (3.6 m/s and 2.4 m/s). The results are
shown in Table 6 and Figures 15, 16 and 17. It is
observed that the yield stress and the ultimate tensile
strength of the specimens increase with the increasing
drawing velocity for each type of the materials. A higher
C content leads to a higher yield and ultimate tensile
C. S. ÇETINARSLAN, ALI GÜZEY: TENSILE PROPERTIES OF COLD-DRAWN LOW-CARBON STEEL WIRES ...
250 Materiali in tehnologije / Materials and technology 47 (2013) 2, 245–252
Table 5: Tensile properties of a rod and coil steel wire at a constant velocity (3.6 m/s) and different reduction ratios – different inlet diameters
and a constant outlet diameter – (from 5.50 to 4.81 and from 6.50 to 4.81)
Tabela 5: Natezne lastnosti palic in `ice v kolobarju pri konstantni hitrosti (3,6 m/s) in razli~nih odvzemih – razli~en vstopni premer in enak
izhodni premer – (od 5,50 do 4,81 in od 6,50 to 4,81)
Dia., /mm
Material
type
V
m/s
Yield Strength
Rp0,2/MPa
Tensile Strength
Rm/MPa
Elongation at Rupture
%
Specimen
No
Wire rod
(before
drawing)
Coil
(after
drawing)
Wire rod
(before
drawing)
Coil
(after
drawing)
Wire rod
(before
drawing)
Coil
(after
drawing)
Wire rod
(before
drawing)
Coil
(after
drawing)
1 5.50 4.81 SAE1006 3.6
258
357
358
477
43
36
2 6.50 4.81 SAE1006 3.6 500 605 24
1 5.50 4.81 SAE1008 3.6
292
388
405
535
42
34
2 6.50 4.81 SAE1008 3.6 531 643 22
1 5.50 4.81 SAE1015 3.6
314
410
436
562
40
33
2 6.50 4.81 SAE1015 3.6 553 676 18
1 5.50 4.81 SAE10B22 3.6
406
480
567
689
38
31
2 6.50 4.81 SAE10B22 3.6 645 811 15
Figure 13: Variation in the ultimate-strength values for a rod and coil
steel wire at a constant velocity (3.6 m/s) and different reduction ratios
– different inlet diameters and constant outlet diameters (from 5.50
to 4.81 and from 6.50 to 4.81)
Slika 13: Spreminjanje natezne trdnosti palice in `ice v kolobarju pri
konstantni hitrosti (3,6 m/s) in razli~nih odvzemih – razli~en vstopni
premer in konstanten izstopni premer (od 5,50 do 4,81 in od 6,50
do 4,81)
Figure 14: Variation in the elongation-at-rupture (%) values for a rod
and coil steel wire at a constant velocity (3,6 m/s) and different reduc-
tion ratios – different inlet diameters and constant outlet diameters
(from 5.50 to 4.81 and from 6.50 to 4.81)
Slika 14: Spreminjanje raztezka pri pretrgu (%) palice in `ice v kolo-
barju pri konstantni hitrosti (3,6 m/s) in razli~nih odvzemih (od 5,50
do 4,81 in od 6,50 do 4,81)
Figure 12: Variation in the yield-strength values for a rod and coil
steel wire at a constant velocity (3.6 m/s) and different reduction ratios
– different inlet diameters and constant outlet diameters (from 5.50
to 4.81 and from 6.50 to 4.81)
Slika 12: Spreminjanje meje plasti~nosti palice in `ice v kolobarju pri
konstantni hitrosti (3,6 m/s) in razli~nih odvzemih – razli~en vstopni
premer in enak izstopni premer (od 5,50 do 4,81 in od 6,50 do
4,81)
strength and a higher drawing velocity. Drawing velocity
slightly affects the elongation, which decreases as the
drawing velocity increases. These values are quite
similar for all the steels.
