VSEBINA – CONTENTS
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Application of a neural network for estimating the crack formation and propagation in sol-gel CeO2 coatings during
processing at temperature
Uporaba nevronske mre`e za ugotavljanje nastanka in {irjenja razpoke v sol-gel premazu CeO2 med postopkom ogrevanja
A. Savran, M. Alcý, S. Yýldýrým, R. Yiðit, E. Çelik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 453
Optimization of the mechanical and tribological properties of extruded AMCs: extension of the algorithm searching area via
multi-strategies
Optimizacija mehanskih in tribolo{kih lastnosti ekstrudiranih AMC: raz{iritev algoritma iskalnega podro~ja z uporabo ve~ strategij
M. O. Shabani, A. A. Tofigh, M. R. Rahimipour, M. Razavi, A. Mazahery, F. Heydari . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 459
Tribo-corrosion properties of a NiTi dental wire
Tribokorozijske lastnosti dentalne `ice NiTi
P. Mo~nik, T. Kosec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 467
Effect of a nano-ceramic mold coating on the fluidity length of thin-wall castings in Al4-1 alloy gravity sand casting
Vpliv nanokerami~nega premaza pe{~ene forme na teko~nost v tankostenskih ulitkih iz zlitine Al4-1 pri gravitacijskem ulivanju
M. Borouni, B. Niroumand, M. H. Fathi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 473
Influence of the thermo-mechanical treatment on the exfoliation and pitting corrosion of an AA5083-type alloy
Vpliv termo-mehanske obdelave zlitine AA5083 na lu{~enje in jami~asto korozijo
A. Halap, M. Popovi}, T. Radeti}, V. Va{~i}, E. Romhanji . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 479
Structural and electrical studies of fullerene (C60) dispersed polymer electrolytes
[tudij strukturnih in elektri~nih lastnosti polimernega elektrolita z dispergiranim fulerenom (C60)
A. Saxena, P. Kumar Singh, B. Bhattacharya . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 485
Effect of tunneling defects on the joint strength efficiency obtained with FSW
Vpliv tunelskih napak na u~inkovitost trdnosti spoja, izdelanega z drsno me{alnim varjenjem
S. Balos, L. Sidjanin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 491
The influence of lamination and conductive printing inks on smart-card operability
Vpliv laminacije in prevodnih tiskarskih barv na delovanje pametnih kartic
M. \oki}, V. Radoni}, V. Mladenovi~, A. Pleter{ek, U. Kav~i~, A. Hladnik, V. Crnojevi}-Bengin, T. Muck . . . . . . . . . . . . . . . . . . . . . . 497
Characterization of the substrates from two cultural-heritage sites and a preparation of model substrates
Karakterizacija gradbenih materialov iz dveh objektov kulturne dedi{~ine in priprava modelnih materialov
S. Kramar, V. Ducman, S. Vucetic, E. Velkavrh, M. Radeka, J. Ranogajec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 505
Long-term durability properties of pozzolanic cement mortars
Dolgoro~na obstojnost pucolanske cementne malte
L. Zavr{nik, J. Strupi [uput, S. Kramar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 509
Estimation of the effective boron-diffusion coefficient in the Fe2B layers grown on gray cast iron
Dolo~anje efektivnega koeficienta difuzije bora pri rasti plasti Fe2B na sivem litem `elezu
B. Bouarour, M. Keddam, O. Allaoui. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 515
Optimization of the secondary cooling in a continuous casting process with different slab cross-sections
Optimizacija sekundarnega hlajenja pri kontinuirnem ulivanju slabov z razli~nimi prerezi
T. Mauder, J. Stetina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 521
Unsteady model-based predictive control of continuous steel casting by means of a very fast dynamic solidification model on a
GPU
Predvidevanje kontrole kontinuirnega litja na podlagi neravnote`nega modela z zelo hitrim dinami~nim modelom strjevanja na GPU
L. Klimes, J. Stetina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 525
DSC/TG of Al-based alloyed powders for P/M applications
DSC/TG prahov na osnovi Al primernih za P/M uporabo
B. [u{tar{i~, J. Medved, S. Glode`, M. [ori, A. Koro{ec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 531
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 48(4)451–606(2014)
MATER.
TEHNOL.
LETNIK
VOLUME 48
[TEV.
NO. 4
STR.
P. 451–606
LJUBLJANA
SLOVENIJA
JULY–AUG.
2014
Microstructure characteristics of the model spring steel 51CrV4
Zna~ilnosti mikrostrukture modelnega vzmetnega jekla 51CrV4
M. Torkar, F. Tehovnik, B. Arh, M. Jenko, B. [arler, @. Raji} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 537
Effect of creep strain on creep rate in the temperature range 550–640 °C
Vpliv deformacije na hitrost lezenja pri temperaturah od 550 °C do 640 °C
B. @u`ek, F. Vodopivec, M. Jenko, B. Podgornik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 545
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Identification of the initial failure and damage of substituents of a unidirectional fiber-reinforced composite using a
micromodel
Uporaba mikromodela za ugotavljanje za~etnih napak in po{kodb sestavin kompozita, enosmerno oja~anega z vlakni
H. Srbová, T. Kroupa, R. Zem~ík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 549
Transformation-induced plasticity in steel for hot stamping
Vpliv pretvorbe na plasti~nost jekla za vro~e {tancanje
B. Ma{ek, C. [tádler, H. Jirková, P. Feuser, M. Selig . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 555
Influence of temperature and binder content on the properties of a sintered product based on red mud
Vpliv temperature in koli~ine veziva na lastnosti proizvoda iz sintranega rde~ega blata
M. Krgovi}, I. Bo{kovi}, R. Zejak, M. Kne`evi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 559
Influence of the precipitation temperature on the thrust force and torque in drilling an Al 2219-SiCp composite
Vpliv temperature izlo~evalnega `arjenja na potisno silo in navor pri vrtanju kompozita Al 2219-SiC
R. Ganesh, K. Chandrasekaran . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 563
Optimization of the recycling processes for magnesium from a highly contaminated waste
Optimiranje postopka recikliranja magnezija iz mo~no kontaminiranih odpadkov
V. Manojlovi}, @. Kamberovi}, M. Soki}, Z. Guli{ija, V. Matkovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 571
Machinability of a Ti-6Al-4V alloy with cryogenically treated cemented carbide tools
Obdelovalnost zlitine Ti-6Al-4V s kriogensko obdelanimi orodji iz karbidne trdine
A. Mavi, I. Korkut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 577
A cost-effective approach to the rapid fabrication of functional metal prototypes
Stro{kovno u~inkovit na~in za hitro izdelavo funkcionalnih kovinskih prototipov
C. C. Kuo . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 581
Thermomechanical treatment of Ti-Nb-V-B micro-alloyed steel forgings
Termomehanska obdelava odkovkov iz mikrolegiranega jekla Ti-Nb-V-B
M. Opiela . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 587
Progressive failure analysis of composite sandwich beam in case of quasistatic loading
Napredna analiza po{kodb sestavljenega kompozitnega nosilca pri kvazistati~nem obremenjevanju
T. Mandys, T. Kroupa, V. La{ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 593
Influence of heat treatment on the microstructure and mechanical properties of aluminium bronze
Vpliv toplotne obdelave na mikrostrukturo in mehanske lastnosti aluminijevega brona
P. Sláma, J. Dlouhý, M. Kövér . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 599
DOGODKI – EVENTS
[olski Center [kofja Loka se je predstavil tudi predsedniku RS Borutu Pahorju . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 605
A. SAVRAN et al.: APPLICATION OF A NEURAL NETWORK FOR ESTIMATING THE CRACK FORMATION ...
APPLICATION OF A NEURAL NETWORK FOR
ESTIMATING THE CRACK FORMATION AND
PROPAGATION IN SOL-GEL CeO2 COATINGS DURING
PROCESSING AT TEMPERATURE
UPORABA NEVRONSKE MRE@E ZA UGOTAVLJANJE NASTANKA
IN [IRJENJA RAZPOKE V SOL-GEL PREMAZU CeO2 MED
POSTOPKOM OGREVANJA
Aydogan Savran1, Musa Alcý1, Serdar Yýldýrým2,3, Recep Yiðit3,4,5, Erdal Çelik2,3,5
1Ege University, Dept. of Electrical and Electronics Engineering, 35100 Bornova, Izmir, Turkey
2Dokuz Eylul University, Dept. of Metallurgical and Materials Engineering, 35160 Buca, Izmir, Turkey
3Dokuz Eylul University, Center for Production and Applications of Electronic Materials (EMUM), 35160 Buca, Izmir, Turkey
4Dokuz Eylul University, Izmir Vocational School, Department of Technical Programs, 35150 Buca, Izmir, Turkey
5Dokuz Eylul University, Dept. of Nanoscience and Nanoengineering, 35160 Buca, Izmir, Turkey
recep.yigit@deu.edu.tr
Prejem rokopisa – received: 2012-09-21; sprejem za objavo – accepted for publication: 2013-10-28
In this study the application of a neural network to estimate the crack propagation and crack size in CeO2 coatings on a Ni
substrate during processing at temperature was evaluated as a function of the Ce content in solutions with increasing processing
temperatures from 24 °C and 700 °C. In this respect, CeO2 coatings were prepared on Ni tapes from solutions derived from
Ce-based precursors using a sol-gel method for YBCO-coated conductors. The crack size of the coating was determined using
an in-situ Hot-Stage ESEM depending on the temperature at a certain time in vacuum conditions. It was determined that the
crack size of the coating increased with the increasing processing temperature. Measuring the crack sizes of the coatings using
Hot-Stage ESEM is an expensive and time-consuming process. In order to eliminate these kinds of problems a neural-network
approach was used to estimate the crack sizes of the coatings at different temperatures. The neural network was constructed
directly from the experimental results. It was concluded that the estimation of the crack propagation of CeO2 coatings on a Ni
tape substrate are reasonable for the processing temperatures.
Keywords: CeO2, sol-gel, neural network, crack
V tej {tudiji je bila ocenjena uporaba nevronske mre`e za dolo~anje {irjenja razpoke in velikosti razpoke v premazu iz CeO2 na
podlagi iz Ni med ogrevanjem kot funkcija vsebnosti Ce v raztopini med ogrevanjem od 24 °C do 700 °C. Za ta namen je bil
pripravljen premaz iz CeO2 na traku iz Ni iz predhodne raztopine Ce z uporabo sol-gel metode za premaz na prevodniku YBCO.
Velikost razpoke v premazu je bila dolo~ena v vro~em z in-situ ESEM v odvisnosti od temperature po dolo~enem ~asu v
vakuumu. Ugotovljeno je bilo, da velikost razpoke v premazu nara{~a z nara{~anjem temperature procesa. Merjenje velikosti
razpoke v vro~em z uporabo ESEM je drag in ~asovno zamuden postopek. Za odpravo teh te`av je bila za dolo~anje velikosti
razpoke v premazu pri razli~nih temperaturah uporabljena nevronska mre`a. Ta je bila postavljena na podlagi eksperimentalnih
rezultatov. Ugotovljeno je bilo, da je mogo~e dolo~anje {irjenja razpoke v premazu CeO2 na traku iz Ni pri razli~nih
temperaturah procesa.
Klju~ne besede: CeO2, sol-gel, nevronska mre`a, razpoka
1 INTRODUCTION
Neural networks (NNs) possess massive parallelism,
a powerful mapping capacity, and a learning property.1
They are capable of making decisions based on incom-
plete, noisy and disordered information. In addition, they
can provide reasonable outputs for inputs not encoun-
tered during training. Because of these properties they
have gained the attention of scientist from many different
areas. There are a huge number of papers in the area of
control, identification and estimation.2–5 Most of the
industrial productive processes exhibit a certain degree
of nonlinearity. By means of their nonlinear mapping
and learning properties, neural networks have been
employed in the identification of unknown nonlinear
systems.
The development of the new, high-temperature super-
conductors has resulted in many applications in the fields
of magnetic, electronic and microwave technologies.
Ceramic superconductors have been deposited as thin
films using various methods, e.g., electron-beam co-eva-
poration, chemical vapor deposition, ion-beam sputter-
ing, pulsed-laser deposition and the sol-gel technique on
many buffered substrates.6 CeO2 is one of the buffer
layers frequently used for surface-coated YBCO super-
conductors on textured nickel substrates.7 The sol-gel
process offers a new possibility for the synthesis of oxide
layers by applying liquid precursors to substrates by
dipping, spinning or spraying methods.8 The major fac-
tors affecting the CeO2 sol-gel coatings are the nature of
the coating and the substrate, and the reaction parameters
at the temperature of YBCO formation.9 The nature of
the coatings are influenced by sol-gel parameters such as
types of precursor, solvent and chelating agents, visco-
sity, Ce content, dilution, chelation, complexation, with-
drawal rate, coating thickness and annealing conditions.
Materiali in tehnologije / Materials and technology 48 (2014) 4, 453–457 453
UDK 666.3/.7:004.032.26 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)453(2014)
In particular, the cracks of buffer layers are the main
problem in sol-gel processing.10
In our initial attempts11 we determined the crack-pro-
pagation rate and the crack size of CeO2 buffer layers on
Ni tapes for YBCO-coated conductors using an in-situ,
hot-stage, environmental scanning electron microscope
(ESEM). In the present investigation, in order to deter-
mine the crack size of the CeO2 sol-gel coatings by
avoiding time-consuming and expensive experimental
works, we propose to use a multilayer feed-forward NN
to estimate the crack size at the processing temperatures.
2 NEURAL NETWORKS
Neural Networks (NNs) are massively parallel, distri-
buted, adaptive, nonlinear information processing tools.
They consist of simple nonlinear processing elements
(PEs), referred to as neurons. Each neuron receives con-
nections from other PEs and/or itself. The block diagram
of a neuron is shown in Figure 1.
The neuron collects the values from all of its input
connections, performs a predefined mathematical opera-
tion, and produces a single output value computed as:
y f w u bj ji i j
i
n
= +
⎛
⎝
⎜ ⎞
⎠
⎟
=
∑
1
(1)
where ui is the input value, wji is the connection weight
between the ith input and the jth neuron, and f is the
activation function. The sum of the weighted inputs and
the bias forms the input to an activation function f. The
neurons may use any differentiable activation function f
to generate their output. The sigmoid function is the one
of the most common activation functions, as follows:
f x
e x
( ) =
+ −
1
1
(2)
The interconnectivity defines the topology of the NN.
An important class of NNs is the multilayer feed-forward
neural network, as shown in Figure 2. They consist of
simple nonlinear processing elements (PEs) and
weighted connections in a layered structure. The input
signal propagates through the network in a forward
direction from one layer to the next. The neurons are
connected with weighted connections represented by the
arrows in Figure 2.
The training of the NN is performed by adjusting the
connection weights. It is performed by iteratively adjust-
ing the weights (w) of the connections and biases (b) in
the network in order to minimize a predefined cost func-
tion. A popular cost function is the sum of the squared
error between the actual output and the desired output
value for each unit in the output layer:
E t yk k
k
= −∑1
2
2( ) (3)
where tk is the desired or target response on the kth unit
and yk is the produced network output on the same unit.
In order to perform training, a training set, including
the input and desired response vectors, is prepared. The
training is performed in two phases referred to as the
forward and backward phases. In the forward phase, the
input patterns from the training set are applied to the
input layer. These inputs are multiplied by the associated
weights, passed through the activation functions and
transferred to the next layer. They propagate through the
network, layer by layer. Finally, the actual response of
the network is produced at the output layer. The errors
between the desired and the network outputs are found.
In the backward phase, the error signal propagates back-
wards through the network and the gradients are com-
puted. These are then used to determine the weight
changes in the net according to used learning rule. A
more common learning algorithm is the back-propaga-
tion algorithm, which is introduced in12.
The standard back-propagation implements the
steepest-descent method (also called the gradient-descent
method). At each step of the steepest-descent method the
weights are adjusted in the direction in which the error
function decreases most rapidly. This direction is
determined by the gradient of the error surface at the
current point in the weight space. The weights are
updated in a negative direction to the gradient with a
certain rate, as given by:
w w
E
w
w w
E
w
w w
ji ji
ji
lj lj
lj
kl
new old
new old
new
= −
= −
=


∂
∂
∂
∂
kl
kl
E
w
old − 
∂
∂
(4)
A. SAVRAN et al.: APPLICATION OF A NEURAL NETWORK FOR ESTIMATING THE CRACK FORMATION ...
454 Materiali in tehnologije / Materials and technology 48 (2014) 4, 453–457
Figure 2: Multilayer feed-forward neural network
Slika 2: Ve~plastno napajanje nevronske mre`e
Figure 1: Artificial neuron model
Slika 1: Model umetne nevronske mre`e
in which  is the learning rate that determines the step
size through the gradient direction. Similarly, the biases
are updated as follows:
b b
E
b
b b
E
b
b b
j j
j
l l
l
k k
new old
new old
new old
= −
= −
= −



∂
∂
∂
∂
∂E
bk∂
(5)
More details regarding the NN training can be found
elsewhere.13
3 EXPERIMENTAL AND NUMERICAL
PROCEDURES
Ce-based solutions were prepared from cerium nitra-
te precursors (Ce(NO3)3 · 6H2O). The powder precursors
were first dissolved in glacial acetic acid, which is used
as a chelating agent. Eleven different amounts of Ce
precursor, including (0.0499, 0.1990, 0.3490, 0.7500,
1.4900, 2.9900, 4.2500, 4.9900, 5.9900, 7.0080,
43.4000) g, which are called as concentration indexes,
were used to estimate the effects on the crack formation
and propagation of the coatings in the solutions as a
processing parameter. The obtained solutions were sub-
sequently diluted with isopropanol. All the solutions
were stirred at room temperature for 60–120 min in
order to yield transparent solutions. The thickness of the
CeO2 thin films measured by ESEM varied between 0.4
μm and 7 μm, depending on the Ce concentration in the
solution. More details regarding the preparations of Ni
tape and CeO2 thin films were given in11,14,15.
In order to investigate the surface morphology, the
crack formation and the propagation of CeO2 films,
in-situ hot stage ESEM was used. For this procedure, Ni
tape substrates were separately dipped into the eleven
different Ce-based solutions at room temperature in air.
Ce-based gel coatings were obtained from this process.
These gel coatings were then placed in the hot-stage
ESEM. The gel coatings were examined in the tempe-
rature range 24 °C to 700 °C for 5 min in vacuum con-
ditions. After placing the Ce-based gel coatings contain-
ing a bubble structure in the ESEM, they were dried and
heat treated from room temperature to 700 °C. The crack
length and the size of the coatings were measured in the
ESEM at several temperatures, i.e., (25, 100, 200, 300,
500, 600 and 700) °C, for a period of 5 min under
vacuum conditions. However, here we utilized the tem-
peratures 400 °C and 700 °C. A computer program of an
in-situ hot-stage ESEM was used to measure the crack
A. SAVRAN et al.: APPLICATION OF A NEURAL NETWORK FOR ESTIMATING THE CRACK FORMATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 453–457 455
Figure 3: The used NN structure
Slika 3: Uporabljena NN-struktura
Figure 4: ESEM microstructures of Ce-based gel and CeO2 coatings
on Ni tape at: a) 25 °C and b) 600 °C11
Slika 4: ESEM-posnetka koloida na osnovi premaza Ce in CeO2 na
traku iz Ni pri: a) 25 °C in b) 600 °C11
size of the coating as a function of concentration index
(Ce content in the solutions) and processing temperature.
In as much as the coating is a gel-like structure at
25–100 °C, and amorphous at 100–350 °C, the CeO2
coating was formed on Ni tape at 420 °C and densified at
700 °C. For these reasons, the crack size of the CeO2
coatings was also measured at 400 °C and 700 °C in
detail. When measuring the crack sizes of the gel and
oxide coatings, ESEM micrographs were taken with the
scale bar 20 μm.
The NN used in this study is shown in Figure 3.
Where v and w represent the weights, and b represents
the biases. The NN has two inputs (u1, u2) and 1 output
(y). There are six neurons in the hidden layer. f(.) and
g(.) are the hidden-layer and the output-layer activation
functions, respectively. The sigmoid-type activation
functions are used by both layers. All the programs for
the NN were developed under Matlab. The inputs of the
NN represent the amount of material and the processing
temperature at which the experiment was performed. The
output of the NN represents the crack sizes of the coat-
ings. The training data set contains 66 elements, includ-
ing the data for temperatures (25, 100, 200, 300, 500 and
600) °C. We left the data for temperatures 400 °C and
700 °C as the testing data sets. Both test sets contain 11
elements. A batch-training procedure was applied to
train the NN using the Steepest Descent (SD) method
and then the crack size of the CeO2 buffer layer was esti-
mated for 400 °C and 700 °C. The details of test data
sets and the NN estimation results are shown in Tables 1
and 2.
4 RESULTS AND DISCUSSION
In-situ, hot-stage ESEM, depending on the tempera-
ture, the crack formation and propagation of the coatings
were evaluated as a function of the Ce content in the
solutions. Figure 4 shows ESEM microstructures of the
Ce-based gel and CeO2 coatings on Ni tape at 25 °C and
600 °C, respectively.11 The ESEM observations revealed
that the crack surface of the Ce-based films changed
slightly with increasing temperatures, from 25 °C to 500
°C, when compared with the other thin films. After CeO2
formation at about 420 °C, microcracks were observed
on the surface. It is interesting to note here that the
bubbles and micro-bubbles were produced on the surface
once the precursor material increased in this solution.
The bubbles showed characteristic properties that caused
the cracks to start on the surfaces of the gel films of the
nitrate-salt-based precursors. Note that there are many
factors influencing the bubbling problem, such as types
of precursors, solvent and chelating agent, viscosity,
dilution, Ce content in solution and so on. It was found
that the size and propagation rate of the cracks increased
with the increasing Ce content. It is clear from the
ESEM observations that the measured values of the
A. SAVRAN et al.: APPLICATION OF A NEURAL NETWORK FOR ESTIMATING THE CRACK FORMATION ...
456 Materiali in tehnologije / Materials and technology 48 (2014) 4, 453–457
Figure 5: The NN estimation and experimental results for CeO2
coatings on Ni tape at 400 °C. In this figure, the x and y axes indicate
the crack size of the coating and the concentration index, respectively.
Slika 5: Dolo~anje NN in rezultati preizkusov za premaz CeO2 na
traku iz Ni pri 400 °C. Na sliki x- in y-os pomenita velikost razpoke in
indeks pogostosti.
Table 1: NN estimated and experimental crack sizes for 400 °C
Tabela 1: Dolo~en NN in eksperimentalna velikost razpok pri 400 °C
NN inputs NN output
Concentra-
tion index
Amount of
precursor
material (g)
Tempera-
ture (oC)
NN
estimated
crack size
(μm)
Experimen-
tal crack
size (μm)
1
2
3
4
5
6
7
8
9
10
11
0.0499
0.1990
0.3490
0.7500
1.4900
2.9900
4.2500
4.9900
5.9900
7.0080
43.4000
400
400
400
400
400
400
400
400
400
400
400
1.6309
1.6980
1.9820
10.2419
14.5338
14.8854
15.2480
15.6821
18.7295
9.6382
33.5412
2.4133
2.4345
2.5000
9.7605
15.0000
15.6250
17.6410
15.7140
25.0500
16.8900
30.5000
Table 2: NN estimated and experimental crack sizes for 700 °C
Tabela 2: Dolo~en NN in eksperimentalna velikost razpok pri 700 °C
NN inputs NN output
Concentra-
tion index
Amount of
precursor
material (g)
Tempera-
ture (oC)
NN
estimated
crack size
(μm)
Experimen-
tal crack
size (μm)
1
2
3
4
5
6
7
8
9
10
11
0.0499
0.1990
0.3490
0.7500
1.4900
2.9900
4.2500
4.9900
5.9900
7.0080
43.4000
700
700
700
700
700
700
700
700
700
700
700
2.2660
2.8976
5.4302
17.2949
18.5257
22.9032
55.6694
71.0257
75.8715
110.6222
142.2619
2.6886
2.6547
3.0000
13.8090
15.6000
17.5000
26.4710
44.4700
71.4200
73.9000
140.0000
crack-propagation rate altered slightly at (300, 400 and
500) °C, whereas at 600 °C they changed considerably.
Figure 5 shows the NN estimation and the experi-
mental results for CeO2 on Ni tape at 400 °C. In this
figure, the x and y axes indicate the crack size of the
coating and the concentration index, respectively. The
concentration index corresponds to eleven type solutions,
including several Ce(NO3)3 · 6H2O contents. As men-
tioned just previously, the Ce(NO3)3 · 6H2O precursors
were dissolved using a solvent and a chelating agent.
The eleven solutions were prepared by changing the
amount of precursors and thus eleven type coatings were
obtained using the hot-stage ESEM. As a result of this,
the crack sizes were measured at 400 °C in the ESEM
and the concentration index was taken as numerical
values, such as from 1 to 11. Actually, the concentration
indexes are not real values indicating the amount of
precursor. It is clear from Figure 5 that the estimated
crack size of the coating is close to the experimental
results for 400 °C. Similarly, the NN estimation and the
experimental results of the crack size for 700 °C are
presented in Figure 6. It is worth noting that the predic-
ted result approaches the mean values when the standard
deviation associated with the measurements is small.
Now that the goal is to fabricate low-cost, high-quality
products in a short time in a modern industry, the
optimization of the processing temperature can be esti-
mated using the NN technique. The key feature of this
research is that the NN approach is useful for the pre-
diction of the cracks size and the crack-propagation rate
during the heat treatment of the sol-gel coatings. This
study can be extended by using other buffer layers,
depending on the sol-gel parameters. This technique is
quite likely to be a key area for the development of all
sol-gel films in the future, prior to the coating processes.
5 CONCLUSIONS
In summary, the NN-based approach was developed
to estimate the crack sizes of a CeO2 coating on a Ni tape
for YBCO-coated conductors, depending on processing
temperatures such as 400 °C and 700 °C using a hot-
stage ESEM. A multilayer feed-forward NN was trained
to estimate the crack size of the coating during process-
ing at temperature in the range 24–700 °C. The experi-
mental measurements of the crack size on the coating on
the Ni tape substrate are a very expensive and a time-
consuming process. The approach proposed in the pre-
sent study provides simplicity and is cost-effective for
preparing the new solutions and subsequent processing
in the sol-gel technique.
Acknowledgements
The authors would like to thank Dr. Y. S. Hascicek
and R. E. Goddard for their assistance in obtaining the
ESEM micrographs at the National High Magnetic Field
Laboratory of Florida State University in Tallahassee,
Florida USA.
6 REFERENCES
1 K. Hornik, M. Stinchocombe, H. White, Neural Networks, 2 (1989),
359–366
2 K. J. Hunt, D. Sbarbaro, R. Zbikowski, P. J. Gawthrop, Automatica,
28 (1992), 1083–1112
3 M. Agarwal, IEEE Control Systems Magazine, 17 (1997), 75–93
4 K. S. Narendra, K. Parthasarathy, IEEE Transactions on Neural Net-
works, 1 (1990), 4–27
5 A. G. Parlos, S. K. Menon, A. F. Atiya, IEEE Transactions on Neural
Networks, 12 (2001), 1411–1432
6 I. H. Mutlu, E. Celik, M. K. Ramazanoglu, Y. Akin, Y. S. Hascicek,
IEEE Transactions on Applied Superconductivity, 10 (2000),
1154–1157
7 E. Celik, I. H. Mutlu, Y. S. Hascicek, IEEE Transactions on Applied
Superconductivity, 10 (2000), 1162–1166
8 E. Celik, H. Okuyucu, I. H. Mutlu, M. Tomsic, J. Schwartz, Y. S.
Hascicek, IEEE Transactions on Applied Superconductivity, 11
(2001), 3162–3165
9 H. Okuyucu, E. Celik, M. K. Ramazanoglu, Y. Akin, I. H. Mutlu, W.
Sigmund, J. E. Crow, Y. S. Hascicek, IEEE Transactions on Applied
Superconductivity, 11 (2001), 2889–2892
10 X. D. Wu, S. R. Foltyn, P. Arendt, J. Townsend, I. H. Campbell, P.
Tiwari, Q. X. Jia, J. O. Willis, M. P. Maley, J. Y. Coulter, D. E. Peter-
son, IEEE Transactions on Applied Superconductivity, 5 (1995),
2001–2006
11 E. Celik, Y. S. Hascicek, Materials Science & Engineering B, 94
(2002), 176–180
12 P. J. Werbos, Beyond Regression, New Tools for Prediction and
Analysis in the Behavioral Sciences, PhD. Dissertation, Harvard
University, Boston, USA, 1974
13 M. S. Al-Haik, H. Garmestani, A. Savran, International Journal of
Plasticity, 20 (2004), 1875–1907
14 E. Celik, J. Schwartz, E. Avci, H. J. Schneider-Muntau, Y. S. Hasci-
cek, IEEE Transactions on Applied Superconductivity, 9 (1999),
2264–2268
15 E. Celik, H. I. Mutlu, E. Avci, Y. S. Hascicek, Advances in Cryoge-
nic Engineering (Materials), 46 (1999), 895–900
A. SAVRAN et al.: APPLICATION OF A NEURAL NETWORK FOR ESTIMATING THE CRACK FORMATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 453–457 457
Figure 6: The NN estimation and experimental results for CeO2
coatings on Ni tape at 700 °C. In this figure, the x and y axes indicate
the crack size of the coating and the concentration index, respectively.
Slika 6: Dolo~anje NN in rezultati preizkusov za premaz CeO2 na
traku iz Ni pri 700 °C. Na sliki x- in y-os pomenita velikost razpoke in
indeks pogostosti.