4 CONCLUSIONS
The wire drawing of SAE1006, SAE1008, SAE1015
(Cq15) and SAE10B22 (20MnB4) low-carbon, malle-
able-steel wires was investigated and their tensile
properties were determined experimentally. This study
contributes to the knowledge of tensile properties and the
behaviour of drawn low-carbon steel wires during the
cold-drawing process. The effect of the process para-
meters (reduction ratio, drawing velocity) were studied
and it was found that the processing parameters have a
major influence on the tensile properties in all four types
of the low-carbon drawn steel wire.
The obtained results can be summarized as follows:
• The experiments have shown that the yield strength
and ultimate tensile strength increase, while the
elongation at rupture decreases for all the steels when
the reduction (deformation) ratio is increased.
• The drawing velocity has a significant effect on the
tensile properties (the yield and the ultimate tensile
strength) of low-carbon steel wires. A high drawing
velocity causes high strength properties. The values
of elongation at rupture also decrease as the drawing
velocity increases.
C. S. ÇETINARSLAN, ALI GÜZEY: TENSILE PROPERTIES OF COLD-DRAWN LOW-CARBON STEEL WIRES ...
Materiali in tehnologije / Materials and technology 47 (2013) 2, 245–252 251
Table 6: Tensile properties of a coil steel wire at a constant reduction ratio and different drawing velocities (3.6 m/s and 2.4 m/s)
Tabela 6: Natezne lastnosti `ice iz kolobarja pri konstantnem odvzemu in razli~nih hitrostih vle~enja (3,6 m/s in 2,4 m/s)
Dia., /mm Material
type
V
m/s
Yield Strength
Rp0,2/MPa
Tensile Strength
Rm/MPa
Elongation at
Rupture, %
Specimen
No
Wire rod (before
drawing)
Coil
(after drawing)
Coil
(after drawing)
Coil
(after drawing)
Coil
(after drawing)
1 5.50 4.81 SAE1006 3.6 357 477 36
2 5.50 4.81 SAE1006 2.4 310 455 38
1 5.50 4.81 SAE1008 3.6 388 535 34
2 5.50 4.81 SAE1008 2.4 350 510 36
1 5.50 4.81 SAE1015 3.6 410 561 33
2 5.50 4.81 SAE1015 2.4 370 536 35
1 5.50 4.81 SAE10B22 3.6 480 689 31
2 5.50 4.81 SAE10B22 2.4 444 650 33
Figure 15: Variation in the yield-strength values for a coil steel wire
at a constant reduction ratio (from 5.50 to 4.81) and with different
drawing velocities (3.6 m/s and 2.4 m/s)
Slika 15: Spreminjanje meje plasti~nosti `ice v kolobarju pri kon-
stantnem odvzemu (od 5,50 do 4,81) in razli~nih hitrostih vle~enja
(3,6 m/s in 2,4 m/s)
Figure 17: Variation in the elongation-at-rupture (%) values for a coil
steel wire at a constant reduction ratio (from 5.50 to 4.81) and with
different drawing velocities (3.6 m/s and 2.4 m/s)
Slika 17: Spreminjanje raztezka pri pretrgu (%) `ice iz kolobarja pri
konstantnem odvzemu (od 5,50 do 4,81) in razli~nih hitrostih vle-
~enja (3,6 m/s in 2,4 m/s)
Figure 16: Variation in the ultimate-tensile-strength values for a coil
steel wire at a constant reduction ratio (from 5.50 to 4.81) and with
different drawing velocities (3.6 m/s and 2.4 m/s)
Slika 16: Spreminjanje natezne trdnosti `ice iz kolobarja pri kon-
stantnem odvzemu (od 5,50 do 4,81) in razli~nih hitrostih vle~enja
(3,6 m/s in 2,4 m/s)
• It was determined that the reduction ratio has a larger
influence on the tensile properties of low-carbon steel
wires than the drawing velocity.