M. O. SHABANI et al.: OPTIMIZATION OF THE MECHANICAL AND TRIBOLOGICAL PROPERTIES ...
OPTIMIZATION OF THE MECHANICAL AND
TRIBOLOGICAL PROPERTIES OF EXTRUDED AMCs:
EXTENSION OF THE ALGORITHM SEARCHING AREA
VIA MULTI-STRATEGIES
OPTIMIZACIJA MEHANSKIH IN TRIBOLO[KIH LASTNOSTI
EKSTRUDIRANIH AMC: RAZ[IRITEV ALGORITMA ISKALNEGA
PODRO^JA Z UPORABO VE^ STRATEGIJ
Mohsen Ostad Shabani, Ali Asghar Tofigh, Mohammad Reza Rahimipour,
Mansour Razavi, Ali Mazahery, Fatemeh Heydari
Materials and Energy Research Center (MERC), Tehran, Iran
aliasghartofigh60@gmail.com
Prejem rokopisa – received: 2013-01-14; sprejem za objavo – accepted for publication: 2013-10-24
The present article focuses on the development of a comprehensive method for the optimization of the mechanical and
tribological properties of metal-matrix composites using multi-strategy ensemble particle-swarm optimization. An
aluminum-alloy matrix reinforced with coated B4C particles was used for the present study. The cohesion of the reinforcing
ceramic particles represents a very important factor, which is mostly poor at temperatures near the melting point of aluminum
and leads to the inferior mechanical and tribological properties of the developed aluminum matrix composites with a non
-uniform distribution of the reinforcement. The main reason for coating the particles is to improve the bonding between the
reinforcement and the molten alloy and thus to eliminate any interfacial reactions. The great enhancement in the strength values
of the composites in this study can be ascribed to the effective load-bearing capacity of the disintegrated B4C particles, which
are adherently bonded to the matrix alloy. Homogeneity and a reduction in the particle size of the B4C during the extrusion
process is evidenced in the microstructural studies.
Keywords: tribological properties, particle, extrusion process
^lanek je osredinjen na razvoj celovite metode za optimizacijo mehanskih in tribolo{kih lastnosti kompozitov s kovinsko
osnovo z uporabo optimizacije ve~ strategij zdru`evanja gru~ delcev. Za {tudij je bila uporabljena Al-zlitina, oja~ana z
oblo`enimi delci B4C. Kohezija opla{~enih kerami~nih delcev za utrjevanje je pomembna: navadno je slaba pri temperaturah
blizu tali{~a aluminija in povzro~a slab{e mehanske in tribolo{ke lastnosti razvitega kompozita na osnovi Al z neenakomerno
razporeditvijo delcev za oja~anje. Glavni namen opla{~enja delcev je izbolj{anje povezave med delcem za oja~anje in staljeno
zlitino ter odprava reakcij na povr{ini stika. Veliko pove~anje trdnosti kompozita v tej {tudiji se lahko pripi{e u~inkoviti
nosilnosti razpadlih delcev B4C, ki so adherentno vezani na osnovno zlitino. [tudije mikrostrukture so pokazale homogenost in
zmanj{anje delcev B4C med iztiskovanjem.
Klju~ne besede: tribolo{ke lastnosti, delec, postopek iztiskovanja
1 INTRODUCTION
The particle-swarm algorithm tries to simulate the
social behavior of a population of agents or particles, in
an attempt to optimally explore a given problem space.1
At a time instant (an iteration in the optimization con-
text), each particle is associated with a stochastic velo-
city vector, which indicates where the particle is moving
to2–5. The velocity vector for a given particle at a given
time is a linear stochastic combination of the velocity in
the previous time instant, of the direction to the particle’s
best position, and of the direction to the best swarm posi-
tions (for all particles).6 The particle-swarm algorithm is
a stochastic algorithm in the sense that it relies on para-
meters drawn from random variables, and thus different
runs for the same starting swarm may produce different
outputs.7 Some of its advantages are that it is simple to
implement and easy to parallelize.8,9 It depends, however,
on a few of parameters that influence the rate of con-
vergence in the vicinity of the global optimum.10 Overall,
it does not require many user-defined parameters, which
is important for practitioners that are not familiar with
optimization.11–13 Some numerical evidence seems to
show that a particle swarm can outperform genetic algo-
rithms on difficult problem classes, i.e., for uncon-
strained global optimization problems. Moreover, it fits
nicely into the pattern search framework.14,15
During the past decade, novel computational methods
have been introduced in some fields of engineering
sciences, including the solidification and deformation of
metal-matrix composites in materials science.16–18
Aluminum metal-matrix composites (AMCs) are gaining
importance as the most sought-after candidate materials
in the space and automotive industries owing to their
excellent properties, such as superior wear resistance,
low density and high specific stiffness.19–32 Several re-
ports have been published addressing the problems asso-
ciated with their developments, mechanical behavior,
microstructure and the distribution of particulates.33–41
Presently, particulate-reinforced composites are being
Materiali in tehnologije / Materials and technology 48 (2014) 4, 459–466 459
UDK 621.777:539.92 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)459(2014)
produced by several different methods: powder metal-
lurgy, liquid metallurgy, diffusion bonding techniques,
infiltration, squeeze casting, compocasting and spray
deposition techniques.42–48 The most simplified approach
to develop near-net -shaped aluminum-based composites
is by the liquid metallurgy route as it is economical and
can result in mass production.18–20 Generally, these com-
posites consist of a metal matrix, which is melted during
casting, and ceramic reinforcement, which is added to
the molten matrix material by a mechanical stirrer.
However, some challenges need to be addressed in the
development of AMCs to intensify their uses in different
engineering fields, including inferior bond and interfacial
reaction product. These problems do have a direct dete-
riorating effect on the mechanical and tribological
properties of the composites, making them unsuitable for
industrial components.3 These challenges have been
addressed to by the use of coated reinforcement and the
addition of reactive metals like magnesium by several
researchers. In addition, instead of conventional stir-cast-
ing techniques, semi-solid agitation processes can be
employed. The benefits include reduced solidification
shrinkage, a lower tendency for hot tearing, suppression
of segregation, settling or agglomeration and faster
process cycles. These advantages are accompanied by a
lack of superheat (lower operating temperatures) as well
as a lower latent heat, which results in a longer die life
together with a reduced chemical attack of the reinforce-
ment by alloy, also a globular, non-dendritic structure of
the solid phase, which then explains the thixotropic
behavior of the material.18–20 Furthermore, these compo-
cast-coated composites can be subjected to secondary
processing such as extrusion and forging to improvise
upon the mechanical properties in particular strength
coupled with practical ductility.39–43 The purpose of this
study is first to investigate the effects of extrusion and
reinforcing coated particles on the microstructures and
mechanical properties of AA6061 aluminum alloy ma-
trix composites produced by compocasting. Another
objective is to solve the global problems using Multi-
strategy ensemble particle-swarm optimization, which
helps to increase the possibility of industrial application.
2 EXPERIMENTAL PROCEDURE
In this study, composites were produced by the com-
pocasting process using the mechanical mixing of the
AA6061 aluminum matrix, i.e., B4C particles. The
AA6061 aluminum alloy was produced from the mass
fraction w = 99.9 % pure aluminum that had been melted
and then a pure silicon master alloy in mass fractions of
Al–75 % Cr, Al–50 % Cu, and pure magnesium were
added in order. Boron carbide (B4C) in powder form is
used as the reinforcement, having a particle range size of
1–60 μm. We attempted to coat the B4C powders with
TiB2, which helps the incorporation of the particles and
reduces interfacial reactions. Titanium tetraisopropoxide
was selected as a sol-gel precursor and diluted with
ethanol. The boron carbide powders were first dispersed
in the ethanol using a stirrer and then titanium tetraiso-
propoxide and distilled water were added to the stirred
suspension. The processing was conducted at room tem-
perature at a solution pH of 7. The solution was then
aged for 105 min at room temperature, with constant,
gentle stirring. When titanium tetraisopropoxide is used
as the precursor, TiO2 can be produced by hydrolysis and
heat treatment. Since titanium oxide does not support
good wettability, it is converted to TiB2. Figure 1 pre-
sents the X-ray diffraction (XRD) analysis of the coated
powders.
The composites were developed using the stir-cast
method. The process involved melting the alloy in a gra-
phite crucible using an electrical resistance furnace. The
stirrer was positioned just below the surface of the slurry.
The furnace is controlled using a J-type thermocouple
located inside the gas chamber. The temperature of the
alloy was raised to about 680 °C and stirred at (400, 500,
600, 700) r/min using an impeller fabricated from
graphite and driven by a variable ac motor. The stirring
times were noted at (5, 10 and 15) min after the addition
of B4C during the process. Both TiB2 coated and un-
coated boron carbide was varied in proportions of volu-
me fractions (2.5, 5, 7.5, 10, 12.5, 15) %. The tempera-
ture of the furnace was gradually lowered until the melt
reached a temperature in the liquid-solid state (corres-
ponding to a 0.2 solid fraction) while the stirring was
continued. The coated particles were added uniformly at
a rate of 50 g/min over a time period of approximately 3
min. The casting was obtained by pouring the composite
slurry into a steel die placed below the furnace. A con-
tinuous purge of nitrogen gas is used inside and outside
the crucible to minimize the oxidation of the molten
aluminum. The cast matrix alloy and the developed
Al6061–B4C composites (both uncoated and TiB2
coated) were machined to 70 mm diameter and 200 mm
length. The machined billets were then subjected to hot
extrusion using a 200 t hydraulic extrusion press. The
extrusion billets were heated in a muffle furnace for 2 h.
An extrusion ratio of 1 : 10 with a constant ram velocity
(extrusion speed) of 2 mm/s was implemented. The ex-
truded Al6061 alloy and Al6061–B4C composites (both
M. O. SHABANI et al.: OPTIMIZATION OF THE MECHANICAL AND TRIBOLOGICAL PROPERTIES ...
460 Materiali in tehnologije / Materials and technology 48 (2014) 4, 459–466
Figure 1: X-ray diffraction (XRD) analysis of the coated powders
Slika 1: Rentgenska difrakcijska analiza opla{~enega prahu
uncoated and TiB2 coated) were subjected to metallo-
graphic studies, microhardness, tensile and wear tests.
Dry sliding wear tests were performed using pin-on-
disk apparatus, under a load of 10 N and 20 N against a
counterface steel disk of hardness 60 HRC. Cylindrical
specimens of 6 mm diameter and 25 mm height were
used as the test samples. Before the abrasion tests, each
specimen was polished to 0.5 μm. Figure 2 shows a
schematic diagram of the abrasion wear test. The expe-
riment was carried out at room temperature with water as
the lubricant. The wear loss was measured in the steady-
state regime using a linear variable differential trans-
ducer of accuracy 1 μm at the end of 30 min. The wear
rates were calculated from height-loss data. A set of
three samples was tested in every experimental condi-
tion, and the average along with the standard deviation
for each set of three tests was measured. The wear tests
were conducted up to the total sliding distance of 2000
m. The tensile test samples were machined according to
the ASTM E8M standard. Hardness measurements were
carried out using a Shimadzu Microhardness tester with
a load of 1 N for a period of 10 s. For each sample, five
hardness tests on randomly selected regions were per-
formed in order to eliminate the possible segregation
effects and obtain a representative value of the matrix
material hardness. During the hardness measurements,
precaution was taken to make an indentation at a distance
of at least twice the diagonal length of the previous
indention.
3 MODELING
3.1 Particle swarm optimization
The particle swarm optimization (PSO) was origi-
nally designed by Kennedy and Eberhart (Kennedy and
Eberhart, 1995) and has been compared to genetic algo-
rithms for efficiently seeking optimal or near-optimal
solutions in large search spaces.3 It is a new search tech-
nique, which simulates preying behavior among birds.2
In contrast to genetic algorithms, PSO does not need
genetic operations such as crossover and mutation. In-
stead, it changes the individuals by their random veloci-
ties in the solution space.49–54 Compared to evolutionary
algebra, the solution group presents greater randomness
and more benefits, such as faster searching speed, easy
implementation and global optimization.3 In the original
PSO with M particles, each particle is represented as a
potential solution to a problem in a D-dimensional space
and its position at the t-th iteration is denoted as:1
X x x xi i
t
i
t
iD
t= ( , , )1 2 (1)
Each particle remembers its own previous best
position and its velocity along each dimension as:49–52
V v v vi i
t
i
t
iD
t= ( , , )1 2 (2)
The velocity and position of particle i at the (t + 1)th
iteration are updated by the following equations:3
V wV c r p x c r Q Xi j
t
i j
t
i j
t
ij
t
ij
t
i j
t
j
t
i, , . .( ) (
+ = + − + −1 1 1 2 2 , )j
t (3)
X V Xi j
t
i j
t
i j
t
, , ,
+ += +1 1 (4)
where c1 and c2 are two positive constants, known as the
acceleration coefficients; r1 and r2 are two uniformly
distributed random numbers on the range (0.1) for the
j-th dimension of particle i. Vector pi = (pi1t, pi2t ... piDt)
is the position with the best fitness found so far for the
ith particle, which is called the personal best (pbest) posi-
tion. And vector Qi = (Qt1, pt2 ... ptD) records the best
position discovered by the swarm so far, known as the
global best (gbest) position. xi,jt, vi,jt and pi,jt are the jth
dimension of the vector of xi t, vi, t and pi t, respectively.
The parameter w is the inertia weight used for the bal-
ance between the global and local search abilities.
Usually, w decreases linearly with the iteration genera-
tions as:1–4
w w
t w w
T
= −
−
max
max min( )
(5)
where wmax and wmin are the maximum and minimum
weights and usually set to 0.9 and 0.4, respectively. T is
a predefined maximum number of iterations, and t
represents the number of the current iteration.
3.2 Multi-strategy ensemble particle swarm optimiza-
tion (MEPSO)
In MEPSO, all the particles are initially divided into
two parts; we denote them as part I and part II, res-
pectively. The two parts are considered to play different
roles in the search of dynamic environments by using
different strategies, which will be introduced as follows.
The role of part I is considered to search the global
optimum in the current environment as quickly as pos-
sible. Thus, similar operations as the standard PSO are
adopted to guarantee good convergence. Furthermore, a
Gaussian local search is introduced to enhance the local
search ability of part I, which is designed as follows:
At every iteration, for each particle, it has the pro-
bability Pls to perform the Gaussian local search defined
as Eqs. (4) and (6), and has the probability (1 – Pls) to
perform the conventional search defined as Eqs. (3) and
(4). The global best used in part I is the best solution
M. O. SHABANI et al.: OPTIMIZATION OF THE MECHANICAL AND TRIBOLOGICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 459–466 461
Figure 2: Schematic diagram of the abrasion wear test
Slika 2: Shematski prikaz preizkusa abrazijske obrabe
found by all particles, both in part I and part II. The
Gaussian local search is defined as follows:
V ci j
t
,
+ = ∗1 3 gaussrand (6)
where i = 1, 2, ... m, gaussrand is a random number ge-
nerated from a standard normal distribution, c3 is a
positive constant. Although both strategies are designed
for local search, the Gaussian local search adopted in
part I of MEPSO is different from the quantum cloud
defined in MQSO; the distribution in quantum cloud is
uniform while Gaussian is not. The Gaussian local
search strategy defined as Eq. (6) has been testified by
many researches to be a good strategy to enhance the
ability of elaborate search. By performing a local search
with the probability Pls, a particle can search for the
optimum around its current position when it is on the
process of "flying" to the best position found by the
entire swarm.
Therefore, each particle has the chance to search for
its neighborhood, and it might be favorable to find the
optimum in dynamic multimodal environments. The role
of part II is considered to extend the searching area of
the algorithm, and to patrol around the part I to track the
changed global optimum possibly "escaped" from the
coverage of part I. To achieve this purpose, in part II,
each particle has a probability 0.5 to fly to get closer to
the personal best of a particle randomly chosen from part
I, and has probability 0.5 to fly to get farther away from
it. The operator is defined as Eq. (7), we call it differen-
tial mutation in this paper. It is implemented by changing
the direction of a particle’s velocity with a certain
probability. The position of the particle is still renewed
by Eq. (4):
V wV r c r p x
c r
i j
t
i j
t
i j
t
ij
t
ij
t
, , ,( . ) ( )
+ = − + − +
+
1
1 1 2
2
05sgn
3 i j
t
j
t
ij
tG X, ( )−
(7)
sgn =
− <
=
>
⎧
⎨
⎪
⎩⎪
1 0
0 0
1 0
t
t
t
(8)
where Gj is the best solution found by particle a, which
is chosen randomly from part I at each iteration; r1, r2,
r3 are uniformly distributed random numbers in the
interval (0, 1); other parameters are the same as the ones
described in Section PSO. The strategy of differential
mutation may enhance the communication between part
I and part II, extend the particle’s search area, and
prevent the algorithm from being premature.
In part II, particles fly in a way totally different to the
standard PSO. There is no global attractor in part II, the
position of each particle in part II is determined by the
particle in part I via differential mutation strategy (each
particle has a probability of 0.5 to fly to get closer to the
personal best of a particle randomly chosen from part I,
and has probability of 0.5 to fly to get farther away from
it). The purpose of this strategy is to keep the particles in
part II flying around part I to extend the coverage of the
particle population to avoid being trapped into a local
optimum. The roles of the two parts in MEPSO are con-
sidered originally to be different. Part I is designed to
enhance the algorithm’s ability of exploitation, while
part II is designed to enhance the algorithm’s ability of
exploration. The two parts work separately, but particles
in these two parts are also interrelated. On the one hand,
the personal best of particles randomly chosen in part I
are used to compose the new velocities of the particles in
part II, and then influence their relative position with
respect to particles in part I. On the other hand, the best
solution found by part II can be the global attractor of
part I (if it is also the best of the entire swarm), which
will guide the part I fly to the new best (maybe the
changed optimum). The overall algorithm is summarized
as follows.
Step1: Randomize the positions and velocities of all the
particles in the search space. Set all the attractors to a
randomized particle position. Divide all the particles
into two parts. Set part I’s attractor to be the best
position of the entire swarm.
Step2: Evaluate the randomly chosen
Step3: IF the new value is different from the last itera-
tion, re-evaluate the function values at each particle
attractor in part I.
Step4: Re-randomize each particle in part II.
Step5: Update part I’s attractor.
Step6: FOR each particle i in part I IF random number <
Pls THEN, the random number within (0, 1), Apply
Eq. (6) to renew the velocity, perform a local search
ELSE, Apply Eq. (3) to renew the velocity.
Step7: FOR each particle j in part II, Randomly choose a
particle a from part I. Apply Eq. (7) to renew the ve-
locity.//including operator of differential mutation
Step8: FOR each particle j both in part I and part II
Apply Eq. (4) to renew the position.
Step9: Evaluate function at updated position. IF new
value better than particle attractor value THEN, Par-
ticle attractor j: position and value of particle
Step10: IF new value better than part I’s attractor value
THEN, Part I’s attractor: = position and value of par-
ticle
Step11: UNTIL number of function evaluations per-
formed > max
There are mainly two parameters that should be set
before the execution of the algorithm: the proportion of
part I to the whole population Pone, the probability Pls of
performing a Gaussian local search. The first parameter
may be used to control the contribution of part I and part
II to the whole performance of the algorithm, and there-
fore, has an influence on the trade-off between the algo-
rithm’s performance on convergence and diversity main-
tenance. The second parameter may be used to control
the proportion of particles in part I that perform the
Gaussian local search other than the conventional stra-
tegy of PSO, and thus has an influence on the trade-off
M. O. SHABANI et al.: OPTIMIZATION OF THE MECHANICAL AND TRIBOLOGICAL PROPERTIES ...
462 Materiali in tehnologije / Materials and technology 48 (2014) 4, 459–466
between the algorithm’s performance on local search and
global search.
4 EXPERIMENTAL RESULTS
The execution of extrusion in this study results in
compressive stresses and the fracturing of the hard cera-
mic particles within the deforming composites. Figure 3
shows the light microphotographs of extruded volume
fraction of Al6061–12.5 % B4C composites. It should be
noted that a combination of the semi-solid technique and
extrusion in the fabrication of these composites leads to a
reasonably uniform distribution of particles in the matrix
and avoids clustering or agglomeration of the reinforcing
phase. The presence of segregated B4C particles can be
easily recognized in the case of uncoated composites.
Coated B4C particles appear more homogeneous
throughout the extruded matrix alloy. It is assumed that
TiB2 coating improves the wetting kinetics in the liquid
aluminum, which results in a uniform distribution of the
coated B4C particles.22–26
The variation in the hardness of the extruded Al6061
alloy and its composites with the volume fraction of B4C
particles is presented in Figure 4. It is observed that the
hardness of the composite samples increases with an
increase in the B4C content. However, coated composites
exhibit a higher hardness compared to the uncoated ones.
Figures 5 and 6 show the results of the tensile strength
in the extruded Al6061 alloy and its composites. In
general, it is noted that both the yield strength and the
ultimate tensile strength increase with the increasing the
volume fraction of the incorporated B4C particles for all
the materials studied. The improvement in the tensile
strength of the composites is the outcome of a higher dis-
location density and plastic constraint in the matrix. The
strain-hardening of the composites is expected to be
influenced by the dislocation density, the dislocation-to-
dislocation interaction and the constraint of the plastic
M. O. SHABANI et al.: OPTIMIZATION OF THE MECHANICAL AND TRIBOLOGICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 459–466 463
Figure 4: The influence of B4C reinforcement on the hardness of the
composites
Slika 4: Vpliv dodatka delcev B4C na trdoto kompozita
Figure 3: The light micrographs: a) unreinforced Al6061, b)
Al6061-12.5 % uncoated B4C, c) Al6061–12.5 % coated B4C
Slika 3: Mikrostruktura: a) neoja~an Al6061, b) Al6061–12,5 %
neopla{~en B4C, c) Al6061–12,5 % opla{~en B4C
Figure 5: The influence of B4C reinforcement on the yield strength of
the composites
Slika 5: Vpliv dodatka delcev B4C na mejo plasti~nosti kompozita
flow due to the resistance offered by the particles. The
matrix could flow only with the movement of the B4C
particles or over the particles during plastic deforma-
tion.31–39
Given the fact that the maximum solubility of B and
C inside Al melt is not high, these two elements dis-
solute and saturate the melt rapidly in the case of un-
coated composites, which leads to the nucleation of other
products on impurity seeds or at the B4C surface from
the supersaturated melt. However, in the case of coated
composites, most of the initial B4C and Al remain un-
reacted, indicating the phases are conserved for desirable
applications. The XRD pattern of the coated B4C-rein-
forced composites is displayed in Figure 7. It can be
seen that B4C, TiB2 and aluminum are present and no
other reaction products are formed in the system. It is
interesting to note that the rate of improvement in the
yield strength and the ultimate tensile strength of the
coated reinforced composites is higher when compared
with the uncoated samples, which can be attributed to a
number of reasons, including the absence of interfacial
reactions in coated reinforced composites. In addition, a
more uniform distribution of particles and a smaller
inter-particle distance in the case of coated composites
cause the matrix to become considerably constrained and
this results in a higher degree of improvement in flow
stress and UTS.
Figure 8 shows the influence of B4C reinforcement
on the wear rates of the Al6061 alloy. It is clear that the
wear rates of the Al6061 alloy decrease with the addition
of the B4C reinforcement. This improvement in the wear
resistance of the composites can be attributed to the for-
mation of mechanically mixed layers (MML) consisting
of oxides of iron and aluminium during the sliding of
composites on hard steel surfaces. It should be noted that
coated composites exhibit a higher wear resistance than
the Al6061 matrix alloy and the uncoated composites.
The higher hardness, the uniform distribution of particles
through the matrix alloy and the strong interfacial bond
that exists between the matrix and the reinforcement in-
crease the load-bearing capacity and minimize the matrix
contact area in case of the coated B4C reinforced com-
posite. Figure 9 shows the worn surfaces of the unrein-
forced Al alloy and the composites. Extensive cracking
and shearing are observed on the worn surfaces of the Al
alloy. SEM micrographs indicate wider grooves and a
greater extent of the damaged area on the worn surfaces
of the uncoated samples, compared to the coated ones.
5 MODELING RESULTS
For MEPSO, the parameters used in Eqs. (4) and (7)
are: w = 0.25, c1 = c2 = 2.0. Unless stated otherwise, the
probability of a Gaussian local search Pls is set to be
0.15, and the coefficient c3 in Eq. (6) is set to be 0.3. The
proportion of part I Pone = 0.3, and the proportion of part
II (1 – Pone) = 0.7, i.e., the ratio of part I and part II was
3 : 7. When the dimensionality of the solution space in-
creases, the problem becomes more and more difficult
due to the increase in the number of local optima;
therefore, the algorithms are more likely to be trapped
into the local optimum. As can be seen, MEPSO has the
M. O. SHABANI et al.: OPTIMIZATION OF THE MECHANICAL AND TRIBOLOGICAL PROPERTIES ...
464 Materiali in tehnologije / Materials and technology 48 (2014) 4, 459–466
Figure 8: The influence of B4C reinforcement on the wear rates of the
composites
Slika 8: Vpliv delcev B4C na hitrost obrabe kompozita
Figure 6: The influence of B4C reinforcement on the UTS of the com-
posites
Slika 6: Vpliv dodatka delcev B4C na natezno trdnost kompozita
Figure 7: XRD pattern of the coated B4C reinforced composites
Slika 7: Rentgenska difrakcija kompozita z opla{~enimi delci B4C
best performance in all the dimensionality conditions,
and the offline error gap between the MEPSO and other
algorithms becomes larger and larger as the dimension-
ality increases. In the following experiments, for
MEPSO we use the method of re-evaluating five ran-
domly chosen "sentry" particles to detect the change of
the environments. When the changes have been detected,
we re-evaluate each particle’s best position and current
position in part I, and then update each memory with the
better position. The re-randomization mechanism is
applied to all the particles in part II after the environment
changes. The methods of detecting changes and handling
outdated memories of the compared algorithms are the
same as that adopted in1–5. It is clear that MEPSO copes
well with both unimodal and multimodal dynamic pro-
blems. A possible reason may stem from the ensemble of
multiple strategies: from the experimental analysis, it can
be observed that the mechanisms used in part I have a
good effect on the convergence performance of the
algorithm; while the mechanisms adopted in part II can
extend the coverage of the particle population to avoid
being trapped into the local optimum, and to enhance the
ability of catching up with the changing optimum in
dynamic multimodal environments. Figure 10 shows the
effect of iteration number on the global fitness of the
developed model. The numbers of the iteration were
selected to be 2000.
The value of Pls varies from 0.0 to 0.9. Good results
are achieved for 0.1 = Pls = 0.2 in various severity and
multimodal environments. When Pls > 0.4, the perfor-
mance gets rapidly worse as the value of Pls increases. It
is obvious that when the value of Pls is set to be close to
or equal to 0, the results are getting much worse than
those attained for 0.1 = Pls = 0.2. From the experimental
results it is clear that, it is effective to perform a proper
local search for the particles in part I. The value of Pls
must not be too large or too small, i.e., 0.1 = Pls = 0.2 is
the best choice. What is more, the same conclusion can
be drawn for various dynamic problem settings; in other
words, the performance of MEPSO is rather robust to
various dynamic problems with the parameter Pls setting
range from 0.1 to 0.2. The final optimized parameters are
617.48 s stirring time, 644.26 r/min speed of stirrer,
31.43 μm particle size of B4C, 11.53 % volume fractions
of B4C, 87.21 VHN hardness, 0.43 % porosity, 277.21
MPa UTS, 133.57 MPa yield strength, 0.06 10–3 mm3/(N
m) wear rates and 5.73 % elongation. The results show
that the novel technique implemented in this investi-
gation has an acceptable performance. Therefore, this
work shows the usefulness of an intelligent way to pre-
dict the performance of aluminum matrix composites
using Multi-strategy ensemble particle-swarm optimiza-
tion.
6 CONCLUSION
In MEPSO, the whole population of particles is
divided into two parts: part I works as a standard PSO
M. O. SHABANI et al.: OPTIMIZATION OF THE MECHANICAL AND TRIBOLOGICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 459–466 465
Figure 10: The effect of iteration number on the global fitness
Slika 10: Vpliv {tevila ponovitev na globalno zmogljivost modela
Figure 9: Worn surfaces of the unreinforced Al alloy and the com-
posites: a) unreinforced Al6061 alloy, b) Al6061–10 % uncoated B4C,
c) Al6061-10 % coated B4C
Slika 9: Obrabljena povr{ina neoja~ane Al-zlitine in kompozitov: a)
neoja~ana zlitina Al6061, b) Al6061–10 % neopla{~en B4C, c) Al6061
–10 % opla{~en B4C
enhanced with a Gaussian local search strategy, part II
works as a patrol team around part I to extend the search
area of the algorithm, and to catch up with the moving
optimum. It is concluded that extrusion of the fabricated
composites helps to reduce the B4C particles size in both
TiB2 coated and uncoated B4C reinforced composites.
This can be attributed to fact that, the execution of the
extrusion results in compressive stresses and therefore
fracturing of the hard ceramic particles within the de-
forming composites. It should be noted that the particle
distribution in the coated composites is much more
uniform than the uncoated ones. The strong interfacial
bond that exists between the matrix and the reinforce-
ment in the case of coated B4C reinforced composites
contributes significantly to the improved wear resistance
by increasing the load-transfer efficiency between the
matrix and the reinforcement. Furthermore, a poor
interfacial bond between the hard B4C reinforcement and
the soft matrix alloy will lead to the three-body abrasive
wear phenomenon.
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M. O. SHABANI et al.: OPTIMIZATION OF THE MECHANICAL AND TRIBOLOGICAL PROPERTIES ...
466 Materiali in tehnologije / Materials and technology 48 (2014) 4, 459–466
P. MO^NIK, T. KOSEC: TRIBO-CORROSION PROPERTIES OF A NiTi DENTAL WIRE
TRIBO-CORROSION PROPERTIES OF A NiTi DENTAL
WIRE
TRIBOKOROZIJSKE LASTNOSTI DENTALNE @ICE NiTi
Petra Mo~nik, Tadeja Kosec
Slovenian National Building and Civil Engineering Institute, Dimi~eva 12, 1000 Ljubljana, Slovenia
petra.mocnik@zag.si
Prejem rokopisa – received: 2013-06-13; sprejem za objavo – accepted for publication: 2013-10-04
NiTi-alloy archwires are used in dental medicine for tooth positioning. Failures are reported during the mounting and operation.
It is supposed that these difficulties are results of a simultaneous presence of corrosion and mechanical wear. First, a corrosive
medium was examined in order to simulate the conditions in the mouth. Different simulated body fluids were compared with
natural saliva using electrochemical methods. The corrosion properties of the NiTi dental wire in the as-received state and
without the surface oxide film were studied with electrochemical impedance spectroscopy. Tribo-corrosion tests of NiTi in
artificial saliva were performed and a relation between the chemical and mechanical wear was studied. After the experiment, the
surface was spectroscopically examined. The relation between the chemical and mechanical wear was determined.
Keywords: NiTi, simulated saliva, passive film, tribo-corrosion, electrochemical impedance spectroscopy
@ice zlitine NiTi se uporabljajo v dentalni medicini za oblikovanje zobnega loka. Med namestitvijo in uporabo se ve~krat
pojavijo pretrgi. Domneva se, da je vzrok so~asna korozija in mehanska obraba. Raziskava je bila najprej usmerjena v iskanje
primernega korozivnega medija, ki nadome{~a razmere v ustih. Z elektrokemijskimi metodami smo primerjali ve~ razli~nih
sinteti~nih slin z vzorcem naravne sline. Korozijski preizkusi dentalne `ice NiTi z oksidno plastjo (dostavljen vzorec) v
primerjavi z `ico brez oksidne plasti so bili izvr{eni z elektrokemijsko impedan~no spektroskopijo. Tribokorozijski preizkusi v
umetni slini so pokazali odnos med kemijskim in mehanskim prispevkom obrabe. Po eksperimentu je bila povr{ina spektro-
skopsko pregledana. Dolo~ena je bila odvisnost med kemijsko in mehansko obrabo.
Klju~ne besede: NiTi, umetna slina, pasivna plast, tribokorozija, elektrokemijska impedan~na spektroskopija
1 INTRODUCTION
A NiTi alloy is known as a shape-memory alloy. It is
able to change its shape from a deformed to its original
shape if it is heated or cooled in a certain temperature
range. In general, the atomic composition of the alloy is
50 % Ni and 50 % Ti. A dental NiTi alloy wire is sub-
jected to the mechanical and corrosive wear, known as
tribo-corrosion processes. Tribo-corrosion is a complex
process and its parameters include corrosion, deforma-
tion, friction and mechanical wear. The total wear of a
material within a tribo-corrosion system is a combination
of all these factors. Destructive effects of one factor on
another one cannot be controlled and they are difficult to
predict. The corrosion environment of dental wires in-
cludes several factors (saliva, food, drink), and, occa-
sionally, also chemical abrasives (toothpaste, gels and
mouthwashes). It is difficult to include all the factors and
influences in the electrochemical experiments, so the
research has been limited to the factor that is constantly
present – the saliva.
In most cases the studies on a NiTi alloy are electro-
chemical researches.1–5 There are some researches that
investigate different surface finishes4,6–9 with the aim to
improve the surface since nickel is known to be a poten-
tially allergenic material. There are some studies that
investigate the problems of pitting1 and crevice corro-
sion2,3,6 on the NiTi alloy.
Since the NiTi alloy is primarily used for biomedical
purposes, the electrochemical research of the material
properties often uses a variety of simulated human fluids.
For the studies of the biomaterial properties in the body,
different simulated body fluids are used like: simulated
body fluid (SBF),7 Hank’s solution1,3,4,8 and solutions of
NaCl with different concentrations.2,9 For the corrosion
studies of dental materials, several different artificial
salivas were used such as the artificial salivas according
to Fusayama5,10,11 and Duffo et al.12
There are very few studies of the NiTi alloy with an
emphasis on tribo-corrosion. Only tribological properties
of NiTi alloys are reported.13,14 Zhang and Farhat13 found
that the NiTi-alloy wear resistance is about 30 times
greater than that of pure titanium, and about 10 times
greater than that of pure nickel. Abedini and colleagues14
tested the tribological behaviour of a NiTi alloy in the
austenitic and martensitic phase forms. They noted that
plastic deformation dominates in the martensitic phase
and fatigue wear dominates in the austenite phase.
The aim of the present study is to investigate the elec-
trochemical properties of a NiTi alloy in an artificial sali-
va to learn about the corrosion properties of a dental wire
and to define the effect of the wear on the corrosion pro-
perties of the dental wire.
Materiali in tehnologije / Materials and technology 48 (2014) 4, 467–472 467
UDK 620.193:669.24'295 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)467(2014)
2 EXPERIMENTAL WORK
The following types on the NiTi alloy were used in
the study:
• a NiTi dental alloy (a 3M orthodontic wire, super-
elastic, 0.48 mm × 0.64 mm, Orthoform III Ovoid,
Upper) for EIS (electrochemical impedance spectro-
scopy) and tribo-corrosion;
• a NiTi sheet (alloy BB NiTi, 2 mm sheet, super-
elastic, flat annealed, surface-oxide free (pickled),
Memry GMBH) for the basic electrochemical study.
NiTi dental-wire samples were investigated for two
different surface finishes, first in the as-received state
(with the suppliers’ finish) and later the surface was
abraded with the 1200-grid SiC paper.
Before measuring, the samples were ultrasonically
cleaned in ethanol, washed with distilled water and then
well dried. For the selection of the most appropriate
simulated saliva, natural saliva was compared to six
different compositions of simulated saliva proposed by
various authors. Duffo12 (0.6 g/L NaCl, 0.72 g/L KCl,
0.22 g/L CaCl2·2H2O, 0.68 g/L KH2PO4, 0.856 g/L
Na2HPO4·2H2O, 0.06 g/L KSCN, 1.50 g/L KHCO3,
0.03 g/L citric acid); Ericsson12 (0.584 g/L NaCl, 0.34
g/L KH2PO4, 1.50 g/L KHCO3, 0.029 g/L citric acid,
0.34 g/L Na2HPO4, 0.166 g/L CaCl2, 0.014 g/L MgCl2);
Hank12 (8.0 g/L NaCl, 0.4 g/L KCl, 0.6 g/L KH2PO4,
0.06 g/L Na2HPO4·12H2O, 0.10 g/L MgCl2·6H2O, 0.35
g/L NaHCO3, 0.06 g/L MgSO4·7H2O, 1.0 g/L glucose,
0.60 g/L Na2HPO4, 0.14 g/L CaCl2); Fusayama11 (0.4 g/L
NaCl, 0.4 g/L KCl, 0.906 g/L CaCl2·2H2O, 0.69 g/L
KH2PO4, 0.05 g/L Na2S·9H2O); simulated body fluid
(SBF)15 (8.18 g/L NaCl, 0.22 g/L KCl, 0.13 g/L KH2PO4,
0.07 g/L Na2SO4, 0.35 g/L NaHCO3, 0.36 g/L CaCl2,
0.30 g/L MgCl2); NaCl + lactic acid12 (5.844 g/L NaCl,
9.008 g/L lactic acid).
The natural-saliva collection was conducted in two
consecutive days. During collecting the saliva was stored
in a refrigerator at 2 °C. The donor of the saliva did not
eat any food or drink for 1 h before, and during the
collection. The procedure took place more than 1 h after
the last teeth cleaning. To stimulate the saliva secretion,
the donor was asked to chew a parafilm and drink only
water.
The electrochemical testing was executed in a
three-electrode corrosion cell with a SCE reference
electrode and a graphite counting electrode. After 2 h
stabilization at the open-circuit potential (OCP) poten-
tiodynamic measurements were preformed starting at
–0.25 V vs. the OCP and progressing up to +2.0 V at the
scan rate of 1 mV/s. A Gamry Instruments potentiostat/
galvanostat (ZRA Reference 600, USA, 2006) was used.
For EIS measurements an Autolab PGSTAT100
potentiostat/galvanostat, with a NOVA 1.6 module was
used. The frequency scan ranged from 65 kHz to 1 mHz
at 10 points per decade with an AC amplitude of ± 10
mV. The absolute impedance and phase angle were
measured at each frequency. The impedance measure-
ments were carried out at the open-circuit potential
(OCP) at different times of the immersion (2, 8, 24 and
72) h in the electrolyte. All the measurements were
conducted in a Duffo simulated saliva. The impedance
data were interpreted on the basis of the equivalent
electrical circuits, using the Zview (Scribner) program
for fitting the experimental data.
The tribo-corrosion tests were performed with a reci-
procal tribometer (Tribotechnic, a pin-on-disc and reci-
procating tribometer, 2009, France) in a Teflon cell with
a platinum wire as the counter electrode and Ag/AgCl as
the reference electrode. The counter body was a Al2O3
ball 6 mm. The length of the wear track was 10 mm, the
sliding speed was 5 mm/s and the loads were 1 N and
2 N.
3 RESULTS AND DISCUSSION
3.1 Selection of an appropriate corrosion media
Figure 1 shows a potentiodynamic measurement of a
NiTi alloy in different artificial saliva solutions and in
natural saliva. The corrosion potential, Ecorr, and the cur-
rent density, jcorr, were defined by adjusting the tangents
of the anodic and cathodic parts of the potentiodynamic
curve in the Tafel region. The corrosion current density
of the NiTi alloys in natural saliva was 46 · 10–9 A cm–2.
The most similar result was obtained for the saliva pro-
posed by Duffo,12 where the current density was 70 · 10–9
A cm–2. Higher corrosion current densities were obtained
for the other artificial salivas (Table 1). The current den-
sities in the passive regions were similar in all the tested
solutions, being approximately 3 · 10–6 A cm–2.
The breakdown potential value, Eb, was the same in
Duffo’s solution and in natural saliva, 1.1 V. The width
of the passive area for the NiTi samples in natural saliva
and in Duffo’s solution were similar, 1.08 V for natural
saliva and 1.06 V for Duffo’s solution. All the other inve-
stigated solutions had their breakdown potentials at
higher potential values and their widths of the passive
areas differed from the values obtained in natural saliva.
The highest values of the breakdown potential were
measured in the solution of lactic acid with an addition
of NaCl (Eb = 1.40 V) and in Fusayama’s artificial saliva
(Eb = 1.28 V).
After comparing the measurement results for the NiTi
alloy in natural saliva with the results obtained in the
artificial salivas, it is evident that the artificial saliva
solution proposed by G. S. Duffo12 shows the greatest
similarity to natural saliva. The artificial saliva prepared
with Duffo’s procedure was used for further electroche-
mical measurements.
3.2 Electrochemical impedance spectroscopy
The electrochemical-impedance-spectroscopy results
for the NiTi dental wire in a simulated saliva solution at
P. MO^NIK, T. KOSEC: TRIBO-CORROSION PROPERTIES OF A NiTi DENTAL WIRE
468 Materiali in tehnologije / Materials and technology 48 (2014) 4, 467–472
different immersion times are presented as Nyquist plots
and Bode diagrams in Figures 2 and 3.
The impedance spectra consist of a high-frequency
intercept with the abscise axis and the main-frequency
semicircle. It can be seen at different immersion times on
Figures 2 and 3 that the impedance response increases
with the time. There is a small difference in the profiles
of the high- and medium- frequency regions, but a great
difference in the responses at the lower frequencies.
The impedance spectra were fitted with the Randles
equivalent circuit, as presented in Figure 3. The equiva-
lent circuit consists of the resistance and capacitance
elements due to the oxide film (RC) that are in the series
with Re, representing the electrolyte resistance.
R and Q represent the properties of the outer porous
and passive film/solution interface reactions. The Q sym-
bol signifies the possibility of a non-ideal capacitance (a
constant-phase element, CPE) with n varying from 0.939
to 0.952 for the impedance data at different immersion
times. The impedance of the CPE is given by:16
Q = ZCPE()= [C(j)
n]–1 (1)
For n = 1, the Q element is reduced to a capacitor
with the C capacitance and for n = 0, to a simple resistor.
The values of the fitted parameters of the equivalent
circuit at four different immersion times and different
potentials are presented in Table 2.
P. MO^NIK, T. KOSEC: TRIBO-CORROSION PROPERTIES OF A NiTi DENTAL WIRE
Materiali in tehnologije / Materials and technology 48 (2014) 4, 467–472 469
Figure 1: Potentiodynamic curves for the NiTi alloy in different
simulated body solutions: a) Fusayama’s saliva, SBF and Hank’s
solution compared to natural saliva, b) Ericsson’s saliva, NaCl with
the lactic acid and Duffo’s saliva compared to natural saliva, at a scan
rate 1 mV/s
Slika 1: Potenciodinami~ne krivulje za zlitino NiTi v razli~nih umet-
nih slinah: a) slina Fusayama, SBF in Hank primerjava z naravno sli-
no, b) slina Ericsson, NaCl z mle~no kislino in Duffo-slina, primerjava
z naravno slino, hitrost preleta 1 mV/s
Table 1: Electrochemical parameters for the NiTi alloy in different
simulated saliva solutions compared with natural saliva
Tabela 1: Elektrokemijski parametri za zlitino NiTi v razli~nih
raztopinah umetne sline v primerjavi z naravno slino
Ecorr/
V
jcorr /
(A cm–2)
Eb/
V
(Eb–Epp)/
V
Natural saliva –0.30 46 · 10–9 1.10 1.08
Fusayama’s saliva –0.13 78 · 10–9 1.28 1.00
Ericsson’s saliva –0.21 114 · 10–9 1.10 0.90
SBF –0.25 81 · 10–9 1.15 1.04
Hank’s solution –0.23 147 · 10–9 1.15 1.03
Lactic acid + NaCl –0.16 124 · 10–9 1.40 1.20
Duffo’s saliva –0.32 71 · 10–9 1.10 1.06
Figure 3: Nyquist and Bode plots for the NiTi dental wire without an
oxide film at the OCP for different immersion times in the simulated
saliva
Slika 3: Nyquistov in Bodejev diagram za dentalno `ico NiTi brez
oksidne plasti pri OCP po razli~nih ~asih namakanja v umetni slini
Figure 2: Nyquist and Bode plots for the NiTi dental wire in the
as-received state at the OCP for different immersion times in the simu-
lated saliva
Slika 2: Nyquistov in Bodejev diagram za dentalno `ico NiTi v do-
stavljenem stanju pri OCP po razli~nih ~asih namakanja v umetni slini
Table 2: Values of the fitted parameters of the equivalent circuit as a
function of the applied potential at different immersion times
Tabela 2: Vrednosti prilagojenih parametrov nadomestnega vezja v
odvisnosti od uporabljenega potenciala pri razli~nih ~asih namakanja
NiTi wire in the
as-received state CPE n
Rp/
(k cm2)
C/
(F cm–2)
2 h 2.25 · 10–5 0.939 517 2.64 · 10–5
8 h 2.10 · 10–5 0.941 704 2.49 · 10–5
24 h 1.91 · 10–5 0.941 1104 2.31 · 10–5
72 h 1.65 · 10–5 0.943 11250 1.99 · 10–5
NiTi without the
oxide film CPE n
Rp/
(k cm2)
C/
(F cm–2)
2 h 1.79 · 10–5 0.938 1470 2.22 · 10–5
8 h 1.64 · 10–5 0.947 4470 2.09 · 10–5
24 h 1.57 · 10–5 0.951 9330 2.03 · 10–5
72 h 1.50 · 10–5 0.952 20400 2.00 · 10–5
The Re parameter has a value from 15  cm2 to 24
 cm2 and it is ascribed to the electrolyte resistance. The
R values are increasing with the immersion time,
showing that the oxide film on the NiTi alloy has an
increasing resistance. The CPE, denoted as Q, was
recalculated using equation17 C = [R1 – n Q]1/n in order to
compare the capacitance values for the NiTi dental alloy
at different immersion times.
A decrease in the C value can be attributed to the
thickening of the oxide layer. The thickening effect is
larger in the case of the as-received oxide film than in the
case of the oxide film naturally grown during the immer-
sion in the simulated saliva. A similar effect was already
observed in18–20.
Based on equation:
d = () (0) (A) / C (2)
where  is the value for the dielectric constant, d is the
thickness of the film, 0 is 8.85 · 10–14 F/cm and A is the
surface area (cm2), the thickness of the oxide layer can
be estimated. The decrease in capacitance points occurs
with the increase in the thickness of the passive layer.
This thickness growth is larger at the early immersion
times at the OCP.
Assuming the value of 48 for the dielectric con-
stant,21 the thickness determined with EIS can be calcu-
lated. The calculated thickness of the oxide film on the
NiTi wire is approximately 1 nm and in the case when
the oxide film was abraded during the 2 h immersion, the
thickness is 2 nm.
The measured impedance spectra and the recalcu-
lated values show that the polarization resistance of the
oxide films, immersed in a simulated saliva increases
with the time of immersion. It is larger for the freshly
polished oxide films on the NiTi wires than in the case of
the as-received NiTi wires. Also, the impedance results
show that the former oxide layer might be thinner and
less protective than the oxide film formed during the
immersion in a simulated saliva solution.
Even if the oxide layer on the NiTi wire is damaged,
a quick repassivation is expected.
3.3 Tribo-corrosion
Due to a slightly more demanding form of the wire
(the wire width of 0.64 mm) and the width of the wear
track made with the 2 N force, which was up to 0.6 mm,
the experiments with higher loads were not conducted.
Figure 4 shows a tribo-corrosion experiment on the
dental NiTi wire worn by the 1 N and 2 N forces. The
coefficients of friction (COF) were high,  0.7, and the
increase was intense in the first 50 s; after that the value
of the COF was fairly constant. Near the end of the
rubbing period, the COF value was registered and it was
0.75 for the load of 1 N and 0.65 for the load of 2 N. The
COF fluctuations were smaller when the load of 2 N was
applied. During the mechanical wear with 2 N load, the
recorded potential was lower by 0.08 V than in the case
of the test with the 1 N load. It was found that during the
mechanical wear the potential and the coefficients of
friction were stable and the fluctuation window was
fairly narrow. These findings can also be attributed to the
homogeneous surface. After the mechanical wear in the
tribo-corrosion tests, the repassivation was fast, appro-
P. MO^NIK, T. KOSEC: TRIBO-CORROSION PROPERTIES OF A NiTi DENTAL WIRE
470 Materiali in tehnologije / Materials and technology 48 (2014) 4, 467–472
Figure 4: Monitoring of the: a) electrochemical potential, b) coeffi-
cient of friction and force vs. time during the tribo-corrosion test on
the dental NiTi wire in a Duffo solution
Slika 4: Spremljanje: a) elektrokemijskega potenciala, b) koeficienta
sile trenja v odvisnosti od trajanja tribokorozijskega preizkusa `ice
NiTi v Duffo-raztopini
ximately 30 s, showing fairly good repassivation pro-
perties of the NiTi dental alloy.
In Figure 5, a SEM image presents the wear track on
the NiTi alloy after the tribo-corrosion test with the load
of 1 N. The wear track is hardly visible, most likely due
to the homogeneous surface and the large hardness of the
NiTi alloy.13 The hardness of the NiTi dental alloy in the
wear track was about 400 HV. The width of the wear
track depending on the normal load was 106 μm for the 1
N load. In the case of the higher load of 2 N (not shown)
the width of the wear track was about 135 μm.
We tried to define the prevailing wear mechanism by
observing the characteristic changes in the surface topo-
graphy and the microstructure of the thin zone bellow the
rubbing surface. The wear particles were not found
during the mechanical wear of the dental alloy. The wear
track was smoothly worn, as already indicated by the
small changes in the coefficient of friction. The SEM
image did not give any additional information on the
possible wear-mechanism properties. Since a consider-
able electrochemical change was detected by measuring
the corrosion potential, one can assume that an oxidative
wear as a tribo-corrosion phenomenon took place, com-
bining the mechanical wear and oxidation.
4 CONCLUSIONS
A comparison between different solutions and natural
saliva was made to study and determine realistic condi-
tions. Furthermore, the corrosion and tribo-corrosion
properties of a NiTi dental alloy in artificial saliva were
studied.
Different solutions simulating the body fluids were
tested with electrochemical techniques in order to find a
solution that most closely represents the appropriate
effects of natural saliva on the electrochemical properties
of the NiTi alloy.
From the electrochemical impedance spectroscopy
results it was clear that a NiTi-alloy archwire has a good
corrosion resistance. By comparing two samples with
different surfaces it could be concluded that the wire
without the oxide film has better corrosion properties
than the wire in the as-received state.
In the tribo-corrosion study it was found that the
effect of the mechanical wear is fairly strong and the
coefficients of friction are high. Heavy pressures on the
surface accelerate the corrosion of the NiTi dental wire.
It can be concluded that the mechanical wear also has
a great effect on the NiTi dental wire used in realistic
conditions. It was evidenced that the NiTi alloy has good
repassivation abilities.
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P. MO^NIK, T. KOSEC: TRIBO-CORROSION PROPERTIES OF A NiTi DENTAL WIRE
472 Materiali in tehnologije / Materials and technology 48 (2014) 4, 467–472
M. BOROUNI et al.: EFFECT OF A NANO-CERAMIC MOLD COATING ON THE FLUIDITY LENGTH ...
EFFECT OF A NANO-CERAMIC MOLD COATING ON
THE FLUIDITY LENGTH OF THIN-WALL CASTINGS IN
Al4-1 ALLOY GRAVITY SAND CASTING
VPLIV NANOKERAMI^NEGA PREMAZA PE[^ENE FORME NA
TEKO^NOST V TANKOSTENSKIH ULITKIH IZ ZLITINE Al4-1 PRI
GRAVITACIJSKEM ULIVANJU
Mansour Borouni, Behzad Niroumand, Mohammad Hosein Fathi
Department of Materials Engineering, Isfahan University of Technology, 8415683111 Isfahan, Iran
m.mahmoudsalehi@ma.iut.ac.ir
Prejem rokopisa – received: 2013-06-27; sprejem za objavo – accepted for publication: 2013-10-29
The thin-wall casting of aluminum alloys provides new opportunities for aerospace industries in producing lightmass structures
with good mechanical properties. The fluidity of molten metal is very important in producing sound castings, particularly
thin-wall castings. In the gravity casting of thin-wall castings of the Al4-1 alloy, the melt that fills the mold with the effect of
self-weight, does not have sufficient fluidity, so causing defects such as cold shut, misrun and shrinkage defects during the
solidification. These defects reduce the quality and mechanical properties of the cast parts. The casting fluidity depends on
many factors, such as temperature, melt composition, casting part shape and mold coating. This study reports the influence of a
nano-ceramics mold coating, a micro-ceramic mold coating, a graphite mold coating and the section thickness of the part on the
fluidity length of Al4-1 alloy castings. The results showed that by applying nano-ceramic mold coatings on the surface of a sand
mold, the fluidity length and the surface quality of the castings can be improved.
Keywords: Al4-1 alloy, nano ceramic coating, micro ceramic coating, fluidity length, thin wall castings
Tankostenski ulitki iz aluminijevih zlitin so nova prilo`nost za letalsko industrijo pri proizvodnji lahkih ulitkov z dobrimi
mehanskimi lastnostmi. Teko~nost staljene kovine je pomembna pri proizvodnji ulitkov brez napak, posebno {e tankostenskih.
Pri gravitacijskem ulivanju tankostenskih ulitkov iz zlitine Al4-1 talina zapolni formo zaradi lastne te`e, ~e pa nima dovolj
teko~nosti, to povzro~i nastanek napak, kot so hladni zvari, slabo te~enje, napake zaradi kr~enja pri strjevanju. Omenjene
napake zmanj{ujejo kvaliteto in mehanske lastnosti ulitkov. Livnost zlitine je odvisna od mnogih dejavnikov, kot so temperatura,
sestava taline, oblika ulitka in prevleka forme. Predstavljen je vpliv nanokerami~nega premaza, mikrokerami~nega premaza,
grafitnega kerami~nega premaza in lokalne debeline na teko~nost v ulitkih iz zlitine Al4-1. Rezultati so pokazali, da je mogo~e z
uporabo nanokerami~nega premaza na pe{~eni formi izbolj{ati teko~nost in kvaliteto povr{ine ulitka.
Klju~ne besede: zlitina Al4-1, nanokerami~ni premaz, mikrokerami~ni premaz, teko~nost, tankostenski ulitki
1 INTRODUCTION
Aluminum is one of the most important non-ferrous
metals. Al–Si alloys are the most widely used aluminum
alloys due to their castability, high strength-to-mass
ratio, corrosion resistance, etc.1 The industrial demand
for thin-wall castings in aluminum alloys is of great im-
portance in order to produce light components, which
enable an increased payload and reduced energy con-
sumption in aerospace applications. The Al4-1(Ak9pch)
alloy is an aluminum-silicon alloy with 9–10.5 % silicon,
0.25–0.35 % manganese and 0.23–0.3 % magnesium.
This alloy is used for reducing mass and good mecha-
nical properties, and is employed for manufacturing
many airplane and aerospace parts with complicated
shapes and thin-wall castings.2–6
However, thin-wall castings of Al4-1 alloy can pose
manufacturing problems associated with mold filling.
The rapid cooling of thin-wall sections of the casting
reduces the fluidity of the molten metal, which could
cause the molten metal to prematurely freeze before it
can completely fill the mold cavity, resulting in an in-
complete fill or cold shuts. Hence, the fluidity of the
melt is an important concern in the foundry for thin-wall
castings and the production of thin-wall castings is
limited by the fluidity of the molten metal. The fluidity
of aluminum alloys has a direct influence, not only on
the material’s castability, but also on the casting
mechanical properties of the casting part. The fluidity of
molten metals, in the foundry environment, is defined as
the length the metal flows before it is stopped due to
solidification. This is because the fluidity limits the geo-
metry of a casting that can be successfully filled. The
study of fluidity is particularly important for the aerospa-
ce and automotive industries in order to produce thinner
and lighter products. Therefore, in recent years, many
foundries and metal suppliers have invested time and
money in studying the fluidity of their foundry alloys.
Fluidity is a complex parameter that is affected by the
properties of the molten metal and mold (such as the
mold coating), the pouring conditions and the solidifica-
tion mechanism.4,7–9 The mold coating significantly im-
proves the fluidity. It also improves, indirectly, the
mechanical properties of the cast products because the
mold coating significantly increases the fluidity and,
Materiali in tehnologije / Materials and technology 48 (2014) 4, 473–477 473
UDK 669.715:621.74 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)473(2014)
hence, a lower casting temperature can be used.8 The
surface roughness with the production of frictional for-
ces reduces the flow of the melt, and consequently redu-
ces the fluidity of the melt. In addition, various experi-
ments show that the mold-coating materials will produce
smooth surfaces and reduce the wetting characteristics of
the mold at the metal-mold interface and reduce the
friction and the melt-mold contact, therefore reducing
the heat-transfer coefficient (HTC) at the metal-mold
interface, and so the fluidity is substantially increased.8,10
The most common application of nano-materials is
nano-ceramic particles. Nano-ceramics are ceramic parti-
cles with dimensions in the nanometer rang.11 The nano-
mold coating has a proper surface, anti-sticking and a
low coefficient of friction. The roughness is related to
the grain size of the mold coating. The apparent fluidity
is represented by the flow distance in the test mold and
was found to be increased significantly by the mold
coating.6,12
In this paper we investigate the effect of nano- and
micro-ceramics and graphite mold coating on the fluidity
length and the soundness of thin-wall castings of the alu-
minum alloy Al4-1 by gravity casting in a sand mold.
2 EXPERIMENTAL PROCEDURE
2.1 Materials
The standard as well as the actual chemical compo-
sition of the Al4-1 alloy are shown in Table 1. A 180 kg
crucible resistance furnace was used to melt the alloy. At
first, 170 kg of Al4-1 alloy was put into the furnace.
Then, additions of 0.05 % of melt-mass beryllium to
aluminium-5 % beryllium master alloy in the ingot form
at the start of the melting treatment. Then the alloy was
melted and heated to 740 °C and kept for 10 min, and a
mobile rotary degassing system was used. The impeller,
made of a graphite pipe, was immersed into the melt and
argon was injected into the melt for 20 min through the
pipe. Then a RE- containing flux was injected into the
melt for 10 min through the pipe. The rotating speed of
the impeller was maintained at 350 r/min. The main
composition of the flux used in this experiment can be
seen in Table 2. For melting qualification treatment, the
addition of 0.05 % of melt mass strontium in rod form,
to the aluminum-10 % strontium master alloy. The stron-
tium master alloy was added to the melt approximately 3
min before the end of the fluxing, because it has the best
effect and fewest losses.
The addition of 0.2 % of melt mass, titanium with
Al-5Ti-1B master alloy in rod form, simultaneously with
strontium master alloy 3 min before the end of the
fluxing treatment. The titanium master alloy was added
at the end stage of fluxing, because without the loss of
grain refiner, refinement of the structure.5,13,14
The sand molds were made from sodium-silicate
bonded, CO2-cured sand. After the preparation of the
mold, a nano-ceramic (MM12), graphite or micron
(ZR1) mold coating was applied to the sand mold surface
using the spray gun and the mold surface was heated
with a natural gas torch to harden the mold coating. The
nano-ceramic mold coating MM12 was product by ltN
Nanovation AG (Germany) in order to study the effects
of mold coating and casting section thickness on the
fluidity. The mold cavity was designed with four straight
channels, 250 mm long, with rectangular cross-sections
of (2, 4, 6, 8) mm thickness and 30 mm across. A mecha-
nical drawing sample for fluidity test, a 3D model and
one cast sample are shown in Figure 1.
The pouring basin that is shown in Figure 2 was used
for reducing the turbulence of the melt and to unify the
velocity of the melt pouring. This basin was made using
AFS (American foundry men society) supervision.15
M. BOROUNI et al.: EFFECT OF A NANO-CERAMIC MOLD COATING ON THE FLUIDITY LENGTH ...
474 Materiali in tehnologije / Materials and technology 48 (2014) 4, 473–477
Figure 1: a) Mechanical drawing of sample for fluidity test, b) 3D
model for fluidity test and c) one cast sample
Slika 1: a) Risba vzorca za preizkus teko~nosti, b) 3D-model za preiz-
kus teko~nosti in c) ulit vzorec
Table 1: Chemical composition of the Al4-1alloy in mass fractions
(w/%)
Tabela 1: Kemijska sestava zlitine Al4-1 v masnih dele`ih (w/%)
Elements Standard Actual
Si 9–10.5 10
Mn 0.25–0.35 0.3
Mg 0.23–0.3 0.3
Ti 0.08–0.15 0.1
Others 0.6 0.6
Al Bal. Bal.
Table 2: Composition of flux used in this research
Tabela 2: Sestava talil, uporabljenih pri tej raziskavi
Composition Content (w/%)
NaCl 45
KCl 45
LaF3 5
Other composition 5
2.2 Fluidity evaluation
In order to identify the effect of the mold coating and
casting cross-sections on the fluidity, 12 samples were
cast. The molten metal was poured manually from the
furnace into the pouring basin with a ladle. The tempe-
rature of the metal was measured by the operator with a
calibrated thermocouple (±1 °C accuracy) in the ladle,
before pouring into the basin. The aim was to pour the
molten metal into the basin as fast as possible and to fill
up the basin completely, in order to have the same initial
metallostatic pressure head on the flowing metal.
Table 3 shows the casting conditions of the samples.
In this study, a temperature of 625 °C was selected for
pouring the melt after some preliminary tests, because at
the mentioned temperature all of the channels have
misrun. Therefore, trying with a change of the coating
type, improves the fluidity length and the castability, for
better filling of the channels. For a phase analysis of the
MM12 coating, the amount of coating drying and pow-
dered then carry out X-ray diffraction (XRD) using a
Philips Xpert-MPD model on a coating and show that the
mentioned coating contains ZrO2, SiO2 and Al2O3 pha-
ses.
Table 3: Casting condition of various samples (for each condition, 3
samples were cast in a sand mold and 625 °C pouring temperature)
Tabela 3: Razmere pri ulivanju razli~nih vzorcev (za vsako stanje so
bili uliti 3 vzorci v pe{~eno formo pri temperaturi ulivanja 625 °C)
Mold coating type Sample
Without coating 1
Graphite coating 2
Micro-ceramic coating (ZR1) 3
Nano-ceramic coating (MM12) 4
3 RESULTS AND DISCUSSION
The beryllium element protects the melt from oxida-
tion and burning of the active element such as magne-
sium.2 The strontium element modifies the eutectic
silicon phase morphology and changes the needle-like to
a fine fibrous morphology, consequently the mechanical
properties and, especially, the elongation was improved.
The fine fibrous eutectic modifies the hypoeutectic Al–Si
alloys that occur when elements such as strontium and
sodium are added, and this has been explained based on
observations of increased twinning. The increased den-
sity of twinning is believed to result from an impurity-
induced twinning (IIT) mechanism that promotes further
growth by encouraging the formation of a perpetuating
twin plane re-entrant edge (TPRE).16–18
The chemical compositions of the nano-ceramic coat-
ing (MM12) and micro-ceramic coating (ZR1) are shown
in Table 4. The XRD patterns of the nano-ceramic coat-
ing (MM12) and the micro-ceramic coating (ZR1) pow-
ders are shown in Figures 3 and 4.
Table 4: Chemical composition of MM12 and ZR1 coatings used in
this study
Tabela 4: Kemijska sestava premazov MM12 in ZR1, uporabljenih v
tej {tudiji
Composition Al2O3 ZrO2 SiO2 H2O Others
Content (w/%) 30 7 1 60 2
With reference to the XRD analysis shown in Figure
3, the nano-ceramic coating used in this research particu-
larly contains Al2O3, ZrO2 and SiO2. Based on the broad-
ening of the most prominent peak in the XRD profile, a
M. BOROUNI et al.: EFFECT OF A NANO-CERAMIC MOLD COATING ON THE FLUIDITY LENGTH ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 473–477 475
Figure 4: XRD pattern of micro-ceramic coating (ZR1) used in this
study
Slika 4: XRD-posnetek mikrokerami~nega premaza (ZR1), uporablje-
nega v tej {tudiji
Figure 3: XRD pattern of nano-ceramic coating (MM12) used in this
study
Slika 3: XRD-posnetek nanokerami~nega premaza (MM12), uporab-
ljenega v tej {tudiji
Figure 2: Pouring basin shape15
Slika 2: Oblika livne ~a{e15
mean crystallite dimension D was calculated using
Scherer’s formula:19
D =
⋅
0 9.
cos