• Due to a high C content, the tensile-strength proper-
ties of the wires increased for all the reduction ratios.
In addition, Mn was also one of the strongly
influential elements and its effect was amplified by
increasing the strain rate for the SAE10B22 steel.12,13
The increase in the C content enhances the work-
hardening rate.14 The work-hardening ability of steel
increases with an increase in the C content. Thus, the
C content causes a significant variation in the tensile
strength of drawn steel wires. Moreover, it is known
that B enhances the tensile properties of low-carbon
steels.15,16
• The strength of rod wires can be improved using the
wire-drawing process according to the experimental
findings in this investigation. Furthermore, the wire-
drawing-process parameters, like the reduction ratio
and drawing velocity, also have a significant effect on
the tensile properties of steel wires.
Acknowledgements
The authors would like to thank Mr. Uður UZ for his
help with the experimental work and to the ARSAY
WIRE PRODUCTION COMPANY, Kirklareli/Turkey
and the YILMAR STEEL WIRE AND SPRING COM-
PANY, Bursa/Turkey for their technical support in the
experimental processes.
5 REFERENCES
1 T. Altan, S. Oh, H. L. Gegel, Metal forming, ASM, New York, 1983
2 S. Beretta, M. Boniardi, Fatigue strength and surface quality of
eutectoid steel wires, Int. J. Fatigue, 21 (1999), 329–335
3 J. Toribio, E. Ovejero, Effect of cumulative cold drawing on the pear-
lite interlamellar spacing in eutectoid, Scripta Mater, 39 (1998) 3,
323–328
4 A. J. Alawode, M. B. Adeyemi, Effects of degrees of deformation
and stress-relief temperatures on the mechanical properties and resi-
dual stresses of cold drawn mild steel rods, J. Mater. Process. Tech.,
160 (2005) 2, 112–118
5 J. Languillaume, G. Kapelski, B. Baudelet, Evolution of the tensile
strength in heavily cold drawn and annealed pearlitic steel wires,
Mater. Lett., 33 (1997), 241–245
6 G. Vega, A. Haddi, A. Imad, An investigation of process parameters
effect on the copper-wire drawing, Mater. Design, 30 (2009),
3308–3312
7 S. Norasethasopon, Influence of process parameters on shape quality
and area fraction in round-to-hexagonal composite wire drawing, J.
Mater. Process. Tech., 203 (2008), 137–146
8 S. K. Lee, D. C. Ko, B. M. Kim, Pass schedule of wire drawing
process to prevent delamination for high strength steel cord wire,
Mater. Design, 30 (2009), 2919–2927
9 A. Güzey, The Investigation of manufacturing process of ferrous
wires, MSc Thesis, Trakya University, Natural and Applied Science
Institute, Edirne 2009
10 M. Suliga, Z. Muskalski, S. Wiewiórowska, The influence of draw-
ing speed on properties of TRIP steel wires, J. Achieve. Mater.
Manuf. Eng., 26 (2008) 2, 151–154
11 H. Cho, H. H. Jo, S. G. Lee, B. M. Kim, Y. J. Kim, Effect of
reduction ratio, inclusion size and distance between inclusions on
wire breaks in Cu fine wiredrawing, J. Mater. Process Tech.,
130–131 (2002), 416–420
12 A. A. Gol’denberg, N. P. Sukhikh, T. M. Mineeva, Effect of man-
ganese and nickel on the strength of steel under rigid loading con-
ditions, Met. Sci. Heat Treat., 13 (1971) 6, 487–489
13 M. Itabashi, K. Kawata, Carbon content effect on high-strain-rate
tensile properties for carbon steels, Int. J. Impact Eng., 24 (2000),
117–131
14 R. Song, D. Ponge, D. Raabe, Improvement of the work hardening
rate of ultrafine grained steels through second phase particles,
Scripta Mater., 52 (2005), 1075–1080
15 M. I. Haq, N. Ikram, The effect of boron addition on the tensile
properties of control-rolled and normalized C-Mn steels, J. Mater.
Sci., 28 (1993) 22, 5981-5985
16 P. Hau{ild, J. Siegl, P. Málek, V. [íma, Effect of C, Ti, Zr and B
alloying on fracture mechanisms in hot-rolled Fe–40 (at.%)Al,
Intermetallics, 17 (2009), 680–687
C. S. ÇETINARSLAN, ALI GÜZEY: TENSILE PROPERTIES OF COLD-DRAWN LOW-CARBON STEEL WIRES ...
252 Materiali in tehnologije / Materials and technology 47 (2013) 2, 245–252
R. CELIN et al.: IMPELLER-BLADE FAILURE ANALYSIS
IMPELLER-BLADE FAILURE ANALYSIS
PREISKAVA PO[KODBE LOPATICE ROTORJA
Roman Celin, Franc Tehovnik, Franc Vodopivec, Borut @u`ek
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
roman.celin@imt.si
Prejem rokopisa – received: 2013-01-10; sprejem za objavo – accepted for publication: 2013-01-31
The axial pumps are used in a wide variety of applications, such as drainage control, power plants and process cooling. A
fresh-water, vertical, axial-pump, stainless-steel impeller failed. The impeller was a single cast from stainless steel. A visual
examination indicated that the fracture originated near the blade-to-hub attachment. During the investigation standard
nondestructive and destructive methods were used. Specimens from the failed blades were taken for a material characterization.
The goal of the examination was to determine a possible cause for the impeller-blade failure. After the examination it was
determined that the most probable cause for the impeller-blade failure was the fatigue stress.
Keywords: blade, impeller, cast stainless steel, crack
Aksialne ~rpalke se uporabljajo za razli~ne namene, kot so namakanje, hlajenje procesne opreme in komponent termoelektrarn.
Lopatica rotorja vertikalno postavljene aksialne ~rpalke se je nepri~akovano po{kodovala. Lopatica in rotor sta bila ulita iz
nerjavnega jekla kot celota. Vizualna preiskava rotorja je odkrila razpoko na vstopnem robu lopatice v bli`ini pesta. Med
preiskavo vzroka nastanka po{kodbe so bile uporabljene standardizirane poru{itvene in neporu{itvene metode. Iz lopatice rotorja
so bili izdelani preizku{anci za karakterizacijo materiala. Namen preiskave je bil odkriti verjeten vzrok za po{kodbo lopatice. Po
kon~ani preiskavi je bilo ugotovljeno, da je najbolj verjeten vzrok za nastanek razpoke utrujanje materiala zaradi vibracij lopatic
rotorja.
Klju~ne besede: lopatica, rotor, lito nerjavno jeklo, razpoka
1 INTRODUCTION
Centrifugal pumps are turbomachines used for
transporting liquids by raising a specified volume of the
flow to a specified pressure level. The basic centrifugal-
pump components are: the casing, the bearing housing,
the pump shaft and the impeller.1 A pump configuration
of the components may vary depending on the fluid-flow
direction that can be radial, semi-axial or axial. Axial-
flow pumps achieve larger flow rates than radial pumps
and are used in drainage control, power plants and pro-
cess cooling. Figure 1 shows a generic design of a verti-
cal, axial, centrifugal pump with the main components.
Any pump operation is determined by the flow rate Q
(m3/s), the head H (m) and impeller revolutions n
(min–1). The head is the measurement of the height (m)
of the liquid column the pump creates from the kinetic
energy that the pump gives to the liquid. The design and
operation of a pump also depends on the operation
efficiency, the stability of the head-capacity characte-
ristic, vibration and noise. An important issue is also a
possible pump failure due to fatigue, cavitation, hydro-
abrasive wear or erosion corrosion. Most vertical pumps
are out of sight during the operation and, for this reason,
Materiali in tehnologije / Materials and technology 47 (2013) 2, 253–258 253
UDK 620.179.1:620.179.11:691.714.018.8 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 47(2)253(2013)
Figure 1: Vertical, axial, centrifugal pump
Slika 1: Vertikalna aksialna centrifugalna ~rpalka
Figure 2: Damaged impeller blade
Slika 2: Po{kodovana lopatica rotorja
it is important to monitor the shaft vibrations and other
operational parameters.