	 

(1)
where  stands for the X-ray radiation wavelength (here
0.15406 nm), 	 stands for the line broadening at half the
maximum intensity in radians. 	 calculated by equation
(2):20
	 	 	2 2 2= −M S (2)
	M = 	 in nano-ceramic coating pattern XRD (MM12)
	S = 	 in micro-ceramic coating pattern XRD (ZR1).
Here, 	Al2O3 = 0.0053, 	ZrO2 = 0.0012, 	SiO2 = 0.0058,
K is a shape factor, for the case of a sphere K = 0.9, and

 is the Bragg angle.19 
Al2O3 = 28.78°, 
ZrO2 = 16.91° and

SiO2 = 13.38°, therefore, DAl2O3 = 30 nm, DZrO2 = 120 nm
and DSiO2 = 25 nm.
The melting temperature of bulk Al2O3 is 2015 °C.
The melting temperature of the Al2O3 nano-particles (25
nm) should be similar to the bulk counterparts, because
the size effect becomes significant only when the particle
size is less than 10 nm.21 Therefore, no sintering of the
nano-coating occurs.
Figure 5 shows a SEM micrograph of the micro-
ceramic coating ZR1 and the nano-ceramic coating
MM12. The phase particles of the ZR1 coating have an
average size of approximately 2–20 μm. The ZR1 and
MM12 coatings have nearly similar compositions, but
different particle sizes. The micro-ceramic coating with a
similar composition to the nano-ceramic coating was
used to study the influence of the particle size coating on
the surface quality of the cast parts and the fluidity
length.
Figure 6 shows the effect of the cross-section and the
mold coating type on the fluidity length of the melt. With
an increasing cross-section of the channel, the fluidity of
melt was improved. The fluidity in the mold with the
nano-ceramic mold coating was a maximum and the
fluidity in the mold with the micro-ceramic coating and
graphite coating was better than in sand mold without
coating.
The surface roughness was measured using a Mahr-
Perthometer M2 roughness tester device. The nano-mold
coating had a very smooth surface, hence it reduced the
friction forces and the fluidity was improved. Figure 7
shows the relationship between the mold coating and the
surface roughness. This figure show that the nano-cera-
mic mold coating (MM12) has a minimum roughness,
and hence the fluidity length in the mold with this
M. BOROUNI et al.: EFFECT OF A NANO-CERAMIC MOLD COATING ON THE FLUIDITY LENGTH ...
476 Materiali in tehnologije / Materials and technology 48 (2014) 4, 473–477
Figure 7: Surface roughness versus mold coating
Slika 7: Hrapavost povr{ine pri razli~nih premazih forme
Figure 5: SEM micrographs of: a) micro-ceramic coating ZR1 and b)
nano-ceramic coating MM12 used in this study
Slika 5: SEM-posnetka: a) mikrokerami~nega premaza ZR1 in b)
nanokerami~nega premaza MM12, uporabljena v tej {tudiji
Figure 6: Fluidity length versus cross-section channel
Slika 6: Teko~nost v odvisnosti od prereza kanala
coating is better than the mold with micro-coating,
graphite coating and a mold without coating.
Research has shown that the wetting angle between
the Al melt and ZrO2 and Al2O3 is 70° and 152°, respec-
tively.12 Thus, the wetting angles are large, and conse-
quently the wettability of the mold when using molten
aluminum is very low and the fluidity in the thin section
can be improved.12 The completely proper characteristics
of the mold with nano-ceramic coatings can be for the
following reasons:12,22,23
1. Thermal and chemical stability, because of ionic and
covalent bonds of the ceramics. After the experiment,
the composition analysis performed on the inner
surface of the coating mold and in the alloy samples
achieved, indicated that no chemical reaction oc-
curred between the alloy melt and the mold materials,
and therefore no gas is produced and this increases
the fluidity.
2. Best flow of melt into mold, because this reduces the
coefficient friction and the surface roughness of the
mold by smooth and clean surfaces with the nano-
mold coating, consequently the melt easily fills the
mold and improves the fluidity.
3. Large wetting angle between the aluminum melt and
the particles of the coating, which reduces the wetta-
bility of the mold coating with the aluminum melt,
and consequently it easily fills the mold.
4. Crystal lattice misfit between the melt and the coat-
ing materials that inhibits the nucleation of the
aluminum melt on a nano-mold coating and conse-
quently there is easy melt flow and an enhancement
of the fluidity.
4 CONCLUSIONS
1. By applying a graphite mold coating, a micro-cera-
mic mold coating and a nano-ceramic mold coating
on a sand mold surface, because the coat operates the
insulation role and reduces the roughness of the mold
surface, this increases the fluidity and castability of
melt.
2. The melt fluidity and the castability in the sand mold
with a nano-ceramic coating has the best result in
comparison with the sand mold with a micro-ceramic
mold coating, a graphite mold coating and a sand
mold without coating.
3. By increasing the cross-section of the channel,
because it reduces the cooling rate and the surface
tension, this improves the fluidity length and the
castability.
4. Mold coating significantly improves the fluidity, and
hence a lower casting temperature can be used. The-
refore, casting defects were reduced and indirectly
mechanical properties could be improved.
Acknowledgements
The authors are grateful for support of this research
by Isfahan University of Technology (IUT).
5 REFERENCES
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2 V. S. Zolotorevsky, N. A. Belov, Casting Aluminum Alloy, 1st ed.,
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eutectic modification and T4 heat treatment on mechanical properties
and corrosion resistance of an Al-9wt%Si casting alloy, Materials
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294–299
6 K. R. Ravi, R. M. Pillai, K. R. Amaranathan, B. C. Pai, M. Chakra-
borty, Fluidity of aluminum alloys and composites: A review, Journal
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7 Casting Design Handbook, ASM, 1962
8 M. Di Sabatino, Fluidity of aluminum foundry alloys, Thesis sub-
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9 Y. D. Kwon, Z. H. Lee, The effect of grain refining and oxide inclu-
sion on the fluidity of Al-4.5Cu-0.6Mn and A356 alloys, Materials
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10 P. R. Beeley, Foundry Technology, Butterworth Scientific, 1972
11 A. N. Goldstein (ed.), Handbook of Nanophase Materials, Marcel
Dekker Inc., 1997
12 L. Feng, Y. Gencang, Solidification of superalloy in a SiO2- ZrO2-
B2O3 coating mould, Journal of Non-Crystalline Solids, 290 (2001)
2–3, 105–114
13 J. E. Gruzleski, B. M. Closset, The Treatment of Liquid Aluminium-
Silicon Alloys, American Foundrymen's Society, Inc., 1990
14 Metals Handbook Aluminium and Aluminium Alloys, 1992
15 Metalcaster’s Reference and guide, first edition, AFS, Augest 1972
16 S. Zulu, A. Hellawell, Solidification processing, (1987), 4–47
17 S. D. McDonald, K. Nogita, A. K. Dahle, Eutectic grain size and
strontium concentration in hypoeutectic aluminium-silicon alloys,
Journal of Alloys and Compounds, 422 (2006) 1–2, 184–191
18 K. Nogita, H. Yasuda, M. Yoshiya, S. D. McDonald, A. Takeuchi, Y.
Suzuki, The role of trace element segregation in the eutectic
modification of hypoeutectic Al-Si alloys, Journal of Alloys and
Compounds, 489 (2010) 2, 415–420
19 P. N. Prasad, Nanophotonics, John Willey & Sons Inc., 2004
20 B. D. Cullity, Elements of X-Ray Diffraction, 2nd ed, Addison-Wes-
ley, 1978
21 J. Y. Huang, In Situ Observation of Quasimelting of Diamond and
Reversible Graphite-Diamond phase Transformations, Nano Letters,
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(2000), 175–182
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M. BOROUNI et al.: EFFECT OF A NANO-CERAMIC MOLD COATING ON THE FLUIDITY LENGTH ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 473–477 477

A. HALAP et al.: INFLUENCE OF THE THERMO-MECHANICAL TREATMENT ON THE EXFOLIATION ...
INFLUENCE OF THE THERMO-MECHANICAL
TREATMENT ON THE EXFOLIATION AND PITTING
CORROSION OF AN AA5083-TYPE ALLOY
VPLIV TERMO-MEHANSKE OBDELAVE ZLITINE AA5083 NA
LU[^ENJE IN JAMI^ASTO KOROZIJO
Akram Halap1, Miljana Popovi}1, Tamara Radeti}1, Veselin Va{~i}2,
Endre Romhanji1
1Department of Metallurgical Engineering, Faculty of Technology and Metallurgy, University of Belgrade, Karnegijeva 4, 11120 Belgrade, Serbia
2Freelance Consultant
akram_halap@yahoo.com, endre@tmf.bg.ac.rs
Prejem rokopisa – received: 2013-07-18; sprejem za objavo – accepted for publication: 2013-10-04
Exfoliation and pitting corrosion were tested on an AA5083-type alloy sheet, after applying different thermo-mechanical
treatments. Hot-rolled plates were laboratory cold rolled with 16–54 % reductions, and subsequently annealed for 2 h in the
temperature range from 220 °C to 280 °C. The performed tests have shown that an unacceptable exfoliation and a corrosion
susceptibility were developed after the 54 % deformation and low-temperature annealing at 220 °C. However, no severe
exfoliation was observed after the cold rolling with low deformations, independent of the applied annealing temperature. The
pitting corrosion was less agressive, and the highest degree of the pitting attack was not experienced for the tested alloy. A slight
tendency to pitting was detected after the high cold deformation and high-temperature annealing for the samples with
recrystallized but significantly flattened grain structures.
Keywords: AA5083 alloy, thermo-mechanical treatment, exfoliation, pitting
Lu{~enje in jami~asta korozija sta bili preizku{eni na plo~evini iz zlitine AA5083 po razli~nih termo-mehanskih obdelavah.
Vro~e valjane plo{~e so bile v laboratoriju hladno valjane z redukcijo 16–54 %, nato `arjene 2 h v obmo~ju temperature od 220
°C do 280 °C. Preizkusi so pokazali nesprejemljivo lu{~enje in ob~utljivost za korozijo po 54-odstotni deformaciji in nizki
temperaturi `arjenja 220 °C. Mo~nej{ega lu{~enja ni bilo opaziti po hladnem valjanju z majhno deformacijo, neodvisno od
uporabljene temperature `arjenja. Jami~asta korozija je bila manj agresivna in mo~ne jami~aste korozije ni bilo pri preizku{eni
zlitini. Rahla tendenca po nastajanju jamic se je pokazala pri veliki hladni deformaciji in visoki temperaturi `arjenja pri vzorcih,
ki so bili rekristalizirani in so imeli mo~no splo{~eno strukturo.
Klju~ne besede: zlitina AA5083, termo-mehanska obdelava, lu{~enje, jamica
1 INTRODUCTION
Al-Mg alloys found a wide range of applications,
especially in the construction of transportation means
due to the attractive strength-to-weight ratio, good weld-
ability and formability as well as high corrosion resi-
stance.1–4 However, the Al-Mg alloys with more than 3 %
Mg, due to a limited room-temperature solubility of Mg
in an Al matrix (w(Mg) = 1.9 %), can be sensitized due
to the formation of a continuous film of the 	-phase
(Mg2Al3) at the grain boundaries, and become suscep-
tible to intergranular corrosion (IGC).5–9 Generally, the
susceptibility to the IGC is influenced, in a complex
manner, by the structure features of a material,10,11 and an
optimization of the overall thermo-mechanical treatment
is the key point in producing a corrosion-resistant
material.10 One of the manifestations of the IGC is a
specific lamellar type of corrosion, with a blister-surface
appearance, known as exfoliation. It is a form of the IGC
that occurs on the surfaces of the wrought aluminum
alloys12–14 with highly flattened grain structures. Another
common corrosion type, also associated with a specific
surface appearance, is pitting corrosion. It is a form of an
extremely localized corrosion15 that leads to a creation of
small holes or pits in the structure, usually covered by
the corrosion products. These types of corrosion were
often observed and analyzed on 2xxx and 7xxx type
alloys, used in aircraft industries, and due to their im-
portance a major attention was paid to the investigation
of these alloys.16,17
The objective of this work was to study the exfoli-
ation and pitting corrosion on a moderate-strength non-
heat treatable AA5083-type-alloy sheet, after applying
different thermo-mechanical treatments (TMTs). Exfoli-
ation and pitting corrosion appeared to be very important
manifestations of a sea-water attack in Al-Mg-based
alloys,18 therefore, marine-grade alloys should be
subjected to the corrosion testing defined in ASTM B
928/B 928M-04a standard,19 including intergranular,
exfoliation and pitting corrosion assesments.
2 EXPERIMENTAL WORK
Material. The material used in this study was an
industrially manufactured hot-rolled thick plate 7.4 mm,
supplied by Impol-Seval Aluminium Rolling Mill. The
chemical composition is listed in Table 1.
Materiali in tehnologije / Materials and technology 48 (2014) 4, 479–483 479
UDK 669.715:620.19 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)479(2014)
The as-received hot-rolled plates were laboratory
cold rolled to 3.5 mm and inter-annealed for 3 h at 350
°C. These samples were further cold rolled with different
degrees of deformation ranging from 16 % to 54 %, and
then finally annealed for 2 h in the temperature range of
220–280 °C and air cooled.
Table 1: Chemical composition of the tested alloy in mass fractions
(w/%)
Tabela 1: Kemijska sestava preizkusne zlitine v masnih dele`ih (w/%)
Mg Mn Cu Fe Si Zn Cr Ti Sr
5.13 0.718 0.013 0.337 0.11 0.513 0.008 0.0254 0.003
Corrosion testing. Susceptibility to exfoliation and
pitting corrosion was determined with a visual inspection
using the ASSET method (the method for a visual
assessment of the exfoliation-corrosion susceptibility of
AA5xxx-series Al alloys), described in ASTM G66
standard.
Brinell hardness was measured on the samples pro-
cessed in the manner described above, as well as the
grain aspect ratio, calculated after revealing the grain
structure using Barker’s etchant and a microstructure
analysis with an optical microscope.
3 RESULTS
Exfoliation and pitting-corrosion susceptibility were
ranked after a visual inspection according to ASTM G66
standard. The following letters were used to describe the
degrees of pitting corrosion: N – no appreciable attack,
or A, B and C levels of attack (A and B denote accept-
able corrosion, C denotes susceptibility to corrosion). In
case of exfoliation (besides N) the ratings are as follows:
A, B, C and D levels of attack (A and B denote accept-
able exfoliation, C and D denote susceptibility to exfoli-
ation). A combination of two letters was used in the case
when the surface morphology could not be described
with only one letter due to a mixed- or boundary-type
surface appearance like the ones on the photographs
given in ASTM G66 standard.
The ratings of the exfoliation and pitting-corrosion-
attack dependence on the applied TMTs, i.e., the
cold-rolling reductions and annealing temperatures, are
shown in Figures 1a and 1b. The results for the general
IGC susceptibility of the tested alloy (expressed as a
mass loss per unit area, according to ASTM G67
standard), considered elsewhere,20 are also attached. The
results shown in Figure 1a indicate that the C level
denoting unacceptable exfoliation appeared after the 54
% cold deformation and low-temperature annealing at
220 °C. However, after a prior low deformation of 16 %,
the exfoliation was absent, even after a low-temperature
annealing at 220 °C, and not related to the IGC
susceptibility that was very pronounced in this case, as
shown in Figure 1a.
The most detrimental pitting-corrosion susceptibility,
described as the C-type surface appearance, was not
recognized for the tested alloy. However, after a higher
degree of the cold-rolling reduction, 54 %, and the final
annealing at 260 °C and 280 °C, the B level of the pitting
susceptibility was identified (Figure 1b). The same B
level of the pitting attack was noticed after the 30 %
deformation and the annealing at the highest temperature
of 280 °C.
A. HALAP et al.: INFLUENCE OF THE THERMO-MECHANICAL TREATMENT ON THE EXFOLIATION ...
480 Materiali in tehnologije / Materials and technology 48 (2014) 4, 479–483
Figure 1: Ratings of the: a) exfoliation and b) pitting corrosion levels (according to ASTM G66 standard) for the cold rolled (16–54 %) and
subsequently annealed (220–280 °C) specimens. The appropriate data for the IGC susceptibility is also shown. The values below the lower
dashed line (IGC = 15 mg/cm2) indicate the IGC resistance, while the ones above the upper line (IGC = 25 mg/cm2) indicate the IGC sus-
ceptibility.
Slika 1: Ocena: a) lu{~enja in b) jami~aste korozije (skladno z ASTM G66-standardom) za hladno valjane (16–54 %) in kasneje `arjene
(220–280 °C) vzorce. Prikazani so tudi ustrezni podatki za ob~utljivost IGC. Vrednosti pod spodnjo ~rtkano ~rto (IGC = 15 mg/cm2) izkazujejo
odpornost proti IGC, medtem ko vrednosti nad zgornjo ~rto (IGC = 25 mg/cm2) izkazujejo ob~utljivost za IGC.
Illustrations of the surface appearances for the sam-
ples subjected to the ASSET test, used for the exfoliation
and pitting assesment, are shown on Figures 2 and 3.
Figure 2 shows the surfaces of the specimens cold rolled
with a low-degree deformation, showing no corrosion
and no influences of the final annealing temperature.
However, a severe C-type exfoliation was detected on
the surfaces of highly deformed samples (54 %) subse-
quently annealed at a low temperature (220 °C), as
shown in Figure 3a. The B-level pitting corrosion was
observed after the 54 % deformation and annealing at
260 °C or 280 °C as shown in Figures 3c and 3d.
The results of the hardness measurements shown in
Figure 4 revealed that an increase in the cold-rolling
reduction up to  35 % caused a hardness increase. In
the cases of higher cold-rolling reductions of > 35 %, the
hardness level stayed almost unchanged (ranging from
85–100 HB), except for the samples annealed at 280 °C.
For those samples, a steep decrease in the hardness was
observed after a 20 % pre-deformation. However, after
the annealing at 260 °C a less steep but continuous drop
in the hardness occurred when the prior cold deformation
was increased to over 30 %.
In order to correlate the corrosion behavior and the
microstructure of the tested alloy, the grain structures of
the samples cold rolled with 16 % and 54 % reductions
and annealed at 220 °C 2 h and 280 °C 2 h, were
revealed as shown in Figures 5a to 5d. In the case of the
prior 54 % cold deformation, the grains were elongated
after the annealing at both temperatures, independently
of the recrystallization processes that occurred after the
annealing at 280 °C 2 h (the recrystallization is indicated
by the hardness drop in Figure 4).
The grain aspect ratio was calculated for the prior
cold deformations of (16, 30 and 54) %, for the whole
range of the annealing temperatures, as shown in Figure
6. It was found that the appropriate grain aspect ratios
were around 5.5 for the sample annealed at 220 °C 2 h
and 5.7 for the sample annealed at 280 °C 2 h after the
54 % deformation.
A. HALAP et al.: INFLUENCE OF THE THERMO-MECHANICAL TREATMENT ON THE EXFOLIATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 479–483 481
Figure 3: Surface appearance for the samples deformed with the 54 %
reduction and annealed for 2 h at different temperatures
Slika 3: Videz povr{ine vzorcev deformiranih s 54-odstotno redukcijo
in `arjenih 2 h pri razli~nih temperaturah
Figure 4: Influence of cold-rolling reductions on the hardness of the
samples subsequently annealed for 2 h in the temperature range of
220–280 °C
Slika 4: Vpliv redukcije pri hladnem valjanju na trdoto vzorcev, ka-
sneje `arjenih 2 h v obmo~ju 220–280 °C
Figure 2: Surface appearance for the samples deformed with the 16 %
reduction and annealed for 2 h at different temperatures
Slika 2: Videz povr{ine vzorcev, deformiranih s 16-odstotno redukcijo
in `arjenih 2 h pri razli~nih temperaturah
After the 16 % deformation the grains were less elon-
gated, with the grain aspect ratio being 1.7 for the sam-
ples annealed at 220 °C 2 h and 3.5 for the samples
annealed at 280 °C 2 h . After the 30 % pre-deformation
the grain aspect ratio was from 2.3 to 4.4 for the tested
range of temperatures. Generally, the grain aspect ratio
significantly increased with the increase in the annealing
temperature.
4 DISCUSSION
Previous studies12–14,21 have shown that the grain
structure developed under different thermo-mechanical
treatments (TMTs) can have an important role in deve-
loping the exfoliation corrosion of the high-strength
alloys such as 2xxx or 7xxx series of Al alloys. The
basic influence of TMTs is related to the grain flatness or
grain aspect ratio enabling a specific long surface path
for the IGC, and a formation of corrosion products on a
few surface planes with a higher specific volume. In this
way the expansion tendency creating the compressive
stress can lift the surface grains and form surface bli-
sters. On the other hand, it is already known that the
general IGC susceptibility is mostly controlled with the
alloy composition and the distribution of grain-boundary
precipitates.22
In this work, the harmful C level of exfoliation was
experienced only after the 54 % cold-rolling deformation
and the subsequent low-temperature annealing at 220 °C
(Figures 1a and 3a). In this case the recovery degree was
supposed to be low, as it was indicated by the hardness
level after the annealing at 220 °C (Figure 4). The mi-
crostructural observations of these samples showed that
the grains were highly elongated (Figure 5c) with the
grain aspect ratio of  5.7 (Figure 6). For the samples
with the lower grain aspect ratio of  2.3 after the 30 %
deformation, or  1.7 after the 16 % deformation, no
unacceptable exfoliation level of attack was observed or
it was completely absent (Figure 1a). This indicates the
grain-flatness importance in the case of the exfoliation-
type corrosion. It should also be emphasized that for the
16 % deformed samples, the general IGC susceptibility
was at the highest level (Figure 1a), but obviously the
grain flatness was low and exfoliation could not develop
at all. The grain aspect ratio seemed to increase with the
increasing annealing temperature (Figure 6), but in spite
of this, no detrimental level of the exfoliation attack took
place (Figure 1a). It was assumed that the general IGC
susceptibility was reduced to such an extent that it sup-
pressed exfoliation, i.e., any type of intercrystalline
corrosion. Therefore, after the low-temperature anneal-
ing, the 54 % deformed sample with a non-recrystallized
structure and the grain aspect ratio of 5.5 was prone to
exfoliation, while the recrystallized structure with the
grain aspect ratio of 5.7 (after the annealing at 280 °C)
was completely resistant to exfoliation, as the IGC was
suppressed too. It was supposed that an improvement of
the corrosion properties with the annealing-temperature
variations can be basically related to the 	-phase preci-
pitation, i.e., the change of a continuous film after the
low-temperature annealing to a discontinuous discrete
	-phase morphology after the high-temperature anneal-
ing at 260 °C and 280 °C, as discussed elsewhere.20
Therefore, in practice, the exfoliation-corrosion control
should be considered in a limited range of the structure
states that are also fully resistant to the IGC. However,
according to our results the tendency to the exfoliation-
type corrosion in all the circumstances was related to a
flattened-grain structure, i.e, it was suppressed in the
case of the structure with a grain aspect ratio below a
certain critical value.
On the other hand, in the case of a lower grain aspect
ratio or a structure with less elongated grains, the pitt-
A. HALAP et al.: INFLUENCE OF THE THERMO-MECHANICAL TREATMENT ON THE EXFOLIATION ...
482 Materiali in tehnologije / Materials and technology 48 (2014) 4, 479–483
Figure 6: Grain aspect ratios for the samples cold rolled with (16, 30
and 54) % reductions and subsequently annealed for 2 h at 220 °C and
280 °C
Slika 6: Razmerje dol`ine proti vi{ini zrn pri hladno valjanih vzorcih
z redukcijo (16, 30 in 54) % ter kasneje `arjenih 2 h pri temperaturah
120 °C in 280 °C
Figure 5: Grain structures of the samples deformed with the: a), b)
16 % and c), d) 54 % cold-rolling reductions and annealed for 2 h at:
a), c) 220 °C and b), d) 280 °C
Slika 5: Zrnata struktura vzorcev po hladnem valjanju, deformiranih
s: a), b) 16 % in c), d) 54 % redukcijo ter `arjenih 2 h pri: a), c) 220
°C in b), d) 280 °C
ing-type corrosion is thought to be the active process of
the intercrystalline corrosion.12,13 In this experiment, no
unacceptable C-level pitting was detected (Figures 1b,
3c and 3d) according to the fact that under the applied
cold deformations and heat treatments no equiaxed grain
structure was developed (Figure 5). The appearance of a
severe but not harmful B-level pitting corrosion,
observed after the annealing at 260 °C and 280 °C of the
54 % deformed samples (Figures 1b and 3b), could be
related to a high degree of the recovery or recrystal-
lization process. This is in a good agreement with the
hardness drop (Figure 4) and cannot be related to the
development of a less flattened or equiaxed grain struc-
ture enabling the pitting corrosion with the progress in
the intercrystalline corrosion over the shorter paths
around the grains.12,13
Using the results shown in Figure 1b it was difficult
to correlate the degree of the pitting-corrosion attack and
the general intergranular-corrosion (IGC) susceptibility
as the pitting was even stronger for the samples showing
a high IGC resistance (indicated by a very low mass loss
in Figure 1b).
5 CONCLUSION
The exfoliation and pitting corrosion of an AA5083-
type alloy sheet were considered after applying different
thermo-mechanical treatments (TMTs): cold-rolling
deformations from 16 % to 54 % and the final annealing
for 2 h at 220–280 °C.
The results of the corrosion testing, performed
according to ASTM G66 standard, showed that the
exfoliation-type corrosion can be significantly affected
by varying the applied TMTs. A detrimental or unaccept-
able exfoliation-corrosion level was experienced after
high cold-rolling deformations ( 50 %) and low-tempe-
rature annealing (at 220 °C) together with a highly
elongated grain structure. However, in the case of the
high-temperature annealing and a certain grain flatness,
the exfoliation could not develop as the general IGC was
suppressed too.
No harmful level of pitting corrosion was detected as
the equiaxed grain structure was not developed after
applying different TMTs. However, a moderate pitting
attack was noticed after the annealing at high tempera-
tures on the samples with recrystallized but highly
flattened grain structures and with a very low IGC
susceptibility.
Acknowledgements
The authors are grateful to the Ministry of Education,
Science and Technological Development, Republic of
Serbia, and Impol-Seval Aluminium Mill, Sevojno, for
supporting this research under contract TR 34018.
6 REFERENCES
1 Transport and Aluminium, The Aluminium Association-IAI, 2008
2 G. M. Raynaud, Ph. Gomiero, Aluminium Alloys for the Marine
Market, Aluminium and its Alloys, 79 (1996), 73
3 K. Baumann, M. Bertram, P. Kistler, In: M. Pekguleryuz (ed.), Conf.
Proc.: Light Metals in Transport Applications, Canadian Institute of
Mining, Metallurgy and Petroleum, 2007, 325–338
4 T. Summe, The Aluminum Advantage – Commercial Vehicle Appli-
cations, SAE, Automotive Aluminum – The Performance Advantage,
from www.autoaluminum.org
5 M. Popovi}, E. Romhanji, Mater. Sci. Eng., A492 (2008), 460–467
6 M. Popovi}, E. Romhanji, B. Minov, D. Gli{i}, SCC susceptibility
and formability in relation to different TMTs of an Al-6.8 wt% Mg
alloy sheet, In: J. Hirsch, B. Skrotski, G. Gottstein (eds.), Conf. Proc.
ICAA11: Aluminium Alloys – Their Physical and Mechanical
Properties, Wiley-VCH Verlag GmbH & Co. KgaA, Weinheim, 2
(2008), 2155–2162
7 A. L. Davenport et al., Mater. Sci. Forum, 519–521 (2006), 641–646
8 J. C. Chang, T. H. Chuang, Metall. Mater. Trans., 30A (1999),
3191–3199
9 J. L. Searles, P. I. Gouma, R. G. Buchheit, Mater. Sci. Forum,
396–402 (2002), 1437–1442
10 L. Tan, T. R. Allen, Corrosion Science, 52 (2010), 548–554
11 M. Popovi}, T. Radeti}, E. Romhanji, Precipitation of the 	-phase
and Corrosion Behavior of an Al-6.8 wt.% Mg Alloy, Proc. of 13th
International Conference on Aluminum Alloys ICAA13, H. Weiland,
A. D. Rollett, W. A. Cassada (eds.), John Wiley & Sons, Inc.,
Hoboken, NJ, USA 2012, 363–370
12 M. J. Robinson, Corrosion Science, 22 (1982), 775–790
13 M. J. Robinson, N. C. Jackson, Corrosion Science, 41 (1999),
1013–1028
14 M. J. Robinson, N. C. Jackson, British Corrosion Journal, 34 (1999),
45–49
15 D. W. Hoeppner, C. A. Arriscorreta, International Journal of Aero-
space Engineering, Article ID 191879, (2012), 1–29
16 D. A. Jones, Principles and Prevention of Corrosion, 2nd ed., Pren-
tice-Hall, Inc., Upper Saddle River, NJ 1996
17 B. W. Lifka, Corrosion of aluminum alloys, Corrosion Engineering
Handbook, P. A. Schweitzer, Marcel Dekker, New York 1996,
129–133
18 H. Bushfield, M. Cruder, R. Farley, J. Towers, Marine Aluminum
Plate – ASTM Standard Specification B 928 and the Events Leading
to its Adoption, 2003, available from: www.boatdesign.net/...what-
best-marine-alluminum-alloy-5083-alloy
19 ASTM standard B 928/B 928M-04a: Standard Specification for High
Magnesium Aluminum-Alloy Sheet and Plate for Marine Service and
Similar Environments
20 T. Radeti}, A. Halap, M. Popovi}, E. Romhanji, Effect of the
thermo-mechanical treatment on IGC susceptibility of AA 5083
alloy, In: J. Granfield (eds.), Light Metals 2014: Proceedings of TMS
2014 Annual Meeting, J. Wiley & Sons, Hoboken, NJ, USA 2014,
297–302
21 D. McNaughtan, M. Worsfold, M. J. Robinson, Corrosion Sci., 45
(2003), 2377–2389
22 J. R. Davis, Corrosion of Aluminum and Aluminum Alloys, ASM
International, Metals Park, Ohio 1999, 101
A. HALAP et al.: INFLUENCE OF THE THERMO-MECHANICAL TREATMENT ON THE EXFOLIATION ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 479–483 483

A. SAXENA et al.: STRUCTURAL AND ELECTRICAL STUDIES OF FULLERENE (C60) ...
STRUCTURAL AND ELECTRICAL STUDIES OF
FULLERENE (C60) DISPERSED POLYMER
ELECTROLYTES
[TUDIJ STRUKTURNIH IN ELEKTRI^NIH LASTNOSTI
POLIMERNEGA ELEKTROLITA Z DISPERGIRANIM FULERENOM
(C60)
Amit Saxena1,2, Pramod Kumar Singh2, Bhaskar Bhattacharya2
1Department of Physics, Hindustan College of Science and Technology, Farah, Mathura- 281122, India
2Material Research Laboratory, School of Engineering and Technology, Sharda University, G. Noida 201310, India
saxena.electronics@gmail.com, b.bhattacharya@sharda.ac.in
Prejem rokopisa – received: 2013-07-18; sprejem za objavo – accepted for publication: 2013-10-29
In this work we report on a modification of the optical and electrical properties of a polymer electrolyte film (PEO : NaI) by
doping with C60. Light microscopy was used to study the surface morphology of the polymer electrolyte film doped with C60.
For the electrical properties, complex impedance spectroscopy (CIP) was used and the conductivity was calculated. The
dielectric, modules and transference numbers were calculated and explained in detail to support the electrical conductivity data.
Keywords: polymer electrolyte, light microscopy (LM), conductivity, dielectric constant, ion transference number
V tem delu poro~amo o spremembi opti~nih in elektri~nih lastnosti polimerne elektrolitske plasti (PEO : NaI) po dopiranju s
C60. [tudij morfologije povr{ine elektrolitske polimerne plasti je bil izvr{en s svetlobno mikroskopijo (LM). Za elektri~ne
lastnosti je bila uporabljena kompleksna impedan~na spektroskopija (CIP), prevodnost pa je bila izra~unana. Dielektri~nost,
moduli in transportno {tevilo so bili izra~unani in podrobno razlo`eni za podporo podatkov elektri~ne prevodnosti.
Klju~ne besede: polimerni elektrolit, svetlobna mikroskopija (LM), prevodnost, dielektri~na konstanta, transportno {tevilo
1 INTRODUCTION
Ion-conducting polymers, obtained by polymer-salt
complexation, are becoming increasingly important
because of their application in solid-state polymer batte-
ries, fuel cells, etc. Many good ion-conducting polymer
electrolytes have been developed by complexing poly-
mers (for example, PEO) with salts like lithium perchlo-
rate, lithium triflate, sodium chloride, sodium iodide,
etc.1–4 To enhance the conductivity of the polymer elec-
trolyte, various approaches have been adopted, like (a)
changing the type of complexing polymer (b) the use of
polymers with different chain lengths (c) the use of
different complexing salts and (d) modifying the degree
of crystallinity by the use of plasticizers and copolyme-
rization.5–8 However, the doping of fullerene (C60) with
a polymer electrolyte is a relatively a new approach to
enhance the optical and electrical properties of polyme-
ric films. In this paper we report an important observa-
tion regarding a change in the physical and electrical
properties of polymer films doped with fullerene
(C60).9–11
2 EXPERIMENTAL
2.1 Material Preparation
Composite polymer films with a thickness varying
between 200 μm to 300 μm were synthesis using the
standard "solution cast" technique.12
Before the preparation of the polymer film, the
fullerene is prepared in the laboratory. To obtain this
fullerene, soot is collected by the burning of phenol-
phthalein, which is dissolved in benzene and stirred
continuously for approximately 5 h. Finally, the solution
was filtered. A wine-coloured solution was obtained,
which was characterized by NMR to confirm the forma-
tion of C60. The NMR results showed the formation of
fullerenes (methanofullerene) (Figure 1) with a small
chain attached with the ball. The obtained structure was
similar to that required for doping in polymer electrolyte
films. Finally, PEO (Mw: 2 000 000) and NaI were
weighed separately in the required amounts. These
weighed components were then dissolved in distilled
methanol. Different weights of prepared fullerene were
added to this combination. After this the solution was
stirred for a long time to achieve uniform mixing and
complexation. Then, the mixture was poured into poly-
propylene/Teflon moulds for casting in ambient, room
conditions. The solvent was then slowly evaporated at
room temperature. Finally, the film was dried under
vacuum to eliminate any traces of solvents. The as-pre-
pared films were subjected to different characterizations.
The polymer (PEO) and salt (NaI) molar ratio was
maintained at 0.065 13 and the amount of C60 dispersed
in the complex matrix was expressed in mass fractions.
The composition is represented as (PEO : NaI) +
w(C60).
Materiali in tehnologije / Materials and technology 48 (2014) 4, 485–490 485
UDK 678.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)485(2014)
2.2 Material Characterisation
The surface morphology of the composite polymer
film was studied with an optical microscope. In order to
evaluate the ionic conductivity of the polymer electrolyte
film, impedance spectroscopic techniques were used.
The conductivity of the polymer films was evaluated
from the bulk resistance calculated by ac complex impe-
dance spectroscopy over a frequency range of 100 Hz to
1 MHz. The transference numbers (tion) of the C60-
doped polymer electrolyte composite sample were
measured using Wagner’s polarization method.4 In this
method the sample is placed between two electrodes in
such a way that it minimizes the contact resistance. In an
ideal case one of the electrodes is blocking, whereas the
other is non-blocking. In the present case we used
stainless-steel electrodes. A DC potential is applied
across the sample in such a way that mobile ion species
move towards the non-blocking electrode and a small
polarization current flows and, finally, the sample
becomes polarized. The initial total current Ii and the
final residual current If are use to evaluate the ionic
transference number.7
3 RESULTS AND DISCUSSION
3.1 NMR
The samples were investigated in the NMR study.
The result confirms the formation of methanofullerene.
Figure 1 shows the structure of methanofullerene C61
(CO2Me) [C6H6].9
3.2 Surface Morphological Study
3.2.1 Morphology
Typical light micrographs of the composite polymer
electrolyte system (PEO : NaI) + w(C60) are shown in
Figure 2.
A. SAXENA et al.: STRUCTURAL AND ELECTRICAL STUDIES OF FULLERENE (C60) ...
486 Materiali in tehnologije / Materials and technology 48 (2014) 4, 485–490
Figure 3: Complex impedance spectra of (PEO : NaI) + w(C60) film at room temperature: a) 0 %, b) 6 %, c) 9 %, d) 20 %
Slika 3: Kompleksni spekter impedance za plast (PEO : NaI) + w(C60) pri sobni temperaturi: a) 0 %, b) 6 %, c) 9 %, d) 20 %
Figure 1: Chemical structure of methanofullerene C61
(CO2Me)[C6H6]
Slika 1: Kemijska struktura metanofulerena C61 (CO2Me)[C6H6]
Figure 2: Light micrographs of (PEO : NaI) + w(C60) composite
polymer electrolyte system: a) 0 % and b) 6 %
Slika 2: Mikrostruktura polimernega elektrolitskega kompozitnega
sistema (PEO : NaI) + w(C60): a) 0 % in b) 6 %
(a)
(b)
The spherulites are observed (Figure 2a) in the pure
polymer complex (i.e., for w(C60) = 0). The dark
boundaries observed between the spherulites show the
partial amorphous phase of the films. The morphological
pattern of the film changed upon the addition of different
amounts of the dispersoid. This effect is well shown in
Figure 2b. Several dark spots are visible in the case of
doping with w = 6 % of C60. This indicates that a larger
addition of dispersoid will make more amorphous phase
of the films.13
3.3 Electrical Properties Study
3.3.1 Complex Impedance Spectroscopy (CIS)
A CIS analysis of the C60 dispersed polymer
electrolyte films was performed to determine the change
in the electrical properties. The impedance patterns for
the different compositions of the films are shown in
Figure 3.
It is clearly shown in the figure that as the concen-
tration of the dispersoid increases the impedance pattern
of the sample is changed. This impedance graphs are
used to calculate the bulk resistance (Rb) of the sample.10
It is clear from Figure 3 that the value of the bulk resi-
stance is decreased as the concentration of C60 is
increases. Now by using the values of bulk resistance,
the area and the thickness of the film we are able to
calculate the conductivity of the sample.
3.3.2 DC conductivity
The electrical conductivity of the composite polymer
films as a function of doping concentration was cal-
culated. Figure 4 shows the change in the conductivity
in the polymer electrolyte film doped with C60.
Figure 4 shows that the conductivity decreases with
the doping of the dispersoid C60. The minimum of the
conductivity was seen at 9 % of doping, but the value
suddenly increases to a maximum of 7.9 × 10–7 S/cm
when doping with w(C60) = 20 %
3.3.3 Dielectric Constant
The frequency-dependent dielectric constant () is
calculated for different mass fractions of dispersoid and
shown in Figure 5. It is clear that the dielectric decreases
as the doping of C60 begins, but at 6 % of doping the
dielectric increases.
Two increases in the dielectric constant are clearly
seen in Figure 5, at 6 % and 20 %. This result for the
dielectric supports the change in the conductivity.10,14
3.4 Dielectric relaxation study
A dielectric formulism is useful in retrieving many
particulars about the conductivity behaviour of any
material. The dielectric constant is a measure of the
stored charge. The frequency dependence of the real (’)
and imaginary (”) parts of the dielectric constants are
shown in Figure 6.
’ and ” are calculated by using the following
relations:


’
”
=
Z
CZ 2
and 

”
’
=
Z
CZ 2
where Z” is the imaginary part of the impedance,  is
the angular frequency and C is the capacitance of the
empty measuring cell for the electrode area A and the
sample thickness L. The capacitance is calculated using
the relation:
C
A
L
=

At lower frequency the high value is due to the
polarization of the electrode.15 But at high frequency, the
periodic reversal of the electric field is so fast that the
ion diffusion in the electric field is not feasible. Hence,
the charge accumulation at the electrode decreases,
leading to a decrease in ’ and ”.16
3.4.1 Modulus Formalism
The sudden change in ’ and ” values at lower
frequency is possibly due to the contribution of the elec-
A. SAXENA et al.: STRUCTURAL AND ELECTRICAL STUDIES OF FULLERENE (C60) ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 485–490 487
Figure 5: Variation in dielectric of (PEO : NaI) + w(C60)
Slika 5: Spreminjanje dielektri~nosti (PEO : NaI) + w(C60)
Figure 4: Variation in conductivity of (PEO : NaI) + w(C60) at room
temperature
Slika 4: Spreminjanje prevodnosti (PEO : NaI) + w(C60) pri sobni
temperaturi
trode polarization effect. Hence, ionic conduction relaxa-
tion is not properly analysed in the dielectric formulism.
Here, the modulus is calculated to overcome this prob-
lem and widely used to analyse the relaxation phenome-
non for different compositions.
The values of M’ and M” were calculated using the
following relations:
M’
’
’ ”
=
+