Twelve months after a plant outage a vertical, axial-
flow, fresh-water pump, stainless-steel impeller failed.
The hub and impeller blades were a single cast made of
the ASTM A743 stainless steel (grade 316). This stain-
less-steel grade was a material of choice because of its
good erosion-resistant properties. The axial-pump flow
rate was Q = 8.453 m3/s, the head was H = 13.72 m and
the number of revolutions was n = 370 min–1. A
complete pump assembly was dismantled and the
maintenance crew immediately discovered a very distinct
crack on an impeller blade (Figure 2). The pump owner
decided to carry out a failure analysis.
A failure analysis is a broad discipline that includes
materials and mechanical engineering. The purpose of
this paper is to present a failure-analysis procedure
applied in the case of an impeller-blade failure investi-
gation at the Institute of Metals and Technology (IMT,
Ljubljana, Slovenia).
2 EXAMINATION PROCESS
The failure analysis is explained in books2–6 and
papers.7–11 The examination process starts when a com-
ponent under observation has lost its designated function
in a system. In general, the examination process or
analysis of a damaged component is performed in several
steps, which are described below. The first step is usually
the on-site visit and the gathering of all the available
information on the failed component and the in-service
conditions of the component.
The next step includes nondestructive examinations.
There is a variety of nondestructive techniques avail-
able12. The search for material imperfections is perfor-
med with X-ray, magnetic particle, ultrasonic, liquid
penetrant, eddy current, and other nondestructive testing
procedures. The most common is the visual examination
aiming to determine the general mechanical and struc-
tural conditions of the components. The result of a visual
examination is a record in the form of a sketch, dimen-
sion-measurement data or a photography, identifying
discontinuities or imperfections on the surface of the
components such as cracks, wear, tear, corrosion,
erosion, etc.
Based on the results of a visual examination and the
on-site information further decisions on the course of the
examination are made. Usually that means establishing a
plan for a destructive testing that involves the cutting off
a sample material from the failed component, and an
investigation of the samples in a laboratory for a chemi-
cal analysis, metallography, a mechanical testing and
others depending on the testing plan.
A chemical analysis is performed on the original
material to verify if the material sample meets the
appropriate specification or standard, and whether a
deviation from the specifications could have contributed
to the failure. A wet chemical analysis, atomic absorp-
tion, X-ray photoelectron, Auger electron and inducti-
vely coupled plasma-mass spectrometry (ICP MS) are
some of the suitable methods of a chemical analysis.
The tensile test is one of the most frequently used
tests for evaluating the mechanical properties of mate-
rials.13 A tensile force is applied with a machine and a
gradual elongation and the final fracture of the test
samples are obtained. The tensile test provides the
force-extension data that can quantify the quasi-static
mechanical properties of a material: yield strength,
ultimate tensile strength, elongation and reduction of
area at fracture.
Charpy (CVN) toughness tests are widely used to
determine the impact toughness of a material and the
effect of temperature on the sensitivity of structural
steels to brittle fracture. Notched specimens are sub-
mitted to the impact of a hammer with the kinetic energy
of 300 J. The fracturing occurs in a ductile, mixed or
brittle mode and, accordingly, very different quantities of
energies are consumed.
A careful investigation of the macrostructure and
microstructure of a failed material can provide the most
important information. A macroscopic examination of a
component sample evaluates the surface of the compo-
nent sample at a low magnification (usually up to 10
times). The type of a fracture such as ductile, brittle or
torsion can be identified.