 2 2
and M ”
”
’ ”
=
+

 2 2
where M’ and M” are the real and imaginary parts of the
electrical modulus.
Figure 7 shows the modulus as a function of lg  for
different w(C60). The peaks at a very high frequency are
invariably present in all the complexes and the height of
the peaks changes with the dispersoid (C60) concen-
tration. This indicates that a relaxation phenomenon is
occurring.17,18
The M’ shows a increasing trend with frequency
whereas the peaks are observed in the M” spectra in the
higher frequency region. The peak in the modulus forma-
lism confirms the ionic conduction in the system.19,20
3.4.2 Transference Number (tion)
This parameter will show the conductivity due to ions
in polymer electrolyte films. The tion of the prepared
films are calculated by dc polarization method. In this
method, the dc current is observed as a function of time
on the application of a fixed dc voltage (0.25 V) across
the sample electrodes. The tion was calculated using the
relation:
T
I I
Iion
initial final
inital
=
−
A. SAXENA et al.: STRUCTURAL AND ELECTRICAL STUDIES OF FULLERENE (C60) ...
488 Materiali in tehnologije / Materials and technology 48 (2014) 4, 485–490
Figure 8: Variation in tion with w(C60)
Slika 8: Spreminjanje transportnega {tevila z w(C60)
Figure 7: a) Variation in M’ with w(C60), b) variation in M” with
w(C60)
Slika 7: a) Spreminjanje M’ z w(C60), b) spreminjanje M” z w(C60)
Figure 6: a) Variation in real part of dielectric constant (’) with
w(C60), b) variation in imaginary part of dielectric constant (”) with
w(C60)
Slika 6: a) Spreminjanje realnega dela dielektri~ne konstante (’) z
w(C60), b) spreminjanje navideznega dela dielektri~ne konstante (”)
z w(C60)
Figure 8 shows a plot of tion vs different w(C60)
doped with polymer electrolyte. In this plot it is clear
that the ionic conductivity in the polymer electrolyte film
initially decreases, but the 6 % dispersed film shows a
high ionic conductivity. As the concentration of dis-
persed C60 is further increased the ionic conductivity of
the film decreases. It may be concluded from the plot
that the higher doping of C60 makes the polymer film
more electronically conductive than ionically conductive.
3.4.3 Concentration and Mobility
The conductivity of polymer electrolytes depends
upon the charge-carrier concentration (n), the charge on
the carrier (e), and the number of charge carrier. Hence,
when the charge concentration is changed by doping, the
conductivity is also expected to change. In polymer–salt
complexes, the charge-carrier concentration is not direct-
ly proportional to the doping concentration. To analyse
the origin of the conductivity changes in the present
system, the concentration factor (n) and the mobility fac-
tor (μ) were calculated with respect to the added w(C60).
Figure 9 shows these changes graphically.
The mobility variation data (Figure 9a) are approxi-
mately following the trend of the conductivity. This con-
firms that ion-pair formation and re-dissociation theory
applicable to polymer–salt complexes10 are also applica-
ble in the present system. The decrease in the mobility
may be due to electrostatic hindrance in the movement
and because of the increased weight of the charge
carriers.21–23
Figure 9b shows the variation in the concentration of
the charge carrier on the doping of different w(C60). It
initially attains a minimum (at 9 %) value, followed by a
maximum (at 20 %). Many theories explaining this phe-
nomenon can be found in literature. The change in the
conductivity can be correlated with a change in the
charge-carrier concentration (Figures 4 and 9b).
4 CONCLUSION
A fullerene-based polymer electrolyte was prepared
by a standard solution cast technique. The ionic conduc-
tivity  7.95 × 10–7 S/cm and the ionic transference
number,  0.87, indicate that the prepared film is mixed
conductor. Parameters such as the charge-concentration
factor (n) and the mobility factor (μ) were studied with
respect to the fullerene concentration. The relationship
between the charge concentration factor (n) and the
mobility factor (μ) with respect to the fullerene w(C60)
is established. At low frequency, the variation of the
dielectric constant with frequency suggests an electrode
interface polarization process. At a low frequency the
high value of the dielectric constant and loss is due to
electrode polarization.
Acknowledgement
This work was supported by the DST project
(SR/S2/CMP-0065/2010) of the government of India.
The Indian Institute of Goat Research Center, Farah,
Mathura is also acknowledged for support with taking
the LM images of the samples.
5 REFERENCES
1 J. R. MacCallum, C. A. Vincent (eds.), Polymer Electrolyte Reviews
1-2, Elsevier, Science Publishers Ltd., New York 1989, 1991
2 R. Saumya, K. Thakur Awalendra, R. N. P. Chaudhary, Ionics, 14
(2008), 255–262
3 D. E. Fenton, J. M. Parker, P. V. Wright, Polymer, 14 (1973), 589
4 H. Al-Ahmad, A. M. Zihlif, Journal of Thermoplastic Composite
Materials, 26 (2013), 263–275
5 T. J. Pinnavaia, G. W. Beall, Polymer-Clay Nanocomposites, John
Wiley and Sons, New York 2001, chap. 2
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630–638
7 A. Chandra, P. K. Singh, S. Chandra, Solid State Ionics, 15 (2002),
154–155
8 P. Kumar, P. K. Singh, B. Bhattacharya, Ionics, 17 (2011), 721–725
9 C. Wang, Zhi-Xin Guo, Fu. Shoukuan, Progress in Polymer Science,
29 (2004), 1079–1141
10 J. Subocoz, A. Valozhyn, M. Zenker, Rev. Adv. Mater. Sci., 14
(2007), 193–196
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Materiali in tehnologije / Materials and technology 48 (2014) 4, 485–490 489
Figure 9: Variation of: a) mobility (μ) and b) concentration (n) of
(PEO : NaI) + w(C60)
Slika 9: Spreminjanje: a) mobilnosti (μ) in b) koncentracije (n) (PEO :
NaI) + w(C60)
11 W. Kratschmer, L. D. Lamb, D. R. Huffman, Nature, 347 (1990),
354–358
12 P. K. Singh, B. Bhattacharya, R. K. Nagarale, Journal of Applied
Polymer Science, 118 (2010) 5, 2976–2980
13 A. Saxena, P. K. Singh, B. Bhattacharya, Mater. Tehnol., 47 (2013) 6,
799–802
14 K. Munindra, T. Tiwari, N. Shrivastava, Carbohydrate Polymer, 88
(2012), 54–60
15 R. Mishra, N. Baskaran, K. J. Rao, Solid State Ionics, 112 (1998),
261–273
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498–505
18 T. Heiser, U. Giovanella, S. Ould-Saad, M. Pa, Thin Solid Films,
(2006), 371–376
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Am. Chem. Soc., Div. Fuel Chem., 49 (2004) 2, 530–531
20 S. Ramesh, A. K. Arof, Material Science and Engineering B, 85
(2001), 11–15
21 D. F. Vieira, C. O. Avellaneda, A. Pawlicka, Electrochimica Acta, 53
(2007), 1404–1408
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rials- Rapid Communication, 7 (2013), 157–160
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Singh, Journal of Optoelectronics and Advanced Materials, 15
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A. SAXENA et al.: STRUCTURAL AND ELECTRICAL STUDIES OF FULLERENE (C60) ...
490 Materiali in tehnologije / Materials and technology 48 (2014) 4, 485–490
S. BALOS, L. SIDJANIN: EFFECT OF TUNNELING DEFECTS ON THE JOINT STRENGTH EFFICIENCY ...
EFFECT OF TUNNELING DEFECTS ON THE JOINT
STRENGTH EFFICIENCY OBTAINED WITH FSW
VPLIV TUNELSKIH NAPAK NA U^INKOVITOST TRDNOSTI
SPOJA, IZDELANEGA Z DRSNO ME[ALNIM VARJENJEM
Sebastian Balos, Leposava Sidjanin
Faculty of Technical Sciences, Department of Production Engineering, Trg Dositeja Obradovica 6, 21000 Novi Sad, Serbia
sebab@uns.ac.rs
Prejem rokopisa – received: 2013-08-14; sprejem za objavo – accepted for publication: 2013-09-13
In this paper, an attempt was made to study the effects of various types of tunneling defects in friction stir welding (FSW). Two
types of welding tool were used, one with a threaded and the other with a polygonal pin. To enable the formation of tunneling
defects, the shoulder-to-pin-volume ratio was larger than 1 for both tools, while the ratio of rotation speed to welding speed was
optimized as well. Tunneling defects were tested on EN-AW 5052 H32 and H38 alloys, having the thicknesses of 3 mm and 8
mm. It was found that the threaded tool produced a single or double triangular, as well as crack-shaped tunnel. The polygonal
tool produced a multiple triangular tunnel (the material flows toward the weld face) with a complex shape or a crack-type tunnel
with a complex shape. The most unfavorable tunnel obtained with the polygonal tool was the crack-shaped one, resulting in the
62 % and 46 % joint yield and ultimate tensile-strength efficiency. The corresponding values for the threaded tool were 40 %
and 36 %. The hardness efficiencies did not follow the trends of the proof and ultimate tensile strengths due to the prevention of
the tunneling during indentation. Another effect on the hardness is the presence of microcracks that may influence the apparent
decrease in the hardness in spite of the detected grain-refinement effect in the nugget. These effects influence the hardness
distribution in the weld, shifting the maximum hardness value into the thermo-mechanical zone.
Keywords: friction stir welding, aluminium alloy, tunneling defect, joint strength efficiency
^lanek predstavlja {tudij vpliva razli~nih vrst tunelskih napak pri drsno me{alnem varjenju (FSW). Uporabljeni sta bili dve vrsti
varilnega orodja, eno z navoji in drugo s poligonalnim nastavkom. Da bi prepre~ili nastanek tunelske napake, je bilo razmerje
med nosilcem in nastavkom ve~je od 1 pri obeh orodjih, medtem ko je bila optimirana tudi rotacija glede na hitrost varjenja.
Tunelska napaka je bila preizku{ena po EN-AW 5052 na zlitinah H32 in H38, debelih 3 mm in 8 mm. Ugotovljeno je bilo, da
orodje z navoji povzro~i enojni, dvojni trikotni tunel, kot tudi tunel v obliki razpoke. Poligonalno orodje povzro~i ve~
kompleksnih tunelov s trikotnim prerezom (material ste~e proti ~elu zvara) ali kompleksnih vrst tunelov v obliki razpok. Najbolj
neza`elen tunel, dobljen s poligonalnim orodjem, je bil v obliki razpoke, kar je povzro~ilo 62-odstotni in 46-odstotni izkoristek
spoja oziroma u~inkovitost natezne trdnosti. Ustrezne vrednosti pri orodju z navoji so bile 40 % in 36 %. U~inkovitost trdote ni
sledila tem usmeritvam in kon~ni natezni trdnosti zaradi izogibanja tunelom med merjenjem trdote. Nadaljnji vpliv na trdoto
imajo mikrorazpoke, ki lahko povzro~ijo navidezno zmanj{anje trdote, kljub odkritemu zmanj{anju zrn v jedru zvara. Ti pojavi
vplivajo na razporeditev trdote v zvaru in potiskajo vrednosti najve~je trdote v termomehansko obmo~je.
Klju~ne besede: drsno me{alno varjenje, aluminijeva zlitina, tunelska napaka, u~inkovitost trdnosti spoja
1 INTRODUCTION
Friction stir welding (FSW) is a solid-state metal-
joining process that uses the third body (a specialized
tool) to join two or more workpieces.1 FSW has attracted
considerable interest due to a number of advantages over
arc-welding processes.2 It has been shown that FSW may
be suitable for joining materials that are difficult to join
using conventional welding techniques such as 7XXX
and 2XXX aluminium alloys.3–7 Furthermore, Mg-alloys
and dissimilar materials have been successfully welded
with FSW.8–13
The microstructure and, therefore, the performance
of the FSW welds strongly depend on the tool geometry
as well as on the welding parameters such as the tool
rotation and welding speeds. One of the main differences
between the traditional arc-welding methods and FSW is
the existence of the thermo-mechanically affected zone
(TMAZ) near the heat-affected zone (HAZ) and the
nugget.14 The nugget takes the central place in the joint
region. It is the most severely deformed region, often
containing the equiaxed grains as the result of a dynamic
recrystallization.14 As the result of its position and the
fact that the tool directly influences it, the nugget, toge-
ther with the region above and under the nugget, may
represent the zone where defects can occur.15 One of the
most critical types of the FSW defect is the tunneling.
The tunneling defect occurs when the material flow
around the tool pin is not adequate, resulting in an irre-
gular weld filling.16–18 In accordance with the results pre-
sented by Radisavljevic et al.15, the ratio between the
rotation and the welding speeds must be optimized. At
the same time, an additional influencing factor might be
the ratio between the pin volume and the reservoir in the
tool shoulder (a concave shoulder volume) that should
not exceed one in order to enable the material to pile up
as a result of a pin indentation. In this work, an attempt
was made to study the influence of the tunneling-defect
shape and size on the mechanical properties of FSW
joints, using two types of aluminium alloy.
Materiali in tehnologije / Materials and technology 48 (2014) 4, 491–496 491
UDK 621.791:669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)491(2014)
2 EXPERIMENTAL PROCEDURE
Al-Mg plates in the half-hard condition (EN-AW
5052-H32) with a thickness of 3 mm and in the full-hard
condition (EN-AW 5052-H38) with a thickness of 8 mm
were used in this study. The chemical compositions of
the plates, determined with an optical emission spectro-
meter ARL 3580, are given in Table 1. The mechanical
properties of the workpiece materials, determined with a
ZDM 5/91 tensile testing machine, are given in Table 2.
The tensile properties were determined as the average of
three specimens. The sample dimensions were 300 mm ×
65 mm, forming a welded sample of 300 mm × 130 mm.
The samples were tightly placed in a steel fixture with a
backing plate into a 130 mm wide groove and secured
with clamps (Figure 1). The fixture was fitted on an
adapted Prvomajska UHG universal milling machine
with a power of 5.2 kW. In this study, two different types
of tool were used: plates 3 mm were welded with a poly-
gonal-type tool, while plates 8 mm were welded with a
threaded-type tool (Figure 2). Both tools were made of
X38CrMoV5-1 (H11) hot-work tool steel, having a che-
mical composition as shown in Table 3. The chemical
composition of the FSW tool material was measured
with an optical emission spectrometer ARL 2460. The
FSW tools were used in a heat-treated condition to 53
HRC. The tool concave-shoulder volume (reservoir) is
larger than the pin volume in both tools used, stimulating
the formation of a tunneling defect. For the polygonal-
type tool, the shoulder-to-pin-volume ratio was 5.3 and
for the threaded type it was 1.1, both being more than
sufficient to store the piled up material as a result of the
tool-pin indentation. Furthermore, the welding para-
meters were set to enable the formation of various types
of the tunneling defect (Table 4). These parameters were
set according to the results presented by Radisavljevic et
S. BALOS, L. SIDJANIN: EFFECT OF TUNNELING DEFECTS ON THE JOINT STRENGTH EFFICIENCY ...
492 Materiali in tehnologije / Materials and technology 48 (2014) 4, 491–496
Figure 2: a) Threaded and b) polygonal tools
Slika 2: Orodje: a) z navoji, b) poligonalno orodje
Table 4: FSW parameters
Tabela 4: Parametri FSW
Sample
Rotation
speed
r/min
Welding
speed
mm/min
Rotation
to welding
speed
r/mm
Tool
rotation
direction
Material
flow
direction
EN-AW 5052-H32; thickness 3 mm; polygonal-type tool
1 1230 17 72.35 right –
2 1230 12 102.50 right –
3 1230 23 53.48 right –
4 925 68 13.60 right –
5 925 91 10.16 right –
EN-AW 5052-H32; thickness 8 mm; threaded-type tool
6 645 5 129.00 left Up
7 645 5 129.00 right Down
8 645 12 53.75 right Down
9 645 49 13.16 right Down
Figure 1: Test setup
Slika 1: Preizkusni sestav
Table 1: Chemical composition of EN-AW 5052 aluminium alloys
(mass fractions, w/%)
Tabela 1: Kemijska sestava aluminijeve zlitine EN-AW 5052 (masni
dele`i, w/%)
Cu Mn Mg Si Fe Zn Ti Al
EN-AW
5052-H32 0.04 0.11 2.46 0.19 0.73 0.061 0.010
bal-
ance
EN-AW
5052-H38 0.03 0.29 2.33 0.19 0.29 0.011 0.022
bal-
ance
Table 2: Mechanical properties of EN-AW 5052-H32 and H38
Tabela 2: Mehanske lastnosti EN-AW 5052-H32 in H38
Proof
strength
Rp0.2 %/MPa
Ultimate
tensile
strength
Rm/MPa
Elongation
%
Vickers
hardness
Number
HV5
EN-AW
5052-H32 157 ± 3 214 ± 2 19 ± 1 67 ± 1
EN-AW
5052-H38 204 ± 7 245 ± 3 12 ± 1 90 ± 2
Table 3: Chemical composition of the X38CrMoV5-1 tool steel (w/%)
Tabela 3: Kemijska sestava orodnega jekla X38CrMoV5-1 (w/%)
C Si Mn P S Cr Mo V Fe
0.37 1.01 0.38 0.017 0.0005 4.85 1.23 0.32 balance
al.15, where the welding speeds were in the range of
6.45–8.22 r/mm and the tool rotation direction was able
to prevent the tunneling defect.
The mechanical properties of the FSW workpieces
were determined with the tensile, root-bending and hard-
ness tests as well as with a metallographic examination.
The tensile and bending tests were carried out with a
WPM ZDM 5/91 mechanical testing machine, according
to EN 895 and EN 910 standards, respectively. The hard-
ness was determined with a VEB HPO-250 Vickers test-
ing machine, with a load 5 kg. The hardness measure-
ments were done on the workpiece material (an average
of five indentations) before and after the welding, to
obtain the hardness profiles for the 1.5 mm distance bet-
ween the indentations. The metallographic examinations
were done with a light microscope Leitz Orthoplan, after
a standard metallographic preparation. The metallographic
preparation consisted of grinding (sandpaper, the grit of
220 to 2000), polishing (a diamond paste with (6, 3, 1 and
¼) μm particle sizes) and etching with Keller’s reagent
(2 mL HF, 3 mL HCl, 5 mL HNO3, 190 mL H2O).
3 RESULTS
3.1 Tensile properties
The tensile properties of FSW samples are presented
in Figure 3. In Samples 1–5, it can be seen that the
highest joint efficiencies are obtained with the higher
tool rotation speed. Within this rotation speed, the
highest joint proof-strength (PS) efficiency was obtained
with a high end-welding speed, while the highest joint
ultimate-tensile-strength (UTS) efficiency was obtained
with the lowest welding speed. The lowest efficiencies
were obtained with the lower tool rotation speed and the
higher welding speed. A similar result was obtained with
the EN-AW 5052-H38 alloy with a thickness of 8 mm,
welded with a conceptually different tool, see Samples
6–9.
3.2 Root-bending test
In Figure 4, the root-bending-test results are shown.
It can be seen that the highest angle of the first crack was
obtained on Sample 1 (a base-material thickness 3 mm),
while the lowest angle was obtained on Sample 6 (an
base-material thickness 8 mm). On samples 1–8, the
fracture did not occur up to 180°, while on Sample 9, the
fracture occurred at 134°.
3.3 Hardness results
The Vickers-hardness (HV5) profiles of the represen-
tative samples are shown in Figure 5. In Figures 5a and
5b, the hardness profiles for Samples 1 and 2 are shown,
depicting the highest and the lowest tensile properties of
the welded plates 3 mm. On the other hand, in Figures
5c and 5d, the hardness profiles for Samples 8 and 9 are
shown, depicting the highest and the lowest tensile pro-
S. BALOS, L. SIDJANIN: EFFECT OF TUNNELING DEFECTS ON THE JOINT STRENGTH EFFICIENCY ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 491–496 493
Figure 5: HV5 hardness profiles in relation to the distance from the
centerline: a) sample 1, b) sample 5, c) sample 8, d) sample 9
Slika 5: Profil trdote HV5 v odvisnosti od razdalje od sredi{~a: a)
vzorec 1, b) vzorec 5, c) vzorec 8, d) vzorec 9
Figure 3: Joint proof and the ultimate tensile strengths of the tested
samples
Slika 3: Preizkus spoja in kon~na natezna trdnost vzorcev
Figure 4: Results of the bending test. The values in parentheses
represent the samples where no fracture occurred.
Slika 4: Rezultati upogibnega preizkusa. Vrednosti v oklepajih so za
vzorce, kjer ni pri{lo do zloma.
perties of the welded plates 8 mm. In Figures 5a, 5c and
5d (Samples 1, 8 and 9) a common FSW hardness profile
may be observed, while for Sample 5 (Figure 5b), an
irregular profile was obtained. For Sample 1, the maxi-
mum hardness in the nugget (the weld centerline) is
higher than that of the base material. For the other sam-
ples, all the hardnesses of the weld, in the thermo-me-
chanical zone and the heat-affected zone, are lower than
that of the base material, indicating an insufficient grain
refinement in certain FSW zones.
3.4 Macro and microstructure
The weld cross-sections of the specimens are shown
in Figure 6, more exactly, in Figures 6a to 6e, the weld
cross-sections presented were obtained with the poly-
gonal tool, while the ones from the second column,
Figures 6f to 6i, were obtained with the threaded tool.
From these results it can be seen that with the polygonal
tool, two types of the tunneling defect were achieved: a
triangular (a closed tunneling defect) and a crack-shaped
defect (an open tunneling defect). Tunnels might occur
in a single (Figure 6b, Sample 2), double (Figures 6a
and 6d, Samples 1 and 4) or even triple configuration
(Figure 6c, Sample 3). The crack-type tunneling defect
obtained consisted of one longer and three smaller cracks
(Figure 6e, Sample 5).
On the other hand, when the threaded-type tool was
used, the tunneling defects with a more complex shape
were obtained. The threaded-type tool with the material
flow direction towards the tool (up) creates multiple
tunnels, closely following the tool pin geometry. Such
tunneling defects are spread across the whole weld
cross-section. The material flow towards the weld root
results in a complex-shaped tunneling defect in the weld
center or at the bottom (Figure 6f to 6i). Furthermore, a
combination of the complex-shaped tunneling defect and
the crack-shaped tunneling defect occurred on Sample 9
(Figure 6i).
The microstructures of the representative welded
samples are shown in Figure 7. In Figures 7a to 7c, the
Sample 5 FSW zones are presented, while in Figures 7d
to 7f, the Sample 9 FSW zones are shown. It can be seen
that the base-material microstructures, EN-AW 5052-
H32 and H38 shown in Figures 7a and 7d, correspond to
the typical cold-rolled alloy, with a clearly visible
deformed grain structure in the rolling direction. On the
other hand, the heat-affected zones (HAZs) (Figures 7b
and 7e, both on the left of the micrograph) show
coarsened recrystallized structures. In the thermo-mecha-
S. BALOS, L. SIDJANIN: EFFECT OF TUNNELING DEFECTS ON THE JOINT STRENGTH EFFICIENCY ...
494 Materiali in tehnologije / Materials and technology 48 (2014) 4, 491–496
Figure 6: Weld cross-sections showing the tunnel profiles: a) sample
1, b) sample 2, c) sample 3, d) sample 4, e) sample 5, f) sample 6, g)
sample 7, h) sample 8, i) sample 9
Slika 6: Prerezi zvara, ki prikazujejo profil tunela: a) vzorec 1, b)
vzorec 2, c) vzorec 3, d) vzorec 4, e) vzorec 5, f) vzorec 6, g) vzorec
7, h) vzorec 8, i) vzorec 9
Figure 7: Microstructures of samples 5 and 9: a) sample 5: base mate-
rial, b) sample 5: HAZ/TMAZ zone, c) sample 5: nugget, d) sample 9:
base material, e) sample 9: HAZ/TMAZ zone, f) sample 9: nugget
Slika 7: Mikrostruktura vzorcev 5 in 9: a) vzorec 5: osnovni material,
b) vzorec 5: podro~je HAZ/TMAZ, c) vzorec 5: jedro, d) vzorec 9:
osnovni material, e) vzorec 9: podro~je HAZ/TMAZ, f) vzorec 9:
jedro
nical zones (TMAZs), shown in Figures 7b and 7e (both
to the right of the micrograph), finer microstructures are
formed as a result of the local deformation. The main
difference between Figures 7c and 7f, where the nugget
is depicted, is the occurrence of an intensive cracking,
seen on Sample 5. In the Sample 9 nugget (Figure 7f),
the microstructure is finer than that of the base material,
HAZ and TMAZ.
4 DISCUSSION
4.1 Joint effectiveness to microstructure correlation
Joint PS, UTS and hardness effectiveness are given in
Figure 8. According to the joint effectiveness shown in
Figure 8, it can be seen that the joint UTS effectiveness
is the lowest, followed by the joint PS effectiveness. The
joint UTS effectiveness is lower than the values obtained
by Radisavljevic et al. (by up to 80.1 %)15, which is the
result of the tunneling-type defect, as depicted in Figure
4. On the other hand, the joint hardness effectiveness
matches the values obtained by Radisavljevic et al.15 This
is the result of the measuring method, according to
which the indentations were made in the tunnel-free
region of the weld, as well as the method of calculation.
Namely, the maximum hardness was taken into account.
However, although there is a correlation between the
proof strength and the ultimate tensile strength, these
two parameters cannot be correlated to the hardness
effectiveness. Nevertheless, the hardness distributions
may be correlated to the microstructures obtained, Fig-
ures 6 and 7, especially for the representative Samples 5
and 9. In Sample 5, the hardnesses of the nugget, TMAZ
and HAZ are all similar, which is in contrast to the hard-
ness values for Sample 9. There the nugget has a higher
hardness than in the case of TMAZ and HAZ. As shown
in Figure 7, the Sample 5 nugget contains a number of
cracks that might have influenced the values that are
similar to the ones for TMAZ and HAZ. On the other
hand, in Sample 9, a crack-free nugget was obtained and,
consecutively, higher hardnesses were measured than
those for TMAZ and HAZ. Such a hardness distribution
is the result of the grain refinement in the nugget. The
nugget-hardness values are lower compared to the base
material, which might be influenced by two factors: the
grain refinement and the recrystallization.
4.2 Influence of the tunnel shape
The mechanical properties of the joints and their effi-
ciencies clearly vary in accordance with the type of the
FSW tool geometry. The two types of the FSW tool have
different influences on the joint properties. It can be seen
that the polygonal tool (Samples 1–5) produces joints
with higher mechanical properties than the threaded tool
(Samples 6–9) in spite of the higher shoulder-to-pin-
volume ratio of 5.3 versus 1.1. The higher mechanical
properties are influenced by the tunnel obtained. Name-
ly, the polygonal tool produces triangular-shaped tunnels
or crack-shaped tunnels that are the least optimal. The
most optimum sample obtained with the polygonal-type
tool is Sample 1 with a double tunnel, while the lowest
mechanical properties were found for Sample 5 with a
crack-shaped tunnel. Regarding the threaded-tool sam-
ples, the mechanical properties of Sample 8 are the
highest, while the lowest mechanical properties were
found for Sample 9. On Sample 9, a combination of a
tunnel and a crack-shaped tunnel was found, which pro-
ved to have a more adverse effect on the mechanical
properties compared to Sample 6, where multiple tunnels
were found, closely following the tool-pin profile. The
relatively high welding speeds applied in this study
clearly did not contribute to the joint quality, regardless
of the tool type.
5 CONCLUSIONS
According to the presented results, the following
conclusions can be drawn:
• The UTS efficiency of the joints with the tunneling
defect was lower by 25–82 % than that of the joints
without the tunneling defect.
• The efficiency of the joint PS is higher than that of
the ultimate tensile strength, while the hardness
efficiency is higher than the joint PS efficiency.
• The tunnel profile is influenced by the pin geometry
and pin direction determining the material flow
direction.
• The polygonal-type tool produces a triangular tunnel
shape, in single and multiple forms, as well as a
crack-type tunneling defect.
• The threaded-type tool produces a triangular and
complex-shaped tunnel, depending on the material
flow direction. If the material flows towards the weld
root a complex-shaped tunnel is formed with superior
mechanical properties. The only exception is the
combination of a complex-shaped and crack-type
tunnel. In the case of an upward material flow,
S. BALOS, L. SIDJANIN: EFFECT OF TUNNELING DEFECTS ON THE JOINT STRENGTH EFFICIENCY ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 491–496 495
Figure 8: Joint PS, UTS and hardness effectiveness
Slika 8: U~inkovitost spoja PS, UTS natezne trdnosti in trdote
multiple triangular tunneling defects are formed that
are very unfavorable with respect to mechanical
properties.
• The lowest mechanical properties both in terms of the
tensile strength and bending were obtained with the
tunneling defects associated with cracks.
• The highest joint efficiency was obtained with the
polygonal-pin profile.
• The pin-to-concave-shoulder-volume ratio greatly
influences the occurrence of the tunneling defect.
Aknowledgements
The authors are grateful to Mr. Michael A. Maier for
his technical assistance.
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496 Materiali in tehnologije / Materials and technology 48 (2014) 4, 491–496
M. \OKI] et al.: THE INFLUENCE OF LAMINATION AND CONDUCTIVE PRINTING INKS ...
THE INFLUENCE OF LAMINATION AND CONDUCTIVE
PRINTING INKS ON SMART-CARD OPERABILITY
VPLIV LAMINACIJE IN PREVODNIH TISKARSKIH BARV NA
DELOVANJE PAMETNIH KARTIC
Miloje \oki}1, Vasa Radoni}2, Vladan Mladenovi~3, Anton Pleter{ek4,
Ur{ka Kav~i~5, Ale{ Hladnik1, Vesna Crnojevi}-Bengin2, Tadeja Muck1
1Faculty of Natural Sciences and Engineering, University of Ljubljana, Sne`ni{ka 5, 1000 Ljubljana, Slovenia
2Faculty of Technical Sciences, University of Novi Sad, Trg Dositeja Obradovi}a 6, 21000 Novi Sad, Serbia
3Cetis, d. d., ^opova ulica 24, 3000 Celje, Slovenia
4ams R&D, d. o. o., Tehnolo{ki park 21, 1000 Ljubljana, Slovenia
5Valkarton Rakek, d. o. o., Partizanska cesta 7, 1381 Rakek, Slovenia
miloje.djokic@gmail.com
Prejem rokopisa – received: 2013-08-15; sprejem za objavo – accepted for publication: 2013-10-24
In this paper a novel design of UHF RFID antennas for smart cards is proposed. Smart cards usually include HF RFID antennas
produced by means of hot foil stamping or etching – a subtractive process. Making HF antennas using printing – an additive
process – is difficult, because it necessitates the printing of three layers: that of conductive lines (a coil), followed by a dielectric
layer and once again a conductive layer. On the other hand, the printing of UHF antenna requires only the single-layer printing
of conductive lines, i.e., antenna, thus enabling a more rational printing.
After designing, simulating and testing a folded UHF dipole antenna for RFID applications, the impact of three different
conductive printing inks, reading distance and card deformation on the backscattered signal power was studied. A three-way
analysis of variance (ANOVA) was implemented to aid in the interpretation of the measurement results obtained on both
non-laminated and laminated cards. It was demonstrated that smart cards with screen-printed UHF antennas can be fabricated
using the proposed design optimization, so achieving good card readability.
Keywords: smart cards, printed antenna, antenna design, lamination, conductive printing inks, UHF RFID, analysis of variance
V ~lanku je predlagana nova izvedba UHF RFID-anten za pametne kartice. Pametne kartice navadno vklju~ujejo HF
RFID-antene, izdelane na podlagi vro~ega tiska s folijo ali z jedkanjem (subtraktivni proces). Izdelava HF-anten v tisku (aditivni
proces) je te`avna, saj mora biti tisk izveden v treh plasteh – tisk prevodnih linij (tuljava), dielektri~ne plasti in nato ponovno
prevodne plasti. Po drugi strani pa je `elen enoplastni tisk prevodnih linij UHF-antene, kar omogo~a racionalnej{i tisk.
Po na~rtovanju, simulaciji in preizku{anju dipolne UHF-antene za RFID-aplikacije je bil preu~evan vpliv treh prevodnih
tiskarskih barv, razdalje od~itavanja in deformacije kartic na mo~ povratnega signala. Pri razlagi dobljenih rezultatov meritev
nelaminiranih in laminiranih kartic je bila uporabljena trosmerna analiza variance (ANOVA). Izkazalo se je, da je mogo~e
izdelati pametne kartice s tehniko sitotiska z opisano optimizacijo izvedbe UHF-antene, kar omogo~a dobro ~itljivost kartic.
Klju~ne besede: pametne kartice, tiskana antena, izvedba antene, laminacija, prevodne tiskarske barve, UHF RFID, analiza
variance
1 INTRODUCTION
Printed flexible electronics is becoming more viable
with recent technological developments, opening up
possibilities to integrate radio frequency (RF) functions
into clothing or moving objects to monitor various para-
meters of interest. Printed electronics that combine elec-
tronic components with different printing technologies
allows the production of miniature low-cost electronic
systems.
The first connection between printing technology and
conventional electronics appeared in the 1980s when
Teng and Vest made metal conductors for solar cells
using inkjet printing.1 This opened up the possibility to
produce various electronic components using different
printing techniques. Generally, two printing techniques
are the most common, i.e., screen printing2–8 and inkjet
printing,9–12 but others such as flexography,13 gravure
printing14 and pad printing15 are also used. Conventional
printing techniques, such as screen printing or gravure
printing, are characterized by high production speed,
high resolution and the possibility of using a number of
different conductive inks, so these techniques are
suitable for the mass production of precise electronics.11
Printed electronics has many advantages, including a
simple production process, the reduced consumption of
raw materials and high productivity (especially in roll-
to-roll processes). One of the major applications of
printed electronics is radio-frequency identification
(RFID). RFID is an automatic identification technology
that uses a wireless non-contact radio frequency to trans-
fer data for the purposes of automatically identifying and
tracking the tags attached to objects. RFID systems are
widely used in various real-life applications, such as
access control or smart-card systems, public transport
and logistics.16–18 RFID systems consist of at least of a
reader, which receives the feedback signals sent back
from the tags, and the tags themselves. The simplest,
passive tags consist of an antenna and an integrated
Materiali in tehnologije / Materials and technology 48 (2014) 4, 497–504 497
UDK 655.1:621.396/.398 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)497(2014)
circuit (IC). An integrated controller and a memory in
the IC chip enable the activity of the tag and contain
electronically stored information. The second part of the
tag is the antenna that receives and transmits the radio
signals.
Since the mechanical and electrical characteristics of
antennas depend on the materials used, several experi-
mental studies have recently been carried out to inve-
stigate the effects of conductive pastes,3 to compare
different printed materials19 and to look at the influence
of substrate folding.12 The effects of different materials
used in the production of UHF RFID antennas on their
performance have been analysed.19 It has been shown
that the conductive material used to print the antenna has
a significant influence on the antenna’s return loss. A
similar study has confirmed that a conductive silver-
based paste outperformed all the other types of conduc-
tive pastes and provided a performance similar to that of
its copper equivalent.3
This paper presents a novel design for UHF RFID
smart cards, realized using the basic functionality offered
by printing technology. The impacts of different conduc-
tive inks that are used to print the UHF antenna are
investigated. In addition, the influence of the lamination
process and card deformation on the performance of the
cards is investigated in detail.
2 THEORETICAL
The reading range in the free space of an RFID
system operating at wavelength  depends on the reader
power density (PTX
rdr ), the antenna gain (Grdr) and the tag
antenna gain (Gtag). The power received by the tag (PRX
tag
)
at distance r from the reader is estimated using the Friis
equation:20
P P G G
rRX
tag
TX
rdr
rdr tag=
⎛
⎝
⎜ ⎞
⎠
⎟
4
2
π
(1)
The transmitted signal power is denoted by TX, while
the received signal is denoted by RX. We can define the
gain of the tag antenna Gtag and the backscatter transmis-
sion loss Tb using the relationships from the Friis equa-
tion, as expressed in equations Eq. 2 to 4:
P P G G
r
TTX
tag
TX
rdr
rdr tag b=
⎛
⎝
⎜ ⎞
⎠
⎟
4
2
π
(2)
P P G G
rRX
rdr
TX
tag
tag rdr=
⎛
⎝
⎜ ⎞
⎠
⎟
4
2
π
(3)
P P T G G
rtagRX
rdr
TX
rdr
b rdr=
⎛
⎝
⎜ ⎞
⎠
⎟2 2
4
4

π
(4)
The backscattered modulated signal received by the
reader is therefore proportional to r–4. Nevertheless, the
communication between the passive tag and the reader is
forward link limited, since the power to power on the IC
(Pmin
tag
) is by at least 6 orders of magnitude higher than the
minimum power (Pmin
rdr ) required to demodulate the
backscattered signal. The forward limited range in free
space (Rforward) follows from Eq. 5:
R
P G G
P
forward
TX
rdr
rdr tag
min
tag) =
⎛
⎝
⎜ ⎞
⎠
⎟
4π
(5)
A simple estimation for a reader with the power PTX =
500 mW, a tag in the form of a /2 dipole, with gain Gtag
= 2.2 dBi and minimum power needed to turn on the IC
Pmin
tag
= 100 μW, gives a forward range of 2.4 m in free
space. Until this point, a perfect matching between the
antenna and load (IC) has been assumed with the
power-transfer coefficient 
 = 1. Generally speaking, the
power-transfer coefficient is defined by Eq. 6:

 =
+
4
2
R R
Z Z
chip ant
chip ant
(6)
where the impedance of the chip (Zchip) and the antenna
(Zant) is the sum of the resistance and the reactance. The
transfer coefficient is a maximum when the reactance of
the chip and antenna cancel each other. In this case,
referred to as conjugate matching, the transfer coeffi-
cient is equal to 1. In all other cases, the transfer coeffi-
cient is lower.
The effective gain (Geff) of the tag antenna (the ability
to increase the power or amplitude of the signal) can be
defined as follows6:

 
G Deff = (7)
where D is the directivity of the antenna,  is the an-
tenna efficiency and 
 is the power transfer efficiency.
3 EXPERIMENTS
The block diagram of the smart-card fabrication
process is shown in Figure 1a. First, the UHF antenna
was designed using the EM simulator CST Microwave
Studio and then printed and dried. In the next step, the
chips were assembled and the lamination performed to
fabricate a final smart card. The fabricated cards were
experimentally tested following the process shown in
Figure 1b. The characteristics of the flat non-laminated
antenna were measured in an anechoic chamber, while
the backscattering characteristics of the smart cards were
evaluated in a real environment using an ams-R902
reader. Furthermore, the effect that bending has on the
initial planar topology in the case of the laminated cards
was also investigated. Multifactorial analysis of variance
(ANOVA) was implemented to aid in the interpretation
of the measurement results obtained for both non-lami-
nated and laminated cards.
3.1 Antenna design and simulation
The novel design configuration of the proposed
folded dipole antenna is shown in Figure 2. The antenna
M. \OKI] et al.: THE INFLUENCE OF LAMINATION AND CONDUCTIVE PRINTING INKS ...
498 Materiali in tehnologije / Materials and technology 48 (2014) 4, 497–504
is designed to operate according to the UHF standard
with a central operating frequency of 868 MHz. In
addition, the dipole is folded in order to fit the standard
credit-card size, with a negligible reduction in antenna
efficiency. The antenna was designed for a thick poly-
carbonate film 120 μm (grammage: 120 g/m2; surface
roughness, ISO 4288: 2.3 μm) with a relative permittivity
of r = 3.2 and a dissipation factor 0.0019. All the
simulations were performed using CST Microwave
Studio. For the conductive material, 16 μm silver was
used and the conductor losses were modelled using the
bulk conductivity for silver.
As the number and thickness of the dielectric layers
during lamination affects the characteristics of the anten-
nas, initially the antenna dimensions were determined for
the laminated case. Nine dielectric layers were used in
the simulation, with the antenna placed between the
fourth and fifth layers (Figure 3).
The dimensions of the antenna are carefully optimi-
zed in order to obtain the central frequency of 868 MHz.
The final antenna dimensions are shown in Figure 2,
with all the marked dimensions in mm. The total antenna
size is only 37 mm × 41 mm. Figure 4 compares the
simulated return loss of the laminated and non-laminated
antennas. It is clear that the lamination process signifi-
cantly influences the resonant frequency of the antenna,
M. \OKI] et al.: THE INFLUENCE OF LAMINATION AND CONDUCTIVE PRINTING INKS ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 497–504 499
Figure 1: a) Block diagrams of smart-card fabrication, b) testing and
measurement process
Slika 1: a) Blokovna diagrama izdelave pametnih kartic, b) procesa
preizku{anja in merjenja
Figure 4: Simulated return losses of the laminated and non-laminated
antennas
Slika 4: Simulacija izgube povratnega signala laminirane in nelami-
nirane antene
Figure 2: Proposed folded dipole antenna (mm)
Slika 2: Nova izvedba dipolne antene (mm)
Figure 3: Cross section of the laminated card with antenna placed
inside
Slika 3: Pre~ni prerez laminirane kartice z vklju~eno anteno v vmesni
plasti
Figure 5: Simulated radiation pattern of the laminated antenna
Slika 5: Simulacija sevalnega diagrama laminirane antene
due to changes in the effective dielectric constant of the
substrate.
The simulated radiation patterns of the proposed
laminated dipole antenna in the E- and H-planes are
shown in Figure 5. The total antenna gain is above 1.9
dBi.
3.2 Printing and drying
Antenna samples were printed using the Siasprint
Novaprint-P screen printer with a monofilament poly-
ester plain weave mesh with 120 l/cm and a theoretical
ink volume of 16.3 cm3/m2. For printing, three different
silver conductive printing inks were used:
• SC-CRSN2442 SunTronic 280 Thermal Drying
Silver Conductive Ink,
• DP-DuPont 5064H Silver Conductor,
• UV-Electra Polymers ELX 30.
Inks denoted as SC and DP require thermal drying,
while UV is an ultraviolet curing ink. The characteristics
of the used inks are summarized in Table 1.
After printing, the optimal drying process for each
ink was determined. The optimal drying was determined
as the point where the sheet resistance of the printed
samples became constant and did not get any lower with
a longer drying time, higher temperature or higher or
longer UV exposure. In accordance with the SC and DP
printing inks’ specifications, drying in an infra-red (IR)
tunnel was performed. The best results – the lowest sheet
resistance – were obtained when the samples passed the
IR tunnel seven times (at 72 °C for 30 s). The UV prin-
ted samples were optimally cured when twice passed
through a UV tunnel with a dose of 700 mJ/cm2.
3.3 Sheet resistance
The resistance, R, of the printed conductive lines was
measured using a DT-890G digital multimeter on the test
element, as shown in Figure 6, between points 1 and 2.
The nominal length, L, between points 1 and 2 was 22
mm and the line width, W, was set to 3 mm. The resis-
tance was measured after 24 h of conditioning with a
50 % relative humidity at 23 °C. The nominal number of
squares Nsq and the final sheet resistance Rsh of the con-
ductive ink layer (m) can be calculated using Eq. 8 and
Eq. 9, respectively. The measured resistances for the
different inks are shown in Table 1.
N
L
WSQ
= =103. (8)
R
R
Nsh SQ
= (9)
3.4 Chip integration
In our study, NXP strap chips were used. The main
chip characteristics are as follows: operating frequency
840–960 MHz, impedance 22 – j195  and Q-factor 9.
The chips were assembled with isotropic conductive glue
based on silver particles (Bison, Netherlands) after the
drying process. After chip integration, tags were dried in
a thermal oven at 120 °C for 30 min.
3.5 Lamination
The lamination process was used to fabricate the final
smart cards. It was separated into two phases: first the
"heat phase", where nine foil layers were assembled at a
temperature of 199 °C and a pressure of 300 N/cm2 for a
fixed period of 19 min. In the second phase, the "cold
phase", the assembled foils were exposed to 25 °C and a
pressure of 500 N/cm2 for 18 min. Finally, cutting into
final cards sized according to ISO/IEC 7810 standard
was carried out. The final prototype of the laminated
smart card is shown in Figure 7.
M. \OKI] et al.: THE INFLUENCE OF LAMINATION AND CONDUCTIVE PRINTING INKS ...
500 Materiali in tehnologije / Materials and technology 48 (2014) 4, 497–504
Figure 6: Test element for resistance measurements
Slika 6: Preizkusni element za merjenje upornosti
Table 1: Properties of investigated printing inks
Tabela 1: Lastnosti tiskarskih barv
Printing ink Solids(%)
Recommended drying
condition
Specified resistance*
Rsh (spec.)/m
(layer thickness)
Measured resistance**
Rsh (meas.)/m
SC
CRSN2442 SunTronic 280 69–71
100–130 °C over 30–90 s
in the hot zone 10–32 (25 μm) 118.12
DP
DuPont 5064H 63–66
140 °C over 120 s in the
hot zone  6 (25 μm) 567.96
UV
Electra Polymers ELX 30
approx.
75
1000–1500 mJ cm–2 300–500 (15 μm) 653.72
*Sheet resistance from the product specification
**Sheet resistance measured on printed samples using the DT-890G digital multimeter. An average thickness of 15 μm was determined on the
prints’ cross-sections with a scanning electron microscope.
4 RESULTS AND DISCUSSION
4.1 Non-laminated antenna measurements
The return loss of the fabricated non-laminated SC
antenna was measured and the comparison of the simu-
lated and measured return losses for the non-laminated
antenna is given in Figure 8, with a photograph of the
fabricated circuit shown in an inset. Despite the small
size of the antenna, good agreement is observed and the
fundamental resonance occurs at 940 GHz, as predicted,
with a reflection better than –10 dB.
Figure 9 shows the comparison of the simulated and
measured radiation patterns for the non-laminated
antenna. The pattern measurements were taken in an
anechoic chamber using a vector network analyser and
the measured patterns agree well with the simulated
ones, except for some ripple, which is ascribed to reflec-
tions in the chamber.
4.2 Card evaluation
The realized readability of the cards with RFID tags
was measured in a real environment using an IDS-R902
reader that consists of the reader electronics and an
A0025 circularly polarized patch antenna (Poynting
GmbH, Dortmund, Germany) with gain of 6.5 dBi
emitting UHF EM radiation at frequency f = 868 MHz.
The reader electronics measures the intensity of the
modulated backscattered signal.
The cards with UHF RFID tags were evaluated by
measuring the backscattered power (dBm), moving the
tag in a straight line perpendicular to the reader in 2 cm
increments, starting at 2 cm from the emitting antenna up
to the estimated free space working distance (Figure 10).
Multifactorial (multiway) analysis of variance
(ANOVA) was implemented to aid in the interpretation
of the measurement results obtained for both non-lami-
nated and laminated cards. The impact of the three selec-
ted factors on the response – backscattered power – was
M. \OKI] et al.: THE INFLUENCE OF LAMINATION AND CONDUCTIVE PRINTING INKS ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 497–504 501
Figure 10: Measurement of the backscattered power
Slika 10: Merjenje mo~i povratnega signala
Figure 8: Comparison of simulated and measured return losses for the
non-laminated antenna with inset photograph of the fabricated circuit
Slika 8: Primerjava simulirane in izmerjene povratne izgube signala
za nelaminirano anteno z vklju~eno sliko natisnjenega vezja
Figure 7: Final laminated smart card prototype
Slika 7: Kon~ni prototip laminirane pametne kartice
Figure 9: Comparison of simulated and measured radiation patterns in
E-plane and H-plane for non-laminated antenna
Slika 9: Primerjava simuliranega in izmerjenega sevalnega diagrama v
E-ravnini in H-ravnini za nelaminirane antene
M. \OKI] et al.: THE INFLUENCE OF LAMINATION AND CONDUCTIVE PRINTING INKS ...
502 Materiali in tehnologije / Materials and technology 48 (2014) 4, 497–504
Figure 12: Laminated cards: a) effect of ink type and b) flatness on
backscattered signal power
Slika 12: Laminirane kartice: a) vpliv vrste tiskarske barve in b) plo-
skosti na mo~ povratnega signala
Figure 11: Non-laminated cards: a) effect of ink type and b) ink-type/
reading distance interaction on the backscattered signal power
Slika 11: Nelaminirane kartice: a) vpliv vrste tiskarske barve in b)
interakcije vrsta tiskarske barve/razdalja od~itavanja na mo~ povrat-
nega signala
Figure 14: Laminated cards: a) effects of interactions – ink-type/
reading distance and b) flatness/reading distance on backscattered
power
Slika 14: Laminirane kartice: a) vpliv interakcije vrsta tiskarske barve
– razdalja od~itavanja in b) ploskost – razdalja od~itavanja na mo~ po-
vratnega signala
Figure 13: Laminated cards effect of ink type on backscattered power
for: a) flat and b) bent cards
Slika 13: Laminirane kartice: vpliv vrste tiskarske barve na mo~ po-
vratnega signala za: a) ravne in b) upognjene kartice
studied: ink type, reading distance and flatness, the later
one only for the laminated cards (flat or bent). The cards
were bent to a final deformed position: 2.6 cm height, 3
cm width, 1.73269 cm radius and an angle of 240°
subtended by an arc. The results of the statistical analysis
at  = 0.05 (i.e., a 95 % confidence level) are presented
in Figures 11a, 12a, 12b, 13a and 13b. The analyses
were made on nine samples for each of the various cards
(non-laminated, flat laminated and bent laminated).
4.3 Non-laminated cards
All of the examined non-laminated cards were flat, so
here only two factors – ink type and reading distance –
were analysed. As shown in the ANOVA table (Table
2a), both factors have a statistically significant effect
(P-value < 0.05) on the response. The very small interac-
tion effect (F-ratioAB = 6.16) is merely due to an extre-
mely strong impact of the reading distance and can be
neglected. Figure 11a and the multiple range test data
(Table 2b) clearly show that the SC ink behaves in a
substantially different manner compared to the other two
inks: the mean backscattered signal power of the UHF
RFID cards with the applied SC ink is –49.9 dBm, while
the mean responses with the DP and UV inks are –51.0
dBm and –51.1 dBm, respectively. The difference in the
responses between the latter two inks is not statistically
significant (Table 2b), since the corresponding mean
values form the same homogeneous group (a). The ink
type/reading distance diagram (Figure 11b) also
supports the above findings.
4.4 Laminated cards
In addition to the impacts of the ink type and reading
distance, the flatness of the substrate was studied in the
case of the laminated cards. As shown in Table 3a,
reading distance has by far the strongest effect on the
backscattered power, followed by the effects of the ink
type and flatness. The impact of the three ink types on
the response (Figure 12a) is similar to that found with
the non-laminated cards, the only difference being that
the mean response with the UV ink (–52.2) is now
statistically higher compared to the one with the DP ink
(–52.0) (Table 3b). The detected backscattered power
with the flat cards is generally lower than with the bent
cards (Figure 12b).
The compared maximum reading distances of the
non-laminated and laminated cards showed a significant
difference (presented in Figures 11b and 14a). Since the
antenna is designed for laminated cards, the lamination
process increases the reading distance from 40 cm to 60
cm.
To find out how flatness affects the measured back-
scattered power, we ran two more ANOVAs, separately
for the flat and for the bend cards. The results (Figure
13) indicate that the differences in the effects of the DP,
UV and SC inks on the response become more pro-
nounced with bending (Figure 13).
It is clear that the readability of the cards depends on
the conductivity of the printing ink used for printing the
antenna. Antennas printed with SC ink (Rsh = 118.12
m) had 5 to 6 times higher conductivity than those with
DP and UV inks and, as a consequence, the cards printed
M. \OKI] et al.: THE INFLUENCE OF LAMINATION AND CONDUCTIVE PRINTING INKS ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 497–504 503
Table 3: Laminated cards: a) ANOVA table, b) multiple range test
results for ink type and c) flatness; SS = Sum of Squares, Df = De-
grees of Freedom, MS = Mean Squares
Tabela 3: Laminirane kartice: a) ANOVA tabela, b) rezultati preiz-
kusov ve~kratnih primerjav za vrsto tiskarske barve in c) ravnole`no-
sti; SS = vsota kvadratov, Df = stopinje prostosti, MS = povpre~je kva-
dratov
Table 2: Non-laminated cards: a) ANOVA table and b) multiple range
test results for ink type; SS = Sum of Squares, Df = Degrees of
Freedom, MS = Mean Square
Tabela 2: Nelaminirane kartice: a) ANOVA tabela in b) rezultati
preizkusov ve~kratnih primerjav za vrsto tiskarske barve; SS = vsota
kvadratov, Df = stopinje prostosti, MS = povpre~je kvadratov
with SC ink had a better readability or a higher back-
scattered power (dBm) of the return signal. The maxi-
mum reading (working) distance showed no significant
differences between the applied inks. Furthermore, the
maximum reading distance was higher for the laminated
cards, with an approximately 20 cm increase, regardless
of the ink used.
With card deformation, we are assuming that the
bending increases the reading distance range. This can
be seen from Figure 14b, where the strength of the
modulated backscattered signal is improved after bend-
ing. The result of the performed analysis shows that the
printed cart printed with SC inks outperforms all the
other printed cards in terms of the backscattered power.
5 CONCLUSIONS
In the present study, a novel folded UHF RFID
antenna that requires only single-layer printing has been
proposed and implemented in the smart-card design.
Despite the small size of the antenna, good readability
was achieved. Through optimizing the antenna design,
more rational printing for low-cost mass production can
be possible.
We showed that a higher ink conductivity increases
the backscattered power and that the novel design of the
UHF antenna is appropriate for final laminated smart
cards, while the lamination process increases the maxi-
mum reading distance.
By the end of our research, we had shown that smart
cards with screen-printed UHF antennas can be realized.
The main effect on the final tag’s operability is that of
the quality of the conductive ink itself, while bending
can improve the strength of the modulated backscattered
signal. The study of chip contacting, especially the
glue/ink interactions and the influence of the mechanical
strength of smart cards, will be the topic of our future
research.
Acknowledgments
Ur{ka Kav~i~ acknowledges assistance from the
European Social Fund ("Operation part-financed by the
European Union, European Social Fund").
6 REFERENCES
1 K. F. Teng, R. W. Vest, Metallization of solar cells with ink jet
printing and silver metallo-organic inks, IEEE Transaction on
Components, Hybrids and Manufacturing Technology, 11 (1988),
291–297
2 D. Y. Shin, Y. Lee, C. H. Kim, Performance characterization of
screen printed radio frequency identification antennas with silver
nanopaste, Thin Solid Films, 517 (2009), 6112–6118
3 K. Janeczek, M. Jakubowska, G. Kozio³, A. M³oz³niak, A. Araz´na,
Investigation of ultra-high-frequency antennas printed with polymer
pastes on flexible substrates, IET Microwaves, Antennas & Propa-
gation, 6 (2012), 549–554
4 T. Elsherbeni, K. E. Mahgoub, L. Sydanheimo, L. Ukkonen, F. Yang,
Laboratory scale fabrication techniques for passive UHF RFID tags,
Antennas and Propagation Society International Symposium
(APSURSI), 2010 IEEE, Toronto, 2010
5 K. Janeczek, M. Jakubowska, A. M³oz³niak, G. Koziol, Thermal cha-
racterization of screen printed conductive pastes for RFID antennas,
Materials Science and Engineering B: Solid-State Materials for
Advanced Technology, 177 (2012), 1336–1342
6 U. Kav~i~, L. Pavlovi~, M. Pivar, M. \oki}, B. Batagelj, T. Muck,
Printed electronics on recycled paper and cardboards, Inf. MIDEM,
43 (2013) 1, 50–57
7 U. Kav~i~, M. Ma~ek, T. Muck, M. Klanj{ek Gunde, Readability and
modulated signal strength of two different ultra-high frequency radio
frequency identification tags on different packaging, Package. Tech-
nol. Sci., 25 (2012) 7, 373–384
8 U. Kav~i~, M. Ma~ek, T. Muck, Impact study of disturbance on read-
ability of two similar UHF RFID tags, Inf. MIDEM, 41 (2011) 3,
197–201
9 X. Nie, H. Wang, J. Zou, Inkjet printing of silver citrate conductive
ink on PET substrate, Applied Surface Science, 261 (2012), 554–560
10 A. Kanso, E. Arnaud, T. Monediere, M. Thevenot, E. Beaudrouet, C.
Dossou-yovo, R. Noguera, Inkjet printing of Coplanar Wire-Patch
antenna on a flexible substrate, 15th International Symposium on An-
tenna Technology and Applied Electromagnetics, Toulouse, 2012,
1–4
11 J. Li, B. An, J. Qin, Y. Wu, Nano copper conductive ink for RFID
application, International Symposium on Advanced Packaging Mate-
rials, Xiamen, 2011, 91–93
12 V. Radonic, K. Palmer, G. Stojanovic, V. Crnojevic-Bengin, Flexible
Sierpinski Carpet Fractal Antenna on a Hilbert Slot Patterned
Ground, International Journal of Antennas and Propagation, Vol.
2012, Article ID 980916, 2012, doi: 10.1155/2012/980916
13 L. B. Keng, W. L. Lai, C. W. A. Lu, B. Salam, Processing and cha-
racterization of flexographic printed conductive grid, Electronics
Packaging Technology Conference, Singapore, 2011, 517–520
14 D. A. Alsaid, E. Rebrosova, M. Joyce, M. Rebros, M. Z. Atashbar, B.
Bazuin, Gravure Printing of ITO Transparent Electrodes for Applica-
tions in Flexible Electronics, Journal of Display Technology, 8
(2012), 391–396
15 S. L. Merilampi, P. Ruuskanen, T. Björninen, L. Sydänheimo, L.
Ukkonen, Characterization of UHF RFID tags fabricated directly on
convex surfaces by pad printing, International Journal of Advanced
Manufacturing Technology, 53 (2011), 577–591
16 Z. J. Tang, Y. G. He, Y. Wang, Broadband RFID tag antenna with
quasi-isotropic radiation performance, AEU-International Journal of
Electronics and Communications, 65 (2011), 859–863
17 D. M. Dobkin, The RF in RFID, Passive UHF RFID in Practice,
Elsevier, 2008
18 K. Finkenzeller, D. Muller, RFID Handbook: Fundamentals and
Applications in Contactless Smart Cards, Radio Frequency Identifi-
cation and Near-Field Communication, 3rd edition, John Wiley &
Sons, Inc., 2010
19 A. Syed, K. Demarest, D. D. Deavours, Effects of antenna material
on the performance of UHF RFID tags, IEEE International Confe-
rence on RFID, Grapevine, 2007, 57–62
20 C. A. Balanis, Antenna Theory: Analysis and Design, 3rd edition,
John Wiley & Sons, Inc., 2005
M. \OKI] et al.: THE INFLUENCE OF LAMINATION AND CONDUCTIVE PRINTING INKS ...
504 Materiali in tehnologije / Materials and technology 48 (2014) 4, 497–504
S. KRAMAR et al.: CHARACTERIZATION OF THE SUBSTRATES FROM TWO CULTURAL-HERITAGE SITES ...
CHARACTERIZATION OF THE SUBSTRATES FROM
TWO CULTURAL-HERITAGE SITES AND A
PREPARATION OF MODEL SUBSTRATES
KARAKTERIZACIJA GRADBENIH MATERIALOV IZ DVEH
OBJEKTOV KULTURNE DEDI[^INE IN PRIPRAVA MODELNIH
MATERIALOV
Sabina Kramar1, Vilma Ducman1, Snezana Vucetic2, Edo Velkavrh3,
Miroslava Radeka4, Jonjaua Ranogajec2
1Slovenian National Building and Civil Engineering Institute, Dimi~eva 12, 1000 Ljubljana, Slovenia
2University of Novi Sad, Faculty of Technology, Bulevar Cara Lazara 1, 21000 Novi Sad, Serbia
3Saning International, d. o. o., Kranj, Ulica Mirka Vadnova 1, 4000 Kranj, Slovenia
4University of Novi Sad, Faculty of Technical Sciences, Trg Dositeja Obradovica 6, 21000 Novi Sad, Serbia
sabina.kramar@zag.si
Prejem rokopisa – received: 2013-09-11; sprejem za objavo – accepted for publication: 2013-10-09
In this study the microstructural characteristics of the materials selected from two cultural-heritage sites (the Dornava Manor,
Slovenia, and the Ba~ Fortress, Serbia) and of the model samples – the control and aged ones – were investigated. The samples
were characterized by means of mercury intrusion porosimetry (MIP) and total-specific-area analysis (BET). A good agreement
was achieved between the samples of the brick and mortar from the Ba~ Fortress, or the samples of the natural stone and render
from the Dornava Manor, and the corresponding model samples.
Keywords: cultural heritage, microstructure, brick, natural stone, mortar, frost resistance
Prispevek obravnava mikrostrukturne lastnosti izbranih materialov iz dveh objektov kulturne dedi{~ine (dvorec Dornava v
Sloveniji in trdnjava Ba~ v Srbiji) ter modelnih materialov – kontrolnih in staranih. Vzorce smo preiskali z `ivosrebrno
porozimetrijo in plinsko sorpcijo. Med vzorci iz objektov kulturne dedi{~ine in modelnimi vzorci smo dosegli dobro ujemanje,
tako v primeru opeke in malte s trdnjave Ba~ kot naravnega kamna in ometa z dvorca Dornava.
Klju~ne besede: kulturna dedi{~ina, mikrostruktura, opeka, naravni kamen, malta, zmrzlinska obstojnost
1 INTRODUCTION
Within the scope of the 7th FP HEROMAT project
(Protection of cultural heritage objects with multifunc-
tional advanced materials ENV-NMP.2011.3.2.1-1
NMP) the main goals is to develop the consolidants and
multifunctional photocatalytic coatings for cultural-
heritage sites in order to conserve degraded materials
and to prevent their further degradation. To meet the set
goals and develop the basis for further applications, the
following steps were implemented: (i) an analysis of the
microstructure of the substrates from the two studied
facilities (the Dornava Manor, Slovenia, and the Ba~ For-
tress, Serbia); (ii) the preparation of a model substrate,
simulating the substrates from such sites for the prelimi-
nary testing of the newly developed products; (iii) an
analysis of the microstructure of the model substrate
before and after the ageing procedure, and a comparison
of this substrate with the substrates from the two selected
sites.
The Dornava Manor Complex, with its accompany-
ing park, is one of the most important monuments of the
late Baroque period in Slovenia. Because of a long expo-
sure to a strong degradation and inappropriate restoration
procedures, the stone elements of the exterior, i.e., the
statues and ornaments of the garden, the balustrades and
the main building, show only a faint picture of the past.
The Ba~ Fortress, Serbia (currently included in the
UNESCO Tentative List) is the best preserved brick-built
medieval fortress in Vojvodina. This, first a Hungarian
and later on an Ottoman stronghold, has been in ruins for
centuries as a result of various physical, chemical and
biological degradation processes. Its physical integrity
has been substantially lost, but the preserved elements
indicate the sophisticated architecture of the fortification
school of the high Gothic style, with the elements of the
early Italian Renaissance.
Among the most active causes for the weathering of
the historical monuments exposed to the outdoor envi-
ronment is frost. A common degradation phenomenon of
the porous materials such as mortars, stone and brick is a
loss of cohesion of the binder-aggregate/grain system.
This process is usually followed by a deposition and for-
mation of new products, a loss of mechanical strength, a
loss of the material from the surface, and chromatic
alterations.1 The cohesion recovery between these parti-
cles, released by the binder loss, can be achieved through
an application of consolidants.2
Porosity and pore-size distribution can have a major
effect on the durability of porous materials, influencing
Materiali in tehnologije / Materials and technology 48 (2014) 4, 505–508 505
UDK 691:620.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)505(2014)
the effectiveness of a consolidation, the depth of a conso-
lidant, etc.3,4 No consolidation action can be reasonably
justified without having access (directly or by deduction)
to the information on all these aspects.5
Taking into consideration the fact that porosity is the
most important factor with regard to the absorption and
fluid transport in porous materials, the present study fo-
cuses on the characterization of the porosity and pore-
size distribution as well as determining the specific
surface area. In this study the microstructural characte-
ristics of the selected materials from the Dornava Manor
and the Ba~ Fortress are determined. In order to assess
the changes that occur in the microstructure after a
freeze/thaw cycles, model substrates were prepared. The
samples were characterized with mercury intrusion poro-
simetry (MIP) and total-specific-area analysis (BET).
2 EXPERIMENTAL WORK
2.1 Materials and sample preparation
Four substrates were selected among various mate-
rials from the two historical sites: (i) sandstone (desig-
nated as Stone Dornava) and (ii) a render sample from
the Dornava Manor (designated as Render Dornava); (iii)
lime mortar with a carbonate aggregate (designated as
Mortar Ba~) and (iv) a brick from the Ba~ Fortress
(designated as Brick Ba~).
Based on the mineralogical and petrographic analy-
ses, a quartzitic-micaceous sandstone from the Dornava
Manor was recognised as the Vundu{ki sandstone out-
cropping in the eastern part of Slovenia. As a model
stone, a bulk sample of the Vundu{ek sandstone (from
the Middle-Miocene geological period) was taken from
the Jelovice quarry, located near the town of Maj{perk,
in the eastern part of Slovenia. Samples were cut having
a dimension of 5 cm × 5 cm × 1 cm and designated as
Model Stone or Model Stone F, where F means that the
sample was exposed to freezing and thawing.
The model render was prepared on the basis of a
petrographic analysis of render sample DOR 137 from
the Dornava Manor. A silicate gravel fraction of 0–4 mm
was obtained from the Ho~e quarry (Slovenia) and sieved
in order to get a grading curve for the aggregate similar
to that of the analyzed render. In order to weaken the
mechanical properties of the cement part of the model
samples, a Portland cement (CEM I 42.4, Lafarge) mixed
with tuff (Ecotrass, Saning, d. o. o.) was used. The tradi-
tionally prepared slaked lime from the local producers
was also used. The substrates were prepared by combin-
ing together coarse and fine renders in order to achieve
the appropriate thickness. On an approximately 3.5 cm
thick coarse layer, a 3–4 mm thick layer of fine render
was applied. Steel moulds with the dimensions of 16 cm
× 4 cm × 4 cm and 15 cm × 15 cm × 4 cm were used for
the preparation of the models that were then cut out to
the dimensions required for the testing. The samples
were cured at 100 % relative humidity at room tempe-
rature for a minimum of 14 d before the cutting was
performed.
The model bricks were prepared by hand in the tradi-
tional way, using an old crushed brick and the clay from
a pit situated near the Ba~ Fortress. Based on the mine-
ralogy and chemical composition of the examined
historical-brick samples, the optimum formulation for
the raw-material mixture and firing temperature was
determined. The clay material was based mainly on
quartz and carbonates such as dolomite and calcite. The
raw material was hand-moulded in the laboratory. The
moulds were filled and then air-dried for a period of two
weeks. When an adequate strength was obtained, the
model bricks were fired in a laboratory kiln (T max of
980 °C). The prepared model-brick samples were then
exposed to a laboratory degradation, followed by a
characterization procedure. Samples with the dimensions
of 5 cm × 5 cm × 1 cm were prepared.
For the preparation of the model mortar, the follow-
ing components were used: lime putty, micronized
quicklime, sand and water (the hot-lime-mix method).
The selected components provided a mineralogical and
chemical composition that was as close as possible to the
characterized historical mortars. All the mortars were
prepared in the traditional way, in a wooden dish. The
prepared mortars were moulded in specially designed
moulds. The interior surfaces of the moulds were coated
with TFE-fluorcarbon (Teflon) that did not affect the
setting of the mortar, also preventing any damage to the
mortar surfaces. During the first weeks, carbonation was
achieved in a moist room and after that all the samples (5
cm × 5 cm × 1 cm) were placed in a chamber with a
constant flow of CO2 (accelerated carbonation).
2.2 Methods
The freeze/thaw cycling was performed according to
standard EN 12371:2010. All the samples were exposed
to 50 cycles of freezing and thawing, where one cycle
consists of (1) a sample being in water for 5.5 h and at
15 °C, (2) sucking the water from the chamber so that
the temperature reaches –4 °C within 2 h, (3) cooling
down over the next 4 h to –10 °C, and (4) pouring the
water again into the chamber.
The porosity and pore-size distribution of the inve-
stigated samples were determined by means of mercury
intrusion porosimetry (MIP). Small blocks, approxi-
mately 1 cm3 in size, were dried in an oven for 24 h at
105 °C, and then analysed using a Micromeritics Auto-
pore IV 9500 Series pore-size analyzer. The samples
were analysed in the range of 0–414 MPa, using solid
penetrometers.
On the basis of the BET (Brunauer-Emmet-Teller)
method, the total specific area of a sample surface was
determined with nitrogen adsorption at 77 K within a
relative pressure range of 0.05 to 0.3 using a Micro-
meritics ASAP-2020 analyzer. Prior to the performance
of these measurements, the samples (2 to 3 small blocks
S. KRAMAR et al.: CHARACTERIZATION OF THE SUBSTRATES FROM TWO CULTURAL-HERITAGE SITES ...
506 Materiali in tehnologije / Materials and technology 48 (2014) 4, 505–508
with a weight of approximately 1 g) were being degassed
for at least 3 h.
3 RESULTS AND DISCUSSION
The results of the porosity and BET specific-surface
tests are presented, for the all prepared samples, in Table
1, from which it can be seen that there is a very good
agreement between the samples from the historical sites
and the model samples obtained, in the case of the brick
and mortars, from the Ba~ Fortress, and a fairly good
agreement with the stone and render from the Dornava
Manor. For example, the overall porosity of the Brick
Ba~ sample was 42.0 %, whereas it was 45.6 % for the
model-brick sample. In the case of the Mortar Ba~ sam-
ple it was 40.3 % compared to 32.2 % for the model-
mortar sample. The porosity of the Stone Dornava sam-
ple was found to amount to 9.8 % compared to 13.4 %
for the model-stone sample, whereas the porosity of the
original render from Dornava reached 26.7 % and that of
the model-render sample was 38.6 %.
After an exposure to the freeze/thaw cycling, only a
slight increase in the porosity of the model-stone sample
was observed. On the other hand, the sample’s average
pore diameter increased significantly, shifting from 0.04
μm to 0.26 μm (Figure 1). The pore-size distribution was
unimodal, with an intrusion peak at 2 μm. The BET
surface-area values also increased. The pressure caused
by freezing range from 14 MPa to 138 MPa, with a
decrease in the temperature of between –1.1 °C and
–12.5 °C.6 During the ageing test performed within the
scope of this study, the temperature fell by up to –10 °C.
The temperature range considered critical for a deteriora-
tion of natural stone is from about –4 °C to –15 °C.6 A
stone with a higher quantity of smaller pores is more
prone to frost deterioration as well as salt crystallisation,
although stone damage is more specifically influenced
by nanopores in the case of salt crystallisation, and by
micropores in the case of frost damage.7
The porosity of the model-render sample reduced
significantly after the freeze/thaw cycling, whereas the
change in the average pore size was not significant (0.04
μm and 0.06 μm). The main intrusion peak after the
freeze/thaw cycling shifted from 1 μm to 0.15 μm (Fig-
ure 2). A significant increase in the BET surface area of
the render was observed after the freeze/thaw cycling.
Both of these phenomena can be ascribed to the disso-
lution/precipitation processes that occur due to the water
exposure in the freezing/thawing test.
The overall porosity of the model-brick sample
decreased slightly after the freeze/thaw cycling tests,
which can be attributed to the natural re-carbonation of
the calcium hydroxide inside the porous media, thus also
contributing to an improvement in the durability.8 Addi-
tionally, the average pore diameter also decreased (from
0.57 μm to 0.21 μm). However, a certain increase was
observed in the case of bigger pores – there are three
peaks above 10 μm, representing either pores or cracks.
This could mean that the brick might be susceptible to
S. KRAMAR et al.: CHARACTERIZATION OF THE SUBSTRATES FROM TWO CULTURAL-HERITAGE SITES ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 505–508 507
Figure 2: Pore-size distributions for the samples of render
Slika 2: Porazdelitev velikosti por v vzorcih ometa
Figure 1: Pore-size distributions for the samples of natural stone
Slika 1: Porazdelitev velikosti por v vzorcih naravnega kamna
Table 1: Comparison of the porosity and specific surface (the average of 2 measurements) of the samples from Dornava Manor and Ba~ Fortress
with the corresponding values for the model substrates before and after the ageing (F – exposure to freezing/thawing)
Tabela 1: Primerjava poroznosti in specifi~ne povr{ine (povpre~je 2 meritev) vzorcev iz dvorca Dornava in vzorcev iz trdnjave Ba~ z ustreznimi
vrednostmi modelnih podlag pred staranjem in po njem (F – izpostavitev zmrzovanju in odtajanju)
Brick Ba~ Model brick Model brick F Mortar Ba~ Model mortar Model mortar F
Porosity (%) 42.0 ± 4.4 45.6 ± 0.4 43.5 ± 0.8 40.3 ± 0.8 32.2 ± 4.1 29.6 ± 0.6
BET (m2/g) 5.4 ± 0.9 2.03 ± 0.7 5.0 ± 1.4 2.8 ± 0.3 2.3 ± 0.2 1.8 ± 0.2
Stone Dornava Model stone Model stone F Render Dornava Model render Model render F
Porosity (%) 9.8 ± 0.2 13.4 ± 1.2 17.1 ± 2.8 26.7 ± 1.6 38.6 ± 1.0 27.8 ± 0.2
BET (m2/g) 1.9 ± 0.3 2.3 ± 0.5 3.5 ± 0.1 10.3 ± 0.5 4.8 ± 0.4 9.3 ± 0.3
freezing. The values of the BET surface area increased
significantly after the ageing. The pore-size distribution
of Brick Ba~ and model-brick samples was bimodal, but
after the freeze/thaw cycling it became unimodal (Figure
3).
The values of the porosity and the average pore
diameter for the model-mortar sample decreased after
the freeze/thaw cycling, most probably due to the disso-
lution/reprecipitation processes inside the pores. The
pore-size distribution of the control sample was bimodal
– this distribution was observed for a significant number
of historical mortars9, becoming unimodal after the
ageing10 (Figure 4). The values of the BET surface area
increased significantly after the freeze/thaw cycling,
indicating a deleterious effect of freezing on the mortar.
4 CONCLUSIONS
Samples of the selected substrates from the two
monuments and model substrates – the control and aged
ones, were investigated by means of mercury intrusion
porosimetry (the total porosity, the average pore dia-
meter, the pore-size distribution) and gas sorption (the
BET surface area). A good agreement between the sam-
ples from the historical sites and the model samples was
obtained for the brick and mortars from the Ba~ Fortress
as well as for the natural stone and render from the
Dornava Manor.
On the basis of these results, the selected model
substrates will be used with newly developed materials
in order to test and otherwise scrutinize them before
applying them to the valuable substrates belonging to
cultural-heritage sites.
Acknowledgements
This work was financially supported by the European
Commission under the 7th FP HEROMAT project
(Protection of cultural heritage objects with multifunc-
tional advanced materials, ENV-NMP.2011.3.2.1-1
NMP).
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1 L. Toniolo, A. Paradisi, S. Goidanich, G. Pennati, Mechanical beha-
viour of lime based mortars after surface consolidation, Construction
and Building Materials, 25 (2010) 4, 1553–1559
2 G. Borsoi, M. Tavares, R. Veiga, A. Santos Silva, Microstructural
Characterization of Consolidant Products for Historical Renders: An
Innovative Nanostructured Lime Dispersion and a More Traditional
Ethyl Silicate Limewater Solution, Microsc. Microanal., 18 (2012) 5,
1181–1190
3 S. P. Gupta, Consolidation of historic porous stone by impregnation
with silica acid-esters, World Journal of Science and Technology, 1
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Valle, M. Alvarez de Buergo, R. Fort, Influence in the porosity and
relative humidity on consolidation of dolostone with calcium
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tive techniques, Materials Characterization, 61 (2010), 168–184
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problem with few practical solutions, Historical Constructions, P. B.
Lourenço, P. Roca (Eds.), Guimarães, (2001), 3–14
6 A. Goudie, H. Viles, Salt weathering hazards, J. Wiley and Sons,
Chichester 1997, 235
7 J. E. Lindquist, K. Malaga, B. Middendorf, M. Savukoski, M. P.
Pétursson, Frost resistance of natural stone, the importance of micro-
and nano-porosity, Geological Survey of Sweden, External research
project report, published online at: http://www.sgu.se/dokument/
fou_extern/Lindqvist-etal_2007.pdf
8 P. Lopez-Arce, D. Benavente, J. García-Guinea, Durability improve-
ment of ancient bricks by cementation of porous media, J. Am.
Ceram. Soc., 88 (2005) 9, 2664–2672
9 A. Klisinska-Kopacz, R. Ti{lova, G. Adamski, R. Koz³ovksi, Control
of the porosity structure to produce Roman cement mortars compa-
tible with the historic substrate, Proceedings of the 11th International
Congress on Deterioration and Conservation of Stone, Torun, Poland,
2008, 931–938
10 A. Izaguirre, J. Lanas, J. I. Álvarez, Ageing of lime mortars with ad-
mixtures: Durability and strength assessment, Cement and Concrete
Research, 40 (2010) 7, 1081–1095
S. KRAMAR et al.: CHARACTERIZATION OF THE SUBSTRATES FROM TWO CULTURAL-HERITAGE SITES ...
508 Materiali in tehnologije / Materials and technology 48 (2014) 4, 505–508
Figure 4: Pore-size distributions for the samples of mortar
Slika 4: Porazdelitev velikosti por v vzorcih malt
Figure 3: Pore-size distributions for the samples of brick
Slika 3: Porazdelitev velikosti por v vzorcih opeke
L. ZAVR[NIK et al.: LONG-TERM DURABILITY PROPERTIES OF POZZOLANIC CEMENT MORTARS
LONG-TERM DURABILITY PROPERTIES OF
POZZOLANIC CEMENT MORTARS
DOLGORO^NA OBSTOJNOST PUCOLANSKE CEMENTNE MALTE
Lina Zavr{nik, Jerneja Strupi [uput, Sabina Kramar
Slovenian National Building and Civil Engineering Institute, Dimi~eva ulica 12, 1000 Ljubljana, Slovenia
lina.zavrsnik@zag.si
Prejem rokopisa – received: 2013-09-23; sprejem za objavo – accepted for publication: 2013-10-30
Within the scope of this study, the durability properties of mortar prepared with pozzolanic cement were examined. The
physical-mechanical properties of the mortars were determined after their exposure to various aggressive media (a sodium
sulphate water solution, a magnesium sulphate water solution, seawater, and a water solution of NO32–, SO42–, and NH4–) and
freeze/thaw cycling. Additionally, the depth of carbonation of the mortars was determined. The results showed that the mortar
containing pozzolanic cement was sulphate resistant, but when it was exposed to a solution of NO32–, SO42–, and NH4– its
strength was significantly reduced. On the other hand, the mortar was found to be sensitive to freezing/thawing, as its
compressive strength decreased by 50 %, and its flexural strength by as much as 90 %. Due, especially, to its lack of resistance
to freezing/thawing, the use of the investigated pozzolanic cement mortar is limited to indoor applications.
Keywords: pozzolanic cement, durability, sulphate attack, frost resistance, carbonation
V prispevku obravnavamo trajnostne lastnosti malte, pripravljene iz pucolanskega cementa. Dolo~ili smo fizikalno-mehanske
lastnosti malte po izpostavitvi razli~nim agresivnim medijem (vodni raztopini natrijevega sulfata, vodni raztopini magnezijevega
sulfata, morski vodi, vodni raztopina ionov NO32–, SO42–, NH4–) in cikli~nemu zmrzovanju in odtaljevanju. Poleg tega smo
dolo~ili tudi globino karbonatizacije. Na podlagi rezultatov ugotavljamo, da je malta s pucolanskim cementom odporna proti
sulfatom, medtem ko se njena trdnost po izpostavitvi raztopini ionov NO32–, SO42– in NH4– mo~no zmanj{a. Prav tako preiskana
malta nima zmrzlinske obstojnosti, saj se njena tla~na trdnost zmanj{a za 50 %, upogibna pa za 90 %. Na osnovi rezultatov
ugotavljamo, da je malta zaradi slabe zmrzlinske obstojnosti uporabna le za notranje povr{ine.
Klju~ne besede: pucolanski cement, obstojnost, sulfatna korozija, obstojnost proti zmrzali, karbonatizacija
1 INTRODUCTION
Over the past decade there has been a clearly percep-
tible trend towards the replacement of Portland cement,
by up to 80 %, with various mineral additives, such as
ground granulated blast furnace slag, natural (volcanic
ash) and synthetic pozzolans (fly ashes, microsilica, rice
husk ash), and ground limestone.1–4 An important reason
for the use of blended or pozzolanic cements lies in their
good resistance to attack by chemical agents, especially
seawater, and by sulphate-rich water.5–7 In addition, the
replacement of Portland cements with pozzolans has
environmental benefits, since in this way waste materials
can be used, thus contributing to the reduced use of
naturally occurring raw materials, and less energy is
consumed in their production compared with the
production of cement clinker, reducing greenhouse-gas
emissions and the overall CO2 footprint.
The most widely used cement-replacement material
is coal combustion fly ash. In the EU, the replacement of
up to 33 % of the Portland cement with fly ash is per-
mitted for a range of cement types, whereas in the case
of pozzolanic cement a limit of 55 % has been specified
(EN 197-1:2011).8 The performance of concrete with
added fly ash is frequently better than that of concrete
mixed with Portland cement only. Although the pozzo-
lanic reaction is slow, and can cause a significant de-
crease in early-age strengths, the use of fly ash in cement
results in a generally improved workability, a higher
late-age strength and increased durability, as well as im-
proved sulphate resistance, and reductions in the heat of
hydration and in the risk of the occurrence of an alkali
silica reaction and efflorescence.4,7,9 There are some situ-
ations, however, where the performance of concrete with
added fly ash may not be so good, e.g., when such con-
crete is subjected to freeze-thaw actions. In this case
concerns have been raised regarding the concrete’s dura-
bility, especially when the fly ash is used at higher con-
tent levels.10 Two main classes of fly ash can be defined:
siliceous fly ash (low calcium fly ash; < 10 % CaO),
which has been widely used as a replacement for cement,
and calcareous fly ash (high calcium fly ash; > 10 %
CaO). Calcareous fly ashes show a more rapid strength
gain during early ages than in the case of concretes made
with siliceous fly ashes.11 This is because calcareous fly
ashes usually exhibit a higher rate of reaction during
early ages than siliceous fly ashes. Since a negative
cement performance has been frequently connected with
the presence of lime and sulphates, the limits of the
existing standards are very strict, so that most calcareous
fly ashes are rejected as unsuitable.11 However, the large
amounts of calcareous fly ashes that are produced have
resulted in a pressing need for their utilization, especially
if sustainability issues are to be adequately resolved in
the construction industry.
Materiali in tehnologije / Materials and technology 48 (2014) 4, 509–513 509
UDK 691.54 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)509(2014)
In the present investigation the durability properties
of mortars prepared with pozzolanic cement have been
investigated. The pozzolanic cement used in the study
was a mix of calcareous fly ash and Portland cement
clinker designated as CEM IV/B-W 32.5 N according to
EN 197-1:20118. The physical-mechanical properties of
mortars made using this type of cement were determined
after the latter had been exposed to different aggressive
media, as well as to freeze/thaw cycling.
2 EXPERIMENTAL
2.1 Materials and methods
A pozzolanic cement CEM IV/B-W 32.5 N (Lafarge
Cement, d. o. o., Trbovlje, Slovenia) and CEN standard
sand were used for the preparation of mortars. A cement:
aggregate:water ratio of 1 : 3 : 0.6 was used, except for
the investigation of the resistance to freezing/thawing
and carbonation. The properties of the samples of fresh
mortar (consistence EN 1015-3:200112, bulk density EN
1015-6:199913, air content EN 1015-7:199914) for the
investigation of the resistance to freezing/ thawing and
carbonation are given in Table 1.
In order to evaluate the sulphate resistance of the
pozzolanic cement mortar, samples having dimensions of
10 mm × 10 mm × 60 mm were cured for 21 d in deioni-
zed water. After this, some of the samples were exposed
to a 4.4 % Na2SO4 water solution or a 3.73 % MgSO4
water solution, whereas the remaining samples were left
in the deionized solution. In order to compare the influ-
ence of the different curing media, another set of sam-
ples was cured in saturated lime water for 21 d, after
which they were exposed to a 4.4 % Na2SO4 water
solution, whereas the remaining samples were left in the
saturated lime water. Compressive and flexural strengths
according to EN 196-1:200515 were determined after (14,
28 and 56) d. The resistance of the investigated mortars
to sulphate attack was investigated and evaluated using
the Koch-Steinegger test.16
In order to assess the influence of various aggressive
media on the long-term durability of pozzolanic cement
mortar, samples having dimensions of 40 mm × 40 mm ×
160 mm were cured for 28 d in saturated lime water.
After this, some of the samples were exposed to a 4.4 %
Na2SO4 water solution, as well as to seawater, and to a
(NO32–, SO42–, NH4–) water solution, whereas the remain-
ing samples were left in the saturated lime water. The
compressive and flexural strengths of the mortar samples
were determined according to EN 196-1:200515, after 1
and 2 years.
In order to investigate the resistance to freezing/
thawing of the pozzolanic cement mortar (in the absence
and presence of de-icing salt) an air-entraining admixture
was added to the mortar (Cementol Eta S, and Cementol
Eta EM, produced by TKK Srpenica). The amount of
air-entraining admixture used in the preparation of the
mortar samples was determined by the consistence of the
fresh mortar (EN 1015-3:2001)12, by keeping a flow
value (175 ± 5) mm (Table 1). For the investigation of
the resistance to freezing/thawing in the absence of de-
icing salt, mortar samples having dimensions of 40 mm
× 40 mm × 160 mm were cured for 28 d in water at (20 ± 1)
°C. After this, some of the samples were exposed to
freeze/thaw cycling, whereas the remaining samples
were left in the water at (20 ± 1) °C. Freeze/thaw cycling
was performed according to the standard SIST 1026:
2008.17 All the samples were exposed to 50 cycles of
freezing/thawing, where one cycle consisted of 4 h at
(–20 ± 2) °C and 4 h at (20 ± 2) °C.
The flexural and compressive strengths of the pre-
pared mortar samples were determined according to EN
196-1:200515 after 50 cycles. In the case of freeze/thaw
cycling in the presence of the de-icing salt solution (3 %
NaCl), samples having dimensions of 200 mm × 200 mm
× 50 mm were cured for 7 d at (20 ± 2) °C at a RH of (65
± 5) %, and their resistance to freezing/thawing was
assessed according to SIST 1026:2008.17
The carbonation of the mortar samples was assessed
on mortar prisms having dimensions of 40 mm × 40 mm
× 160 mm, according to EN 13295:2004.18 The samples
were cured for 24 h in a mould, wrapped in film. After
de-moulding, the samples were cured for 48 h, wrapped
in film, and then for 25 d at (21 ± 2) °C and at a RH of
(60 ± 10) %. After this, they were placed in a chamber
containing CO2 1 %, at a RH of (60 ± 10) %. The depths
of the carbonation were determined by spraying a phe-
nolphthalein indicator solution onto the freshly broken
surface. Another set of samples was exposed to the natu-
ral environment, i.e., to outdoor conditions in a sheltered
area. The depth of carbonation was determined after (28,
56, 90, 180 and 270) d, using the phenolphthalein indi-
cator solution.
3 RESULTS AND DISCUSSION
3.1 Resistance to the investigated aggressive media
The compressive and flexural strengths of the mortar
specimens that had been immersed in the sodium sul-
phate solution and the magnesium sulphate solution were
L. ZAVR[NIK et al.: LONG-TERM DURABILITY PROPERTIES OF POZZOLANIC CEMENT MORTARS
510 Materiali in tehnologije / Materials and technology 48 (2014) 4, 509–513
Table 1: Properties of the samples of fresh mortar for resistance to freezing/thawing and carbonation
Tabela 1: Odpornost vzorcev sve`e malte proti zmrzovanju/odtaljevanju in karbonatizaciji
Additive – cement
(%) Water/cement ratio
Consistence (mm),
EN 1015-3:200112
Bulk density (kg/m3),
EN 1015-6:199913
Air content (%),
EN 1015-7:199914
– 0.5 174 2180 3.6
0.15 (Cementol Eta S) 0.5 175 2100 4
0.6 (Cementol Eta EM) 0.47 171 2105 7.4
determined and compared with those of the samples that
had been immersed in deionised or saturated lime water.
A difference in the initial strengths was observed, which
depended upon which solution, i.e., deionised water or
saturated lime water, had been used to cure the specimens
prior to the test. The lower initial strengths obtained in
the case of the deionised water suggest that the latter
behaves as an aggressive solution, if compared to the
saturated lime water.
According to the Koch-Steinegger test,16 the criterion
that can be used to classify a material as resistant or
durable in a specific aggressive medium is that the
corrosion index (i.e., the relative strength of the
aggressive-solution-stored samples to the water-stored
samples) must be higher than 0.7. As can be seen from
Figures 1 and 2, the corrosion index for the mortar
samples that were exposed to the investigated aggressive
media was always above 0.7, in the case of both flexural
and compressive strength. The corrosion index is slightly
higher in the case of the magnesium sulphate solution.
These results suggest that the pozzolanic cement mortar
could be classified as durable or resistant to a sodium
sulphate solution, as well as to a magnesium sulphate
solution. This can be mainly attributed to the pozzolanic
reaction, where the reactive siliceous and aluminous
phases in the fly ash react with portlandite to form new
C-S-H- or C-AS-H-type phases, provoking a refinement
in the pore structure, and thus impeding the penetration
and movement of potentially aggressive ions.6,19
The flexural and compressive strengths of the mortar
samples, measured after long-term exposure to the
investigated aggressive media, are shown in Figures 3
and 4. The results indicate that the mortar containing
pozzolanic cement continued to be sulphate resistant
after 2 years of exposure, but when it was exposed to
seawater or a solution of NO32–, SO42–, and NH4– its
strength was significantly reduced. This could be due to
L. ZAVR[NIK et al.: LONG-TERM DURABILITY PROPERTIES OF POZZOLANIC CEMENT MORTARS
Materiali in tehnologije / Materials and technology 48 (2014) 4, 509–513 511
Figure 4: Compressive strength of the investigated mortar samples
after long-term exposure to the investigated aggressive media
Slika 4: Tla~na trdnost preiskovanih malt po dolgotrajni izpostavitvi
preiskovanim agresivnim medijem
Figure 1: Corrosion index (flexural strength) for the mortar samples
exposed to various solutions (DW – deionised water, SLW – saturated
lime water)
Slika 1: Korozijski indeks (upogibna trdnost) za vzorce malte, izpo-
stavljene razli~nim raztopinam (DW – deionizirana voda, SLW –
nasi~ena apnena voda)
Figure 3: Flexural strength of the investigated mortar samples after
long-term exposure to the investigated aggressive media
Slika 3: Upogibna trdnost preiskovanih malt po dolgotrajni izposta-
vitvi preiskovanim agresivnim medijem
Figure 2: Corrosion index (compressive strength) for the mortar
samples exposed to various solutions (DW – deionised water, SLW –
saturated lime water)
Slika 2: Korozijski indeks (tla~na trdnost) za vzorce malte, izpostav-
ljene razli~nim raztopinam (DW – deionizirana voda, SLW – nasi~ena
apnena voda)
the removal of calcium hydroxide, the hardened cement
paste will be decalcified, causing the pH value to
decrease.20
3.2 Resistance to freezing/thawing
The results showed (Table 2) that, after the freeze/
thaw cycling, the compressive strength of the investi-
gated mortar samples decreased by more than 50 %, and
their flexural strength by as much as 90 %. The observed
decrease in strengths was higher in the case of the mortar
with a higher air content.
In the presence of de-icing salt the loss of material
exceeded the requirement according to the standard SIST
1026:200817 (0.35 mg/mm2) already after 10 cycles of
freezing/thawing. It was, however, higher in the case of
the mortar with a higher air content, as the loss of mate-
rial was 2.32 mg/mm2 and, in the case of the mortar with
a lower air content, 1.58 mg/mm2. Thus, the investigated
mortar is non-resistant to frost action, despite the addi-
tion of an air-entraining admixture. In general, the
freeze-thaw durability of fly ash concrete, when a fly ash
content of more that 20 % is used, is similar to or worse
than that of Portland cement concretes, which could be
attributed to compatibility problems between the fly ash
and the air entraining agents.10
Table 2: Strengths of the investigated mortar samples after freeze/
thaw cycling
Tabela 2: Trdnost preiskovanih malt po cikli~nem zamrzovanju/odta-
ljevanju
Additive – sample
(%)
Flexural strength
(N/mm2)
Compressive
strength
(N/mm2)
cured in
water
freeze/
thaw
cycling
cured in
water
freeze
/thaw
cycling
0.15
(Cementol Eta S) 7.7 1.3 52.1 31.2
0.6
(Cementol Eta EM) 9.3 0.5 57.3 23.1
3.3 Carbonation
The depth of carbonation of the mortar samples
exposed to the accelerated carbonation was 10.2 mm
after 56 d, and 15.4 mm after 90 d. On the other hand,
the mortar samples that were exposed to the natural
environment showed a reduced depth of carbonation,
which amounted to 3.1 mm after 56 d, 3.9 mm after 90 d,
and 6.1 mm after 270 d.
It has been reported that the carbonation of concrete
increases significantly with increased fly ash content,
i.e., when the fly ash content is higher than 30 %, as a
large amount of fly ash delays the hydration and the
increase of porosity.21 However, when the binder used in
the mortar samples contained just 10 % of fly ash, the
carbonation level was equal to that of the samples
containing Portland cement only, or was slightly higher.
On the other hand, when the concrete incorporated calca-
reous fly ash, carbonation was reduced in comparison
with the siliceous fly ash.21
4 CONCLUSIONS
The investigated pozzolanic cement CEM IV/B-W
32.5 N is resistant to the selected sulphate solution,
whereas both of the selected (NO32–, SO42–, NH4–)
solution and seawater had a negative effect, since the
strengths of the tested mortar samples reduced over a
period 1–2 years. However, there was a difference in the
initial strengths, depending upon which solution, i.e.,
deionised water or saturated lime water, was used to cure
the samples prior to testing. Deionised water behaves as
an aggressive solution when compared to saturated lime
water. On the other hand, the mortar was found to be
sensitive to freezing/thawing in both the presence and the
absence of de-icing salt, as its compressive strength
decreased by 50 %, and its flexural strength by as much
as 90 %. When exposed to a selected natural environ-
ment the mortar samples showed a reduced depth of
carbonation with respect to the accelerated carbonation
test.
Due, especially, to its lack of resistance to freezing/
thawing, the use of the investigated pozzolanic cement
mortar is limited to indoor applications.
Acknowledgements
This study was financially supported by the Metrolo-
gy Institute of the Republic of Slovenia.
5 REFERENCES
1 K. Sobolev, Mechano-chemical modification of cement with high
volumes of blast furnace slag, Cement and Concrete Composites, 27
(2005) 7–8, 848–853
2 B. Lothenbach, G. Le Saout, E. Gallucci, K. Scrivener, Influence of
limestone on the hydration of Portland cements, Cement and Con-
crete Research, 38 (2008), 848–860
3 V. Ramasamy, Compressive strength and durability properties of rice
husk ash concrete, KSCE Journal of Civil Engineering, 16 (2012) 1,
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high volume fly ash cement pastes and mortars in aggressive solu-
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zolanic cements in relation to strength, Cement and Concrete
Research, 12 (1982), 726–734
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the mortar made from pozzolan cement with zeolite, Journal of
Thermal Analyses and Calorimetry, 94 (2008) 1, 7–14
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mity criteria for common cements, 2011
9 W. Ma, W. P. Brown, Hydrothermal reactions of fly ash with
Ca(OH)2 and CaSO4·2H20, Cement and Concrete Research, 27
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raw calcareous fly ash in mortar mixtures, Coal Combustion and
Gasification Products, 2 (2010), 9–14
12 EN 1015-3:2001, Methods of test for mortar for masonry - Part 3:
Determination of consistence of fresh mortar (by flow table), 2001
13 EN 1015-6:1999, Methods of test for mortar for masonry - Part 6:
Determination of bulk density of fresh mortar, 1999
14 EN 1015-7:1999, Methods of test for mortar for masonry - Part 7:
Determination of air content of fresh mortar, 1999
15 EN 196-1:2005, Methods of testing cement — Part 1: Determination
of strength, 2005
16 A. Koch, H. Steinegger, A rapid test for cements for their behaviour
under sulphated attack, Zement-Kalk-Gips, 7 (1960), 317–324
17 SIST 1026:2008, Concrete-part 1: Specification, performance, pro-
duction and conformity – Rules for the implementation of SIST EN
206-1 (Revised edition), 2008
18 EN 13295:2004, Products and systems for the protection and repair
of concrete structures - Test methods - Determination of resistance to
carbonation, 2004
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tious materials, Cement and Concrete Composites, 32 (2010) 9,
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Materiali in tehnologije / Materials and technology 48 (2014) 4, 509–513 513