Metallographic examinations are performed with an
appropriate magnification of optical and scanning
electron microscopes (SEM). During an optical micro-
scopic examination the grain size, microcracks, the gene-
ral microstructure and inclusion content are determined.
Scanning electron microscopy is used to determine small
details of a microstructure, fracture, precipitate size and
distribution and characteristics of crack initiation and
propagation. Furthermore, with the use of an energy
dispersive analysis (SEM/EDS) corrosion products on a
fracture surface can be identified. It is useful to compare
the microstructure of the samples removed from a failed
component with the samples removed from the sound
sections of a component.
A collection of visual, metallographic and SEM
results along with a chemistry analysis, mechanical data
and on-site information provides a solid ground for an
examiner to put together conclusions on the causes and
mechanism of a component failure. This is no easy task,
because, in many cases, the failure reasons are not
obvious even if a lot of information is available. Suffi-
cient experience in a failure analysis is necessary to
identify the cause of a failure.
3 RESULTS AND DISCUSSION
3.1 Visual examination
Only a visual examination14 was performed on the
failure site. The stainless-steel impeller (Figure 3) was a
cone-shaped hub with an approximate thickness of 35
mm with six blades attached. The blades vary in thick-
R. CELIN et al.: IMPELLER-BLADE FAILURE ANALYSIS
254 Materiali in tehnologije / Materials and technology 47 (2013) 2, 253–258
ness from approximately 25 mm near the hub to an
average of 10 mm at the top of the blade. Figures 2 and
3 show that the cast surface of the blades and the hub
was not ground. Several areas were detected on the
blades and at the attachment zone of the blades and the
hub (Figures 4 and 5) where weld-repair work on the
cast was performed by the supplier.
The top of the blade started to rub against the inner
ring surface and possibly the impeller operated in a
damaged condition for some time. The material loss at
the top of the blade, caused by friction wearing, is
noticed on Figure 3.
The most obvious feature on Figures 2 and 3 is the
blade splitting over the entire hub. The crack propagated
at an angle of approximately 30°–35° from the blade
leading edge toward the outer edge. The visual exami-
nation of the other blades revealed a through-thickness
crack at the leading edge of the blade marked as sample
1 and it is shown, in detail, on Figure 5.
The crack initial point on both blades (Figures 3, 4
and 5) was above the weld-heat-affected zone. It was not
possible to determine the exact initial point of the crack
on the blade leading edge. The crack surface was smooth
and mostly plastically deformed due to the surface
grinding during the pump operation.
Tree samples (Figure 3) were chosen for further
analysis. Figure 5 shows sample 1 which was used for a
fracture analysis, while samples 2 and 3 were used for
manufacturing the tensile and Charpy V-notch speci-
mens.
3.2 Chemical analysis
A quantitative chemical analysis was performed on
the impeller-blade sample material. It was made with an
ICP mass spectrometer. The results of the composition
analysis are shown in Table 1.
R. CELIN et al.: IMPELLER-BLADE FAILURE ANALYSIS
Materiali in tehnologije / Materials and technology 47 (2013) 2, 253–258 255
Figure 5: Crack at the leading edge of the blade
Slika 5: Razpoka na vodilnem robu lopatice
Figure 3: Positions of Figures 4 and 5 with samples 1, 2 and 3 before
the removal
Slika 3: Polo`aj slik 4 in 5 ter oznake vzorcev 1, 2 in 3 pred odvze-
mom
Table 1: Chemical-composition comparison in mass fractions (w/%)
Tabela 1: Kemijske sestave materiala rotorja v masnih dele`ih (w/%)
Cr Ni Mo Mn Si P C S N
sample 1 17.5 9.73 2.05 1.07 1.03 0.022 0.03 0.003 0.044
material test report 19.2 9.6 2.3 1.08 1.28 0.03 0.01 0.01 –
ASTM A 743 CF3M 17-21 9-13 2-3 1.5 1.5 0.04 0.03 0.04 –
Figure 4: Crack surface with a repair-weld build up at the blade
leading-edge attachment to the hub
Slika 4: Povr{ina razpoke z reparaturnim zvarom na stiku ~ela
lopatice in pesta
The sample-1 chemical-composition-analysis results
and impeller-manufacturer material-test report were in
accordance with the ASTM grade A743 (316L) cast
stainless-steel specification requirements presented in
Table 1.
3.3 Mechanical testing
A series of standard tensile tests15 and impact tests16
were performed with the specimens machined from the
samples shown in Figure 1. The goal of the standard
mechanical tests was to find out if the sample material
was in accordance with the manufacturer’s material-test
report and the ASTM specification. The tensile-test
results are presented in Table 2.
Table 2: Tensile-test results
Tabela 2: Rezultati nateznega preizkusa
Yield strength
Rp0,2 /MPa
Tensile strength
Rm/MPa
Elongation
A/%
sample 2
1 262 524 57
2 284 527 59
3 301 562 64
material test report
US units 48 PSI 91 PSI 30
ASTM A 743 specification
SI units min. 205 min. 485 30
US units min. 30 PSI min. 70 PSI 30
The yield-strength and tensile-strength results for the
three machined specimens are in accordance with the
manufacturer’s material-test results and ASTM specifi-
cation requirements.
The Charpy impact tests were performed at room
temperature, according to the SIST EN standard, on a
testing machine with an impact pendulum of a 300 J
capacity. The Charpy-test results on the three V-notch
samples were 250 J, 203 J and 261 J of the absorbed
pendulum energy, which is above the minimum required
value of 100 J from the owner’s specification.
Figure 6 shows sample 1, cut from the hub, clamped
on the tensile-testing machine with the purpose to obtain
the fracture surface without any grinding or other dama-
ge.
3.4 Metallographic examination
Figure 7 shows the cast A743-steel microstructure of
the hub at the blade side B of specimen 1 and Figure 8
shows the microstructure of the blade leading edge (side
A).
It consists of dendritic austenite grains with the
inserts of  ferrite at some grain boundaries. The content
of the  ferrite was of about 5 %. The microstructure
shown on both pictures is without any peculiarity and it
is typical for the cast stainless steel.
3.5 Scanning electron microscopy (SEM)
Figure 9 shows the side-A fracture-surface view. The
bright area (marked as 1.2) on the right is the fracture
R. CELIN et al.: IMPELLER-BLADE FAILURE ANALYSIS
256 Materiali in tehnologije / Materials and technology 47 (2013) 2, 253–258
Figure 8: Microstructure of the blade leading edge (sample 1 – side
A)
Slika 8: Mikrostruktura roba lopatice (vzorec 1 – stran A)
Figure 6: Sample-1 crack opening on a tensile-testing machine
Slika 6: Razpiranje vzorca 1 na trgalnem stroju
Figure 7: As-cast microstructure of the hub-to-blade attachment
Slika 7: Mikrostruktura ulitka na stiku pesta in lopatice
surface obtained by tensile-machine crack opening
(Figure 6).
The sample-1 side-A fracture surface has the same
features as the fracture surface shown on Figure 4. The
surface is smooth and the mostly plastically deformed
area marked as 1.1 (Figure 9) was carefully examined
with an optical microscope. Because of the plastically
deformed surface, the exact crack point of origin was
impossible to ascertain. Also, the crack propagation
mode could not be distinguished on the damaged surface
(Figure 10) and the fatigue striations were not detected
either.
The area (marked as 1.2 on Figure 6) enabled a SEM
investigation of the fracture surface, and Figure 11
shows two fracture modes. One (Figure 11 a) presents
an interphase crack propagation on the interface of
austenite and the inserts of  ferrite, while the other
(Figure 11 b) looks like a ductile fracture with small
dimples.