B. BOUAROUR et al.: ESTIMATION OF THE EFFECTIVE BORON-DIFFUSION COEFFICIENT ...
ESTIMATION OF THE EFFECTIVE BORON-DIFFUSION
COEFFICIENT IN THE Fe2B LAYERS GROWN ON GRAY
CAST IRON
DOLO^ANJE EFEKTIVNEGA KOEFICIENTA DIFUZIJE BORA PRI
RASTI PLASTI Fe2B NA SIVEM LITEM @ELEZU
Boudjema Bouarour1, Mourad Keddam1, Omar Allaoui2
1Laboratoire de Technologie des Matériaux, Département de Sciences des Matériaux, Faculté de Génie Mécanique et Génie des Procédés,
USTHB, B.P N°32, 16111, El-Alia, Bab-Ezzouar, Algiers, Algeria
2Laboratoire de Génie des Procédés, Université Amar Tellidji de Laghouat, B.P. 37 G, 03000 Laghouat, Algeria
keddam@yahoo.fr
Prejem rokopisa – received: 2013-09-29; sprejem za objavo – accepted for publication: 2013-10-09
The present work evaluates the effective diffusion coefficient of boron in the Fe2B layers grown on a substrate from gray cast
iron. The boride layers were generated using powder-pack boriding in the temperature range of 1173–1273 K for (2, 4, 6 and 8)
h of treatment. First, the diffusion coefficient of boron in Fe2B (free of chemical stresses) was obtained through solving the
mass-balance equation at the (Fe2B/substrate) interface using the experimental values of parabolic-growth constants taken from
the literature.
Second, a simple equation was derived to evaluate the effective diffusion coefficient of boron in Fe2B by considering the effect
of chemical stresses. The estimated values of boron activation energies were 154.8 kJ/mol and 164.8 kJ/mol with and without
the presence of chemical stresses. Furthermore, the calculated values of effective diffusion coefficients were found to be
sensitive to the change in the boriding temperature and to the increase in the upper boron concentration in Fe2B.
Keywords: boriding, chemical stresses, kinetics, Fick’s law, effective diffusion coefficient, incubation time
To delo ocenjuje efektivni koeficient difuzije bora v plasteh Fe2B, zraslih na podlagi iz sivega litega `eleza. Boridne plasti so
nastale med boriranjem v prahu v {katli, v temperaturnem obmo~ju 1173–1273 K in trajanju (2, 4, 6 in 8) h. Najprej je bil
dolo~en koeficient difuzije bora v Fe2B (brez kemijskih napetosti) z re{itvijo ena~be za masno ravnote`je (Fe2B/podlaga) na
stiku z uporabo eksperimentalnih vrednosti za konstanto paraboli~ne rasti, dobljeno v literaturi.
Nato je bila izpeljana enostavna ena~ba za oceno efektivnega koeficienta difuzije bora v Fe2B z upo{tevanjem kemijskih
napetosti. Dolo~eni sta bili vrednosti aktivacijske energije bora 154,8 kJ/mol in 164,8 kJ/mol s kemijskimi napetostmi in brez
njih. Ugotovljeno je bilo tudi, da so izra~unane vrednosti efektivnega koeficienta difuzije ob~utljive za spremembo temperature
boriranja in za nara{~anje koncentracije bora v Fe2B.
Klju~ne besede: boriranje, kemijske napetosti, kinetika, Fickov zakon, efektivni koeficient difuzije, ~as inkubacije
1 INTRODUCTION
Boriding is a well-known thermochemical treatment
for creating boride layers with interesting properties by
saturating the surface layer of a material with boron. It is
used to improve the surface hardness, the frictional wear,
the fatigue endurance and the corrosion resistance of fer-
rous and non-ferrous alloys. It can be performed between
1123 K and 1323 K to form iron borides on the material
surface with the treatment times varying from 0.5 h to
10 h1. Different boriding methods exist in practice. How-
ever, the powder-pack boriding has the advantages of
simplicity, flexibility with respect to the powder compo-
sition, minimum equipment and cost-effectiveness.2–4
The boriding temperature, the time duration, the
boron potential of the medium and the chemical compo-
sition of the substrate are the key parameters controlling
the morphology, the growth kinetics and the microstruc-
tural nature of boride layers.5–7
The present model considers the effect of chemical
stresses when evaluating the effective diffusion coeffi-
cient of boron in the Fe2B layers grown on a gray-cast-
iron substrate. In this context, some recent reference
studies regarding the evaluation of the effective diffusion
coefficient of boron in Fe2B were reported for borided
AISI 1018 and AISI 4140 steels and borided Armco
iron.8–10
Certain reference works11,12 report on the chemical
stresses resulting from the composition gradient similar
to those caused by the thermal stresses in an isotropic
medium. Diffusion-induced stresses (or chemical stres-
ses) are then built up due to the composition inhomo-
geneity during mass transfer. Diffusion-induced stresses
also change the mechanical properties of metal systems
during mass transfer. Larché and Cahn13–15 investigated
the stresses arising from the inhomogeneities of mate-
rials. Consequently, chemical stresses enhance both the
diffusion coefficient and the concentration.16
The objective of this work was to evaluate the diffu-
sivity of boron in Fe2B (with and without the presence of
chemical stresses) in the temperature range of 1173–
1273 K by applying a kinetic model of a monolayer con-
figuration of Fe2B grown on gray cast iron.
Materiali in tehnologije / Materials and technology 48 (2014) 4, 515–520 515
UDK 546.271:532.72 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)515(2014)
2 MATHEMATICAL MODEL
A schematic non-linear concentration profile of
boron in the Fe2B layer is depicted in Figure 1 where the
boron potential allows the formation of Fe2B on the
surface of the material.
C up
Fe B2 denotes the upper boron concentration in Fe2B,
C low
Fe B2 (= 59.2 × 103 mol m–3) represents the lower boron
concentration in Fe2B and t0(T) is the boride incubation
time depending on the boriding temperature. Cads is
defined as the effective boron concentration17. Distance u
is the layer depth at the (Fe2B/substrate) interface as a
function of treatment time t. C0 is the boron solubility in
the matrix. The upper boron concentration in Fe2B is
between 59.2 × 103 mol m–3 and 60 × 103 mol m–3 since
this phase exhibits a narrow composition range as
pointed out by Massalski18.
When building a mathematical model of the problem,
the following assumptions are made:
• The growth kinetics is controlled by the boron diffu-
sion in the Fe2B layer
• The growth of the Fe2B boride layer is a consequence
of the boron diffusion perpendicular to the sample
surface
• The solid solution of boron in the matrix behaves
ideally (i.e., the activity coefficient of boron in the
solid solution is independent on the concentration)
• The Fe2B iron boride nucleates after a certain incu-
bation time
• The boride layer is thin in comparison to the sample
thickness
• Local equilibrium is held at the (Fe2B/substrate)
interface
• A planar morphology is assumed for the (Fe2B/sub-
strate) interface
• The volume change during the phase transformation
is neglected
• A uniform temperature is assumed throughout the
sample
• No effect of the alloying elements on the boron
diffusion is assumed
• The pressure effect on the boron effective diffusion
coefficient in Fe2B is ignored
The initial condition of the diffusion problem is
expressed with Equation (1):
C x CFe B2 ( , )0 0= (1)
The boundary conditions of the diffusion problem are
given with Equations (2) and (3):
[ ]{ }C x t t T t T CFe B upFe B2 2= = =0 00( ) , ( )
for C ads  59.2 × 10
3 mol m–3 (2)
{ }C x t t u t CFe B low
Fe B
2
2( ) ,= = =
for C ads  59.2 × 10
3 mol m–3 (3)
for 0  x  u.
The Fick’s second law of diffusion19 relating to the
change in the boron concentration throughout the Fe2B
layer with time t and location x(t) is:
∂
∂
∂
∂ 2
C x t
t
D
C x t
x
Fe B
B
Fe B Fe B2 2 2
( , ) ( , )
=
2
(4)
where DB
Fe B2 is the diffusion coefficient of boron in Fe2B
free of chemical stresses and CFe2B(x,t) is the boron
concentration. The boron concentration along the boride
layer, which is the solution of the Fick’s second law
(Equation (4)) in a semi-infinite medium, is consistent
with Figure 1. Its expression is given with Equation (5):
C x t C
C C
erf
u
D t
Fe B up
Fe B low
Fe B
up
Fe B
B
Fe B
2
2
2 2
2
( , ) = +
−
⋅
⎛
⎝ 2
⎜
⎜
⎜
⎜
⎞
⎠
⎟
⎟
⎟
⎟
⋅
⋅
erf
x
D t2 B
Fe B2
(5)
for 0  x  u.
The continuity equation at the (Fe2B/substrate) inter-
face is expressed with Equation (6):
w
x
t
D
C
xx u x u
Fe B B
Fe B Fe B
2
2 2
d
d
⎡
⎣⎢
⎤
⎦⎥ = −
⎡
⎣⎢
⎤
⎦⎥= =
∂
∂
(6)
with [ ]w C C CFe B upFe B lowFe B2 2 2= × + −05 0. ( ) .
Boride layer thickness u follows the parabolic-growth
law expressed with Equation (7), where k represents the
parabolic-growth constant at the (Fe2B/substrate) inter-
face:
u k t t T= − 0 ( ) (7)
where [t –t0(T)] is considered as the effective growth
time of the Fe2B layer9,10,20–22. According to Brakman et
al.23, boride incubation time t0(T) is decreased as the
temperature goes up. The (T) parameter, which takes
B. BOUAROUR et al.: ESTIMATION OF THE EFFECTIVE BORON-DIFFUSION COEFFICIENT ...
516 Materiali in tehnologije / Materials and technology 48 (2014) 4, 515–520
Figure 1: Schematic non-linear concentration profile of boron in the
Fe2B layer
Slika 1: Shemati~ni prikaz nelinearnega profila koncentracije bora
skozi plast Fe2B
into account the effect of the boride incubation time, is
given with Equation (8):
	( )
( )
T k
t T
t
= −1 0 (8)
The 	(T) parameter can be approached with a linear
relationship (Equation (9)), illustrated also in Figure 2
(with a correlation factor = 0.982):
	( ) ( . )T T= ⋅ × +−5 7 10 0 247224 (9)
By derivation of Equation (5) with respect to distance
x(t) and Equation (7) with respect to time t, Equation (6)
can be rewritten24 as follows:
[ ]05
2
0. ( )
(
C C C k
D C C
up
Fe B
low
Fe B
B
Fe B
low
Fe B
up
F
2 2
2 2
+ − =
= ⋅
−
π
e B
B
Fe B
B
Fe B
2
2
2
)
( )
exp
( )
( )
erf
k T
D
T k
D
T
	
	
	

2
4
2
⋅ −
⎛
⎝
⎜
⎞
⎠
⎟ ⋅
(10)
To evaluate the diffusion coefficient of boron in Fe2B
DB
Fe B2 (free of chemical stresses), it is necessary to use
the Newton-Raphson routine25 to solve the problem.
For this purpose, a computer program was written in
Matlab (version 6.5) to find the roots of Equation (10).
By ignoring the pressure effect on diffusion, the effective
diffusion coefficient of boron in Fe2B (see reference10 for
more details) can be evaluated from Equation (11) as
follows:
D D
C x t V E
RT vB
eff
B
Fe B Fe B2 2= +
−
⎡
⎣
⎢
⎤
⎦
⎥1
2
9 1
2( , )
( )
(11)
where DB
eff and DB
Fe B2 are the diffusion coefficients of
boron in Fe2B with and without the chemical stresses,
respectively. The origin of Equation (11) can be found
elsewhere in13–15.
Since the solid solution of boron in the matrix is
ideal, the second term in Equation (11) is positive
regardless of V . Partial molar volume V is taken to be
equal to (1.01 × 10–5 m3 mol–1).26
E = 290 GPa and v = 0.3 are the Young’s modulus
and Poisson’s ratio of the Fe2B layer, respectively27,28.
Taking the mean value of the boron concentration
throughout the Fe2B layer10, i.e.,
[ ]C x t C CFe B upFe B lowFe B2 2 2( , ) . ( )= × +05 , Equation (11) yields:
D D
V C C E
RT vB
eff
B
Fe B up
Fe B
low
Fe B
2
2 2
= +
+
−
⎡
⎣
⎢
⎤
⎦
⎥1
9 1
2 ( )
( )
(12)
3 RESULTS AND DISCUSSION
To evaluate the diffusion coefficient of boron in Fe2B
using Equation (10), the experimental results found in
the reference work29 in terms of the parabolic growth
constants at the (Fe2B/ substrate) interface and boride
incubation times were used. The chemical composition
(in mass fractions) of the gray cast iron (class 30, ASTM
A48) to be borided is the following: (3.44–3.45 % C,
1.7–1.77 % Si, 0.5–0.6 % Mn, 0.2 % Cr, 0.45–0.5 %
Cu). The boriding process was performed on the gray
cast iron under an argon atmosphere in a conventional
furnace at three temperatures (1173, 1223 and 1273) K
and for variable times (2, 4, 6 and 8) h. The boriding
medium is composed of boron carbide (B4C) as the bo-
ron source and KBF4 as the activator. Fifty measure-
ments were taken on different cross-sections of the
borided samples of the gray cast iron to estimate the
thickness of the Fe2B layer. As the input data, the model
uses the following parameters: the time, the temperature,
the upper and lower boron concentrations in the Fe2B
iron boride and the experimental values of the parabolic
B. BOUAROUR et al.: ESTIMATION OF THE EFFECTIVE BORON-DIFFUSION COEFFICIENT ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 515–520 517
Figure 3: Evolution of the squared value of the Fe2B layer thickness
versus the boriding time at different temperatures
Slika 3: Razvoj kvadratne vrednosti debeline plasti Fe2B v odvisnosti
od ~asa boriranja pri razli~nih temperaturah
Figure 2: Temperature dependence of the 	(T) parameter
Slika 2: Temperaturna odvisnost parametra 	(T)
growth constants at the (Fe2B/substrate) interface. The
effective diffusion coefficient of boron in Fe2B was
estimated on the basis of Equation (12). Figure 3 shows
the squared value of the experimental boride-layer
thickness as a function of the boriding time, according to
Equation (13):
[ ]u k t t T2 2 0= − ( ) (13)
The squared values of parabolic-growth constants (k2)
are obtained from the slopes of the straight curves using
the least-square method. The intercept with the abscissa
determines the boride incubation time at each tempe-
rature.
Table 1 gives the experimental values of parabolic-
growth constants k in the temperature range of
1173–1273 K, along with the corresponding boride
incubation times deduced from Figure 3.
Table 1: Experimental values of parabolic-growth constants at the
Fe2B/substrate interface with the corresponding boride incubation
times
Tabela 1: Eksperimentalne vrednosti konstant paraboli~ne rasti na
stiku Fe2B/podlaga z ustreznimi inkubacijskimi ~asi boriranja
T/K
Experimental para-
bolic-growth constant
k/μm s–0.5
Incubation boride time
(s)
1173
1223
1273
0.3150
0.4321
0.6122
3200
2428
1902
Table 2 lists the simulated values of DB
Fe B2 and DB
eff
obtained at each boriding temperature using Equations
(10) and (12), for the upper value of the boron content in
Fe2B 10 being equal to 59.80 × 103 mol m–3.
Table 2: Computed values of the boron-diffusion coefficients for Fe2B
and the effective diffusion coefficients of boron in Fe2B, for the upper
boron content in the Fe2B phase equal to 59.80 × 103 mol m–3
Tabela 2: Izra~unane vrednosti koeficientov difuzije bora v Fe2B in
efektivni koeficienti difuzije bora v Fe2B za zgornjo vsebnost bora v
Fe2B fazi 59,80 × 103 mol m–3
T/K DB
Fe B2 /(m2 s–1)
using Eq.(10)
DB
eff /(m2 s–1)
using Eq.(12)
1173 6.31 × 10–12 3.67 × 10–10
1223 11.881 × 10–12 6.64 × 10–10
1273 23.86 × 10–12 12.82 × 10–10
When incorporating the effect of chemical stresses on
the boron diffusion, the effective diffusion coefficient of
boron in Fe2B is increased with the boriding temperature.
So, the dependence between the diffusivity of boron in
Fe2B and the boriding temperature can be expressed with
the Arrhenius equation. The temperature dependence of
the diffusivity of boron in Fe2B is then depicted in
Figure 4. The activation energy of boron (with and
without the presence of chemical stresses) can be easily
obtained from Figure 4.
As a result, the diffusion coefficient of boron in Fe2B
in the temperature range of 1173–1273 K is given with:
D
/
RTB
eff kJ mol= ×
−−135 10
16484. exp
.
m2/s (14)
The effective diffusion coefficient of boron in Fe2B is
also determined as:
D
/
RTB
eff kJ mol= ×
−−28 10
15483. exp
.
m2/s (15)
where R is the universal gas constant (8.314 J/mol K)
and T represents the absolute temperature. The boron
activation energy obtained in this work, in the absence
of chemical stresses, was compared with the values
found in the literature29–34.
Table 3 shows a comparison of the boron activation
energies (in the absence of chemical stresses) obtained
from different borided materials. The reported values in
Table 3 differ from each other depending on different
factors such as the chemical composition of the
substrate, the boriding method and the kinetic approach
used to estimate the boron activation energy.
Table 3: Comparison of the boron activation energies (in the absence
of chemical stresses) for borided ferrous alloys
Tabela 3: Primerjava aktivacijskih energij bora (pri odsotnosti
kemijskih napetosti) za primere boriranja zlitin `eleza
Material Boron activationenergy (kJ mol–1) Reference
Armco iron
Armco iron
AISI H13
AISI 1045
Gray cast iron
Gray cast iron
Gray cast iron
151
157
186.2
169.6
177.4
175
164.8
30
31
32
33
29
34
Present work
It is seen that the calculated value of the boron
activation energy (154.8 kJ mol–1) under chemical
stresses is lower than the one for 164.8 kJ mol–1 due to
an enhancement of the boron diffusion. The temperature
dependence of the computed values of the effective
B. BOUAROUR et al.: ESTIMATION OF THE EFFECTIVE BORON-DIFFUSION COEFFICIENT ...
518 Materiali in tehnologije / Materials and technology 48 (2014) 4, 515–520
Figure 4: Temperature dependence of the diffusivity of boron in Fe2B
Slika 4: Temperaturna odvisnost difuzivnosti bora v Fe2B
diffusion coefficient of boron in Fe2B for the increasing
values of the upper boron contents in Fe2B is shown in
Figure 5.
With the presence of chemical stress, the diffusion of
boron atoms is accelerated with an increase in the
boriding temperature since the diffusion process is a
thermally activated phenomenon. At a given value of the
upper boron content in Fe2B, the computed values of
DB
eff are affected by the change in the boriding tempera-
ture.
The variation in the calculated values of the effective
diffusion coefficient of boron in Fe2B as a function of the
upper boron content in the same phase (C up
Fe B2 ) is dis-
played in Figure 6 for different boriding temperatures.
The effective diffusion coefficient of boron in Fe2B is
decreased with an increase in the upper boron content in
Fe2B. It can be explained with the saturation of the
material surface with the active boron atoms for longer
boriding times. It is concluded that the calculated values
of the effective diffusion coefficients of boron in Fe2B
vary notably with the upper boron content in Fe2B at a
fixed boriding temperature.
4 CONCLUSION
In the present work, the diffusion coefficient of boron
in the Fe2B layers grown on gray cast iron was estimated
through the mass-balance equation at the (Fe2B/sub-
strate) interface under certain assumptions. The model
included the effect of boride incubation times by forming
the Fe2B layers. Afterwards, the boron effective diffusion
coefficient in Fe2B was evaluated by applying a simple
equation based on the elasticity theory.
A lower value of the boron activation energy under
chemical stresses was obtained (154.8 kJ mol–1) for an
upper boron content in Fe2B equal to 59.80 × 103
mol m–3. This result is a consequence of the chemical
stresses enhancing the boron diffusion through the Fe2B
layers.
In addition, it is also shown that the calculated values
of the boron effective diffusion coefficient in Fe2B vary
notably with the boriding temperature. The values of the
boron effective diffusion coefficients in Fe2B are found
to be decreased with an increase in the upper boron
content in Fe2B at a given boriding temperature.
The model can be reformulated to be applied to a
bilayer configuration (FeB + Fe2B) as an extension of the
present work to estimate the boron diffusion coefficients
for the FeB and Fe2B layers in borided ferrous alloys,
incorporating the effect of chemical stresses.
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Figure 6: Variation in the calculated values of effective diffusion
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Slika 6: Spreminjanje izra~unanih vrednosti efektivnega koeficienta
difuzije bora v Fe2B v odvisnosti od zgornje vsebnosti bora v isti fazi
pri razli~nih temperaturah boriranja
Figure 5: Temperature dependence of the computed values of the
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values of C up
Fe B2
Slika 5: Temperaturna odvisnost izra~unanih vrednosti efektivnega
koeficienta difuzije bora v Fe2B pri nara{~ajo~ih vrednostih C up
Fe B2
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B. BOUAROUR et al.: ESTIMATION OF THE EFFECTIVE BORON-DIFFUSION COEFFICIENT ...
520 Materiali in tehnologije / Materials and technology 48 (2014) 4, 515–520
T. MAUDER, J. STETINA: OPTIMIZATION OF THE SECONDARY COOLING IN A CONTINUOUS CASTING PROCESS ...
OPTIMIZATION OF THE SECONDARY COOLING IN A
CONTINUOUS CASTING PROCESS WITH DIFFERENT
SLAB CROSS-SECTIONS
OPTIMIZACIJA SEKUNDARNEGA HLAJENJA PRI
KONTINUIRNEM ULIVANJU SLABOV Z RAZLI^NIMI PREREZI
Tomas Mauder, Josef Stetina
Brno University of Technology, Faculty of Mechanical Engineering, Technicka 2, Brno, Czech Republic
mauder@fme.vutbr.cz
Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2013-10-28
Although the continuous casting of steel began almost 60 years ago, its production still suffers from many serious defects in the
final structure. Cracks in the solidifying slab are mainly caused by variable thermal conditions and mechanical stresses. It is
well known that the secondary cooling zone has an important effect on the internal and surface quality. Thus, the optimal control
of the cooling intensity in secondary cooling is inevitable in order to obtain high-quality products. Nowadays, the control of the
slab temperature via numerical temperature-field models is a common practice for controlling the quality of the slabs. This leads
to cooling regulation according to the actual casting speed, casting temperature, cross-section of casting slab, chemical com-
position of steel, etc. Unfortunately, the actual practice in many steelworks is that the cooling regulation is determined as a
simple linear function of the casting speed. In order to deal with this problem, the fuzzy-optimization algorithm and numerical
model of the temperature field were created. Their combination can provide instructions for how to control the secondary
cooling and obtain high-quality of steel. This paper mainly describes the differences between the optimal cooling of different
slab cross-sections (width between 800 mm and 1600 mm, thickness between 180 mm and 250 mm). The results show that the
proper setting of secondary cooling cannot be done without a consideration of all the main casting factors.
Keywords: secondary cooling, fuzzy optimization, temperature field, continuous casting
^eprav je kontinuirano ulivanje jekla poznano `e skoraj 60 let, se v proizvodnji {e vedno pojavljajo {tevilne resne napake
slabov. Razpoke pri strjevanju slaba so v glavnem posledica spremenljivih toplotnih razmer in mehanskih napetosti. Znano je,
da ima sekundarna hladilna cona pomemben u~inek na kakovost notranjosti in povr{ine. Zato je za visoko kakovost proizvoda
neizogibna optimalna kontrola intenzivnosti sekundarnega hlajenja. Danes se zagotavlja kontrola slabov z uporabo numeri~nih
modelov temperaturnega polja. To vodi do regulacije ohlajanja, skladno z dejansko hitrostjo ulivanja, temperaturo ulivanja,
prereza ulitega slaba, kemijske sestave jekla itd. Da bi obvladali te te`ave, sta bila postavljena algoritem mehke optimizacije in
numeri~ni model temperaturnega polja. Njuna kombinacija lahko da napotke, kako kontrolirati sekundarno hlajenje in zagotoviti
veliko kvaliteto jekla. Ta ~lanek opisuje razlike med optimalnim ohlajanjem razli~nih prerezov slabov ({irine med 800 mm in
1600 mm in debeline med 180 mm in 250 mm). Rezultati ka`ejo, da pravilna nastavitev sekundarnega hlajenja ni mogo~a brez
upo{tevanja glavnih dejavnikov pri ulivanju.
Klju~ne besede: sekundarno hlajenje, mehka optimizacija, temperaturno polje, kontinuirano ulivanje
1 INTRODUCTION
Nowadays, more than 95 % of the world’s steel pro-
duction is being produced by continuous casting (CC).1,2
Using CC molten steel is formed into semi-finished pro-
ducts such as slabs, blooms and billets. The CC instal-
lation is divided into three parts. The water-cooled mold
(primary cooling zone), secondary cooling where the
steel is transported by rollers and cooled down by groups
of nozzles, and the tertiary cooling zone where the sur-
face is cooled down by free convection and radiation
only. The importance of the intensity of cooling in the
secondary cooling is well known and has been discussed
in many papers.1,3,4 These papers give the casting recom-
mendations and possible mathematical tools which can
be applied in the real system.
Unfortunately, the practice in many steel works does
not reflect the actual state of the art in this field and the
regulation of the secondary cooling is far from optimum.
There can be many reasons for this. Some steel workers
are not progressive enough; many papers are rather theo-
retical and the validation is not sufficient; mathematical
models and optimization algorithms are often less gene-
ral than is necessary; etc.
This paper deals with optimal cooling curves in the
secondary cooling zone for different slab cross-sections
(width 800–1600 mm, thickness 180–250 mm). These
cooling curves were found by using a combination of
numerical modeling and the optimization technique. The
next section describes the modeling concept, but for a
detailed description we recommend our previous work.5,6
2 DESCRIPTIONS OF THE MATHEMATICAL
ALGORITHMS
The mathematical part of the research is created by
two models. The first model is a numerical model of the
temperature field based on the governing equation of
transient heat conduction, also called the Fourier-Kirch-
hoff equation:1,5
Materiali in tehnologije / Materials and technology 48 (2014) 4, 521–524 521
UDK 621.74.047:519.61/.64 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)521(2014)
∂
∂
∂
∂
H
v
H
z
k T T