4 CONCLUSIONS
The examination procedure described was performed
to determine the mechanism of a failure of stainless-steel
impeller blades. According to the owner’s data the
impeller operated within the design parameters. For the
given application, the impellers appear to have been
fabricated within the tolerances and specifications
required by the owner. The surface of the hub and blades
was as cast. In general, an as-cast condition of a surface
as well as the weld-repair places are sensitive regions
prone to a crack formation because of dynamic loading.
The metallographic examination did not reveal a
deteriorative influence of a weld on the microstructure
around the crack-initiation area. The crack-initiation
point could not be determined because of the deformed
surface.
After the examination it was concluded that the most
probable cause for the impeller-blade failure was the
fatigue caused by the flow-induced vibrations due to a
turbulent flow over the blades and the internal stresses
caused by the welding repair of the casting defects.
The crack-initiation point on the blade leading edge
was probably a weak spot between the columnar den-
dritic grains of the cast.
5 REFERENCES
1 J. F. Gülch, Centrifugal pumps, Springer-Verlag, Berlin 2008, 61 or
368
2 K. A. Esakul, Handbook of case histories in failure analysis, Vol. 1,
ASM International, 1992, 251
3 H. M. Tawancy, A. Ul-Hamid, N. M. Abbas, Practical engineering
failure analysis, Marcel Dekker, New York 2004, 14
4 W. T. Becker, R. J. Shipley, ASM Handbook, Vol. 11, Failure analy-
sis and prevention, Materials Park: ASM International, 2002, 316
5 A. K. Das, Metallurgy of failure analysis, Mc Graw-Hill, New York
1997, 55
R. CELIN et al.: IMPELLER-BLADE FAILURE ANALYSIS
Materiali in tehnologije / Materials and technology 47 (2013) 2, 253–258 257
Figure 11: a) Interphase fracture and b) ductile fracture
Slika 11: a) Medfazni in b) `ilav prelom
Figure 10: SEM image of plastically deformed surface area 1.1
Slika 10: SEM-posnetek deformirane povr{ine, ozna~ene z 1.1
Figure 9: Side-A fracture surface of sample 1
Slika 9: Povr{ina preloma na strani A vzorca 1
6 C. Brooks, A. Choudhury, Failure analysis of engineering materials,
Mc Graw-Hill, New York 2002, 8
7 M. Torkar, F. Tehovnik, M. Godec, Mater. Tehnol., 46 (2012) 4,
423–427
8 A. van Bennekom, F. Berndt, M. N. Rassol, Eng. Failure Analysis, 8
(2001) 2, 145–156
9 M. Muhi~, J. Tu{ek, F. Kosel, D. Klob~ar, M. Pleterski, Thermal
fatigue cracking of die-casting dies, Metallurgy, 49 (2010) 1, 9–12
10 G. Kosec, A. Nagode, I. Budak, A. Anti}, B. Kosec, Failure of the
pinion from the drive of a cement mill, Eng. Failure Analysis, 18
(2011) 1, 450–454
11 R. B. Setterlund, V. K. Malhotra, L. K. Friesen, Failure of 13 %
Chromium Stainless Steel Impeller From Hydrogen Recycle Com-
pressor, Proc. of the Corrosion 2000, Orlando, USA, 2000
12 ASM Handbook, Volume 17, Nondestructive Evaluation and Quality
Control
13 ASM Handbook, Volume 8, Mechanical Testing and Evaluation
14 SIST EN 13018, Non-destructive testing – Visual testing – General
principles
15 SIST EN ISO 6892-1, Tensile testing, Method of test at ambient
temperature
16 ISO 148-1, Charpy pendulum impact test
R. CELIN et al.: IMPELLER-BLADE FAILURE ANALYSIS
258 Materiali in tehnologije / Materials and technology 47 (2013) 2, 253–258