+ = ∇ ∇( ( ) )eff (1)
where keff (W/mK) is the effective thermal conductivity,
T (K) is the temperature, H (J/m3) is the volume en-
thalpy, 
 (s) is time and v (m/s) is the casting speed and
z (m) is the direction of casting. This model represents a
unique combination of numerical modeling and a large
number of experimental measurements.7 The model is
able to predict the temperature distribution in the whole
slab, the solid shell thickness and the position of the
metallurgical length (the distance below the meniscus).
Its results are validated by measurements in the real
casting process. In order to speed up the computational
time, the model could run in a parallel GPU (Graphic
Processing Unit) architecture.8
The second model is the optimization/regulation
algorithm based on fuzzy logic. The model is created
with the aim to optimize the parameters of the casting for
a given particular grade of steel, cross-sections of the
slab, casting temperature, casting speed, etc. Figure 1
shows a block scheme of the connection between the
regulator and the numerical model. The fuzzy regulator
is subordinated to the numerical model and in every time
iteration the regulator progressively tunes the cooling
parameters as the closed-loop system.
The optimization strategy is to keep the surface and
core temperatures in the specific ranges corresponding
with the hot ductility of the steel.9 The reason for this is
to avoid surface and core defects.
The presented concept creates a very general
approach to optimally control any CC process (geometry
of the slab, steel grades, caster geometry, casting limita-
tion such as allowed casting speed, water flows in secon-
dary cooling, etc.). Moreover, the algorithm can be used
for both off-line and on-line regulation. The example of
an optimal result is shown in Figure 2 for the steel grade
S355J0H. Figure 2 shows the temperature distribution at
the slab surface in different positions and the tempera-
ture distribution in the slab core. The black boxes are the
recommended temperature intervals at the slab surface to
ensure the good surface quality of the steel.
The most important indicator of iterative optimiza-
tion algorithms is usually the number of evaluations of
the problem. The evaluation of the numerical model is
very time consuming and therefore each repetition can
significantly prolong the computation. Our algorithm
was able to find the optimal parameters in less than 30
evaluations, on average. The tests were run several times
for the different grades of steels and with completely
random initial states and the number of evaluations never
exceeded 50. The results shown in Figure 2 were ob-
tained after 27 evaluations.
3 RELATIONSHIP BETWEEN CASTING SPEED
AND COOLING INTENSITY
The steelmakers operating with CC of slabs are
forced by their customers to cast different slab cross-
sections. The width of the slab is generally from 800 mm
to 1600 mm and thickness of slab from 180 mm to 250
mm, depending on the caster installation. Each slab
cross-section has typical defects. In order to avoid these
defects the caster should be set-up with a consideration
of the slab cross-section.
The casting process is influenced by many parame-
ters, but only a few of them can be controlled in
reasonable ranges. For instance, controlling the casting
temperature is not really possible and the safety proto-
cols restrict the water flows through the mold. The
typical control parameters are the casting speed and the
cooling intensity in the secondary cooling zone.
The goal of the optimization is to set the casting
speed as high as possible and still keep the good quality
of the cast steel. Controlling the surface quality can be
achieved by the presented temperature intervals in
certain points (the black boxes in Figure 2, the so-called
control points), while the inner quality is influenced by
the position of the metallurgical length. Thus, there are
two optimization constraints: the limitation of the metal-
lurgical length and the artificial value, the so-called
maximum error (the sum of temperature residuums in the
controlled points).
The search for the optimal relationship between the
casting speed and the cooling intensity is simply based
on a gradual increase of the casting speed (from 0.7
m/min to 1.3 m/min). The algorithm for every value of
the casting speed is able to find a corresponding cooling
T. MAUDER, J. STETINA: OPTIMIZATION OF THE SECONDARY COOLING IN A CONTINUOUS CASTING PROCESS ...
522 Materiali in tehnologije / Materials and technology 48 (2014) 4, 521–524
Figure 1: Block scheme of the regulation approach
Slika 1: Shematski prikaz na~ina regulacije
Figure 2: Temperatures distribution after fuzzy regulation
Slika 2: Razporeditev temperature po mehki regulaciji
intensity. From the optimization results the data where
the metallurgical length exceeded the given limit (from
14 m to 24 m) and where the maximum error exceeded
100 °C were discarded.
4 RESULTS AND DISCUSSION
The results were calculated for a casting machine
SMS Demag. The caster specifications are in Table 1.
The number of independent cooling circuits is nine,
placed according to Figure 3. The examined grade of
steel is typical low-carbon steel from Steel Group 4 No.
4038 specified by the chemical composition in Table 2.
Table 1: Slab caster machine specification
Tabela 1: Lastnosti naprave za ulivanje slabov
Ladle capacity 180 t
Tundish capacity 40 t
Mold level control Berthold mould levelmeasuring system
Torch cutting machine type
Slab Marking Machine type
Caster Machine has 9 cooling loops and electrical mold
oscillation system
Width 650–1880 mm
Length short slab 4.5–4.75 m
Length long slab 9.1–9.9 m
Table 2: Chemical composition of examined steel (w/%)
Tabela 2: Kemijska sestava preiskovanega jekla (w/%)
C Si Mn Cr
0.190 0.250 1.350 0.035
Ni Mo Cu Al
0.045 0.035 0.045 0.040
Nb Ti V
0.005 0.005 0.005
The optimal cooling curves were calculated for the
three frequently cast slab cross-sections, 1600 mm × 250
mm, 1200 mm × 200 mm, and 800 mm × 180 mm. The
results are shown (Figures 4 to 7) for the cooling circuits
on the top surface where the risk of the occurrence of
cracks is higher than on the bottom surface.
The metallurgical length and temperature intervals on
the surfaces restrict the allowed casting speed range for
all the simulated slab cross-sections. For instance, the
slab cross-section 1600 mm × 250 mm has a recom-
mended casting speed interval of between 0.7 m/min and
1.0 m/min. This result was expected, but the casting
limits are slightly different than the casting limits pre-
sented by the casting equipment. The most flexible
casting range is for the 1200 mm × 200 mm slab, while
the most critical interval is for the 800 mm × 180 mm
slab, and it probably needs special treatment. There are
more reasons for this, but we should realize that this
caster was mainly designed to cast larger cross-sections.
The second problem is the shape of the cooling
curves. The initial hypothesis that the linear relationship
between the casting speed and cooling intensity is not
suited is the clearly seen from Figures 4 to 7. The
cooling results were fitted by both linear and quadratic
regression. The results with the higher value of the
T. MAUDER, J. STETINA: OPTIMIZATION OF THE SECONDARY COOLING IN A CONTINUOUS CASTING PROCESS ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 521–524 523
Figure 5: Optimal cooling curves for cooling circuit No. 4
Slika 5: Optimalne krivulje hlajenja za hladilni krog {t. 4
Figure 3: Position of cooling circuits
Slika 3: Polo`aj hladilnih tokokrogov
Figure 4: Optimal cooling curves for cooling circuit No. 1
Slika 4: Optimalne krivulje hlajenja za hladilni krog {t. 1
coefficient of determination were chosen. Only for the
cooling circuit No. 5 and for the slab cross-section 1200
mm × 200 mm the linear regression has better results
than the quadratic. But the rest of them were clearly
non-linear.
5 CONCLUSION
The problem of optimally cast steel for different slab
cross-sections can be efficiently solved using the
described algorithm. The algorithm based on numerical
modeling with fuzzy logic is very robust and is easily
adaptable to any grade of steel, caster and slab geometry,
etc. The results show that the casting of steel slabs
should be made according to more casting indicators
than just the casting speed. The cooling behavior is
different for the different slab cross-sections and the
steelmaker should take into account this fact. Otherwise
the steel product will still be cast with the surface and
the core defects.
Acknowledgement
The authors gratefully knowledge for the financial
support of project ED0002/01/01 - NETME Centre,
project NETME CENTRE PLUS (LO1202), Junior
research project FSI-J-13-1977, project No. CZ.1.07/
2.3.00/20.0188, HEATEAM - Multidisciplinary Team for
Research and Development of Heat Processes and
project GAP107/11/1566 founded by the Czech Science
Foundation.
6 REFERENCES
1 J. P. Birat et al., The Making, Shaping and Treating of Steel: Casting
Volume, 11th edition, The AISE Steel Foundation, A. W. Cramb,
Pittsburgh, PA, USA 2003, 1000 p.
2 A. Flick, Ch. Stoiber, Trends in Continuous Casting of Steel –
Yesterday, Today and Tomorrow, In METEC InSteelCon 2011 Pro-
ceedings, Düsseldorf, Germany, TEMA Technologie Marketing AG,
2011, 8
3 Y. Ito, et al., Development of Hard Secondary Cooling by High Pre-
ssure Water Spray, In METEC InSteelCon 2011 Proceedings,
Düsseldorf, Germany, TEMA Technologie Marketing AG, 2011, 5
4 A. A. Ivanova, V. A. Kapitanov, A. V. Kuklev, Method of calculating
the optimum parameters for the air-mist cooling of a continuous-cast
slab, Metallurgist, 56 (2012), 173–179
5 T. Mauder, C. Sandera, J. Stetina, A Fuzzy-Based Optimal Control
Algorithm for a Continuous Casting Process, Mater. Tehnol., 46
(2012) 4, 325–328
6 T. Mauder, C. Sandera, J. Stetina, M. Masarik, A Fuzzy Logic
Approach for Optimal Control of Continuous Casting Process, In
METAL 2012 Conference proceedings, Ostrava, Czech Republic,
2012, 35–40
7 M. Raudensky, M. Hnizdil, J. Hwang, S. Lee, S. Kim, Influence of
Water Temperature on The Cooling Intensity of Mist Nozzles in
Continuous Casting, Mater. Tehnol., 46 (2012) 3, 311–315
8 L. Klimes, J. Stetina, Parallel dynamic solidification model of conti-
nuous steel casting on GPU, Proceedings of 22nd Conference on
metallurgy and materials, Ostrava, 2013, 34–39
9 G. S. Jansto, Steelmaking and Continuous Casting Process Metal-
lurgy Factors Influencing Hot Ductility Behavior of Niobium
Bearing Steels, In METAL 2013 Conference proceedings, Ostrava,
2013, 32–39
T. MAUDER, J. STETINA: OPTIMIZATION OF THE SECONDARY COOLING IN A CONTINUOUS CASTING PROCESS ...
524 Materiali in tehnologije / Materials and technology 48 (2014) 4, 521–524
Figure 7: Optimal cooling curves for cooling circuit No. 8
Slika 7: Optimalne krivulje hlajenja za hladilni krog {t. 8
Figure 6: Optimal cooling curves for cooling circuit No. 6
Slika 6: Optimalne krivulje hlajenja za hladilni krog {t. 6
L. KLIMES, J. STETINA: UNSTEADY MODEL-BASED PREDICTIVE CONTROL OF CONTINUOUS STEEL CASTING ...
UNSTEADY MODEL-BASED PREDICTIVE CONTROL OF
CONTINUOUS STEEL CASTING BY MEANS OF A VERY
FAST DYNAMIC SOLIDIFICATION MODEL ON A GPU
PREDVIDEVANJE KONTROLE KONTINUIRNEGA LITJA NA
PODLAGI NERAVNOTE@NEGA MODELA Z ZELO HITRIM
DINAMI^NIM MODELOM STRJEVANJA NA GPU
Lubomir Klimes, Josef Stetina
Brno University of Technology, Faculty of Mechanical Engineering, Energy Institute, Technicka 2896/2, 616 69 Brno, Czech Republic
klimes@fme.vutbr.cz, stetina@fme.vutbr.cz
Prejem rokopisa – received: 2013-10-01; sprejem za objavo – accepted for publication: 2013-10-07
The aim of the paper is to develop and test a model-based predictive control system for continuous casting of steel billets with
an emphasis on unsteady casting situations, often accompanied by abrupt changes in the casting speed. A very fast dynamic
solidification model was developed for this purpose. This fully 3D model runs on graphics processing units, GPUs, and it is
significantly faster than the recently used commercial models. Therefore, a scenario approach can be utilized, which means that
the control system, in real time, predicts and evaluates the thermal behavior of cast billets for various scenarios of the control
strategy. The concept of the effective casting speed was utilized for the proposed control strategies. The results show that the
developed control system can provide an effective control of the casting process and brings new control possibilities for
continuous steel casting.
Keywords: continuous casting, model-based predictive control, effective casting speed, dynamic solidification model, GPU,
computing
Namen prispevka je razvoj in preizku{anje nadzornega sistema kontinuirnega ulivanja gredic, ki temelji na modelu
napovedovanja, z upo{tevanjem neenakomernih razmer pri litju, pogosto spremljanih z nenadnimi spremembami hitrosti
ulivanja. Za ta namen je bil razvit zelo hiter dinami~en model strjevanja. To je popoln 3D-model, ki te~e na grafi~nih procesnih
enotah GPU in je bistveno hitrej{i kot pred kratkim uporabljani komercialni modeli. Zato se lahko uporabi scenarijski na~in, kar
pomeni, da sistem nadzora v realnem ~asu napoveduje in ocenjuje toplotno vedenje litih gredic pri razli~nih scenarijih strategije
kontrole. Za predlog nadzornih strategij je bil uporabljen koncept u~inkovite hitrosti ulivanja. Rezultati ka`ejo, da razviti sistem
nadzora ponuja u~inkovito metodo nadzora procesa litja in prina{a nove mo`nosti nadzora pri kontinuirnem ulivanja jekla.
Klju~ne besede: kontinuirno ulivanje, napovedovanje na osnovi modela, efektivna hitrost litja, dinami~ni model strjevanja, GPU,
ra~unalni{tvo
1 INTRODUCTION
Recently, the optimum control of continuous casting
has been among the main objectives of steelmakers
around the world. A proper control, especially in the se-
condary cooling with water or air-mist cooling nozzles,
is an essential issue directly related to productivity and
quality. Various techniques and approaches can be used
for this purpose. An experimental setup of a casting ma-
chine is seldom used due to a large complexity of the
problem and due to fallible results. Instead, many
researchers and steelmakers use numerical dynamic
solidification models. These systems enable the calcula-
tions of the temperature field of a cast blank and its
solidification and they can be, therefore, used for the
casting control,1,2 its optimization3–5 and for investigating
the thermal behavior.6,7
The crucial issue of dynamic solidification models in
relation to casting control is their computational de-
mands limiting their wider use in the production control
process. In particular, a dynamic solidification model
usually computes the transient heat transfer and the tem-
perature distribution of an entire cast blank in specified
points, in the so-called computing grid. This procedure
generates a large amount of calculations (in order of
billions) determining the transient temperature field of
cast blanks. The recently used commercial dynamic soli-
dification models use the state-of-the-art CPU computing
and perform these calculations within tens of minutes.
Nevertheless, tens of minutes are rather long times to
utilize the solidification models for the real-time control
or optimization of a casting process. However, a new
computing approach has recently become available for
highly parallelizable problems. This technique, called the
GPGPU (general-purpose computing on graphics pro-
cessing units), utilizes the computing on graphics cards,
GPUs. A GPU, though primarily intended for the use in
computer graphics and gaming, consists of a large num-
ber (from hundreds to thousands) of rather simple pro-
cessors allowing a huge computing performance that can
be used for solving various scientific and technical prob-
lems. The GPGPU has been mainly used for the simu-
lations of molecular dynamics,8 image processing9 and
Monte Carlo simulations.10 In spite of this, only a few
papers relating to heat-transfer problems have been pub-
Materiali in tehnologije / Materials and technology 48 (2014) 4, 525–530 525
UDK 621.74.047 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)525(2014)
lished11 by researchers and none of these covers conti-
nuous casting. However, Klimes and Stetina12,13 recently
published papers on the use of the GPGPU for a parallel
dynamic solidification model of continuous casting on a
GPU and the results show that the GPGPU can greatly
enhance the computing performance of solidification
models, allowing a significant reduction in the com-
puting time and, therefore, new opportunities for the use
of solidification models in the optimum casting control
have become available.
Based on the literature review, researchers have
recently used various approaches to the optimum control
of the continuous casting process: frequently by applying
PI or PID controllers,1 probabilistic metaheuristics (e.g.,
simulated annealing2), heuristic searches,3 fuzzy logic
techniques4 or mathematical programming methods5.
However, several issues may appear when these methods
are used or intended for a practical use in steelworks.
One of the most important issues of PI and PID control-
lers is that they regulate a process according to the
history of that process. Considering a system delay, it
means that the regulation may react improperly. Heuri-
stic methods and, particularly, the approaches via mathe-
matical programming suffer from huge computational
demands, usually because of the repeated computations
of the model. This issue can, therefore, significantly
prolong the regulation process and, for this reason, the
use of these methods is often limited in a real-time
regulation.
Due to the above reasons and encouraged by the
computational performance of the developed parallel
dynamic solidification model running on a GPU, the aim
was to propose a new control approach for continuous
casting. For this purpose the model-based predictive-
control (MPC) approach was selected. One of the
features of the MPC is that it uses a model as a nume-
rical sensor to predict the future evolution of a controlled
process under certain conditions.
2 PARALLEL DYNAMIC SOLIDIFICATION
MODEL RUNNING ON A GPU
Nowadays, commercial dynamic solidification mo-
dels for continuous steel casting utilize the state-of-the-
art computing on central processing units (CPUs). These
models usually solve heat transfer and solidification of
cast blanks in discretized points (in a number of hun-
dreds of thousands or millions) and these models require
tens of minutes to compute the stationary state under
constant casting conditions.12 These models are, there-
fore, rather awkward for a fast and real-time regulation
due to time-consuming computations. However, many
computational problems with a possibility to be paralle-
lized can be computed more efficiently with the use of
GPGPU techniques and graphics processing units
(GPUs).14 The GPGPU divides a computational problem
into independent parts that can be computed con-
currently in a parallel manner. For this reason a GPU
consists of many simple computing units that are
designed to process an identical code, but on different
data. The number of units depends on the type of the
GPU, generally varying between several hundreds and
several thousands.
The transient heat transfer and the solidification of
cast blanks can be modeled with the Fourier-Kirchhoff
equation:
∂
∂
∂
∂
H
t
k T v
H
zz
= ∇⋅ ∇ +( ) (1)
where H is the volume enthalpy (J m–3), T is the tem-
perature (K), t is the time (s), k is the thermal conduc-
tivity (W m–1 K–1), vz is the casting speed (m s–1) and z
is the spatial coordinate (m) in the direction of casting.
The mass-transfer and fluid-flow phenomena inside a
blank are usually neglected, but they can also be con-
sidered, e.g., with the use of the effective heat-con-
ductivity approach. The volume enthalpy that appears in
Equation (1) is used to include the latent heat released
during the solidification.15 This thermodynamic func-
tion can be defined as:
H T c L
fT
( ) = −
⎛
⎝
⎜ ⎞
⎠
⎟∫   
 
f
s
d
∂
∂0
(2)
where  is the density (kg m–3), c is the specific heat
(J kg–1 K–1), Lf is the latent heat (J kg–1) and fs is the
solid fraction (1).
The developed parallel dynamic solidification model
for continuous casting of steel billets (Figure 1) based
on Equation (1) and the necessary initial and boundary
conditions was created using the control-volume method
and explicit time discretization.12,13 The heat withdrawal
from the mould, within the secondary cooling zones with
the cooling nozzles and other casting conditions, can be
adjusted through the initial and boundary conditions. The
explicit discretization in time is an essential condition for
the parallel model, allowing it to run on GPUs. Basically,
the calculations related to various control volumes are
solved concurrently by various computing units of a
GPU.12,13 The CUDA C/C++ computing architecture was
L. KLIMES, J. STETINA: UNSTEADY MODEL-BASED PREDICTIVE CONTROL OF CONTINUOUS STEEL CASTING ...
526 Materiali in tehnologije / Materials and technology 48 (2014) 4, 525–530
Figure 1: Billet caster and the mesh definition
Slika 1: Naprava za ulivanje gredic in dolo~anje mre`e
used for the development of the presented solidification
model.16
A comparison of the computing performances
between the developed parallel solidification model
running on a GPU and an identical model utilizing the
state-of-the-art computing on a CPU is shown in
Figure 2. The results are presented for various numbers
of control volumes (i.e., for various computational mesh
densities) and identical casting conditions.12 GPU
NVIDIA Tesla C2075 having 448 computing units was
used as the representative of GPUs. The CPU model was
run on a computer with Intel Core 2 Quad CPU with 4
cores, each having a frequency of 2.4 GHz. As can be
seen from Figure 2, the GPGPU computing can greatly
enhance the computational performance of the soli-
dification models. For the mesh with 100000 control
volumes, the GPU model processes the computations in
only about 3 s which is 30-times faster than in the case
of the model running on a CPU. Moreover, with an
increasing mesh density the parallelism and the
computational performance of the GPU become more
significant: in the case of 3 million control volumes, the
model running on a GPU is even about 50-times faster
than the CPU model. In this case, the GPU model
performs all the computations with a very fine mesh in
only 4 min. On the contrary, the CPU model needs more
than 200 min to do the same job. These results unambi-
guously prove the benefits of the parallel GPU solidifi-
cation model, opening new possibilities for its real-time
use in continuous-casting control.
3 MODEL-BASED PREDICTIVE CONTROL FOR
CONTINUOUS CASTING
3.1 Model-based predictive control systems
On the basis of the literature review, the model-based
predictive control approach17,18 was chosen to be utilized
with the developed GPU model for the optimum control
of continuous casting process. Though primarily utilized
in the petroleum industry, the model-based predictive
control has been recently used in many engineering
applications, mainly due to a rapid development of
computers and their performance.17–19 The main principle
of the general model-based predictive control (MPC)
system is as follows: in defined consecutive time instants
with the measurements of the current and past process
outputs and with the past values of the control inputs, the
control inputs for the current and future instants are
determined with the model, so that these control inputs
minimize the differences between the predicted
controlled outputs and the required reference values, the
so-called set-points over a certain control horizon17. In
other words, the MPC utilizes the model of the process
to predict the behavior of the system resulting from the
changes in the inputs and to evaluate the consequences
of these modifications. This is actually an opposite to the
PI and PID controllers that control and regulate the
process according to the known behavior in the past. A
very illustrative comparison between the PI and PID
versus the MPC is as follows: the PI and PID controlling
is like driving a car using the rear-view mirror. On the
other hand, the MPC approach, in the context of driving
a car, is like the normal driving using the windshield of
the car.17
Inspired by the described idea of the MPC, we pro-
pose an MPC system for the continuous casting of steel
utilizing the developed, very fast GPU model. The sce-
nario approach and the effective casting speeds are then
used for determining the control inputs.
3.2 Proposal of the model-based predictive control
system for continuous steel casting
The main idea of the proposed MPC system for
continuous steel casting is based on using the developed
parallel GPU model as a numeric sensor of the real
casting machine. However, due to the developed, very
L. KLIMES, J. STETINA: UNSTEADY MODEL-BASED PREDICTIVE CONTROL OF CONTINUOUS STEEL CASTING ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 525–530 527
Figure 3: Concept of the model-based predictive control system for
continuous casting using a very fast dynamic solidification model
running on a GPU
Slika 3: Zasnova kontrolnega sistema napovedovanja na osnovi mode-
la pri kontinuirnem ulivanju z uporabo modela zelo hitrega dinami~ne-
ga modela strjevanja, ki te~e na GPU
Figure 2: Comparison of computational performances between GPU
and CPU dynamic solidification models
Slika 2: Primerjava ra~unske zmogljivosti med GPU in CPU pri
dinami~nih modelih strjevanja
fast dynamic solidification model running on a GPU, the
concept of the MPC can be extended and modified. The
entire concept of the control system is presented in
Figure 3. The main idea is to use the scenario approach.
It means that in every time step the MPC system
retrieves the actual casting parameters (e.g., the casting
speed, the casting temperature, etc.) from the casting
machine. The control system then generates several
control strategies, the so-called control scenarios in order
to control the process. Each of these scenarios represents
a possible control of the casting machine, particularly the
cooling setup (i.e., the water-flow volume) of the cooling
nozzles in the cooling circuits within the secondary zone.
The control scenarios are generated making use of the
experiences provided by the experts, the traditional rela-
tionships between the water-flow volume through the
nozzles and the casting speed, and the concept of the
effective casting speed20. The control system then pre-
dicts the future thermal behavior of cast billets for a cer-
tain future-time horizon and for all the generated control
scenarios. This task can be performed in real time
(between tens of seconds and several minutes depending
on the number of scenarios) solely due to the very fast
GPU dynamic solidification model. The control system
consequently analyses the results of the scenario thermal
behaviors and chooses the resultant control strategy used
for the real casting process. The main objective of the
resultant control strategy is to ensure that the surface
temperatures of cast billets fit the predefined temperature
intervals (to avoid subcooling or overcooling because of
a low ductility of steel21 resulting in surface defects, e.g.,
cracks) and that the metallurgical length fits the defined
range.3,4 The control loop is, thereby, closed and the
procedure is repeated in the next time instant.
3.3 Effective casting speed
The concept of the effective casting speed was
adopted for the determination of the control cooling
strategies in secondary cooling.20 When the actual cast-
ing speed rapidly varies in time (e.g, due to a change of
tundishes), the determination of the water-flow volume
through the nozzles as a function of the actual casting
speed may result in inappropriate cooling, particularly in
the cooling circuits far from the mould.20 Researchers
and steelworkers usually tend to use the dependencies of
the water-flow volume in the cooling circuits as the qua-
dratic functions of the casting speed.20 The mentioned
problem can be overcome with the use of the effective
casting-speed approach, based on considering the dwell
time that a blank divided into slices spends in each of the
cooling circuits. The effective casting speed ve,i for the
i-th cooling circuit (m min–1) can be calculated as:
v v vi i i i ze, a,= + − ( )1 (3)
where i is the weight coefficient (1), vz is the actual
casting speed (m min–1) and the average casting speed
va,i for the i-th cooling circuit (m min–1) is determined
as:
v
n L
t
a i
i i
r i j
j
n i,
, ,
=
−
∑
1
(4)
where ni is the number of slices, Li is the distance from
the mould (m) and tr,i,j is the residential time (s).20
Figure 4 shows the effective casting speed for the
fluctuating actual casting speed and for the casting
machine with three cooling circuits within the
secondary cooling zone.
As can be seen from Figure 4, the effective casting-
speed approach can be suitably used for the characteriza-
tion of the fluctuating casting conditions. Particularly in
the cooling circuits far from the mould, the effective
casting speed has a smoothness effect, which fairly
corresponds to the reality.
4 USE OF THE CONTROL SYSTEM, RESULTS,
DISCUSSION
The developed MPC system was tested for the con-
trol of the secondary cooling in the case of a temporary
change in the casting speed, e.g., due to the change of
tundishes. The dynamic solidification model was confi-
gured for a caster with 6 cooling circuits within the secon-
dary cooling incorporating 180 JATO cooling nozzles of
several types. The caster casts 200 mm × 200 mm steel
billets, normally with the casting speed of 1.5 m/min.
The temporary drop in the casting speed from 1.5 m/min
to 0.8 m/min was assumed for 6 min as depicted in
Figure 5 where the effective casting speeds for each of
the cooling circuits are presented as well.
The aim of the control was to determine the cooling
strategy of the cooling nozzles within the secondary
cooling so that the surface temperatures on the cast
billets would be maintained as close as possible to the
surface temperatures in the case when no change is made
to the casting speed. Thereby, the steady-state surface
L. KLIMES, J. STETINA: UNSTEADY MODEL-BASED PREDICTIVE CONTROL OF CONTINUOUS STEEL CASTING ...
528 Materiali in tehnologije / Materials and technology 48 (2014) 4, 525–530
Figure 4: Concept of the effective casting speed
Slika 4: Zasnova u~inkovite hitrosti litja
temperatures were considered as the optimum target tem-
peratures. The resultant surface temperatures on the top
(small-radius) surface of the control process are shown
in Figure 6. These temperatures for the cooling circuits
4, 5, 6 are plotted for the time when 5 min have elapsed
from the beginning of the change in the casting speed
(depicted by the vertical dotted line in Figure 5) and
when the actual casting speed begins to increase from
0.8 m/min back to 1.5 m/min.
As for the described control problem of the tempo-
rary change in the casting speed, the developed model-
based control system generated 15 cooling scenarios for
the cooling strategy within the secondary cooling. These
scenarios were generated taking into account the effec-
tive casting speeds. The system then predicted the future
thermal behavior (calculating the complete temperature
distribution) for all these cooling scenarios. The compu-
tations were performed for a computational mesh with
1 million control volumes and all these computations for
all the cooling scenarios were performed within 1 min
using GPU NVIDIA Tesla C2075. The temperature
distributions for three representative scenarios (denoted
as CS-1, CS-2, and CS-3) are plotted in Figure 6
together with the optimum steady-state temperature
distribution and with the temperature distribution for the
case when no change in the cooling strategy is
performed. As can be seen from Figure 6, when no
change in the cooling strategy is considered, a significant
subcooling occurs within the secondary cooling. Further,
the generated cooling strategy CS-1 is also rather
inappropriate due to the subheating, mainly occurring
within the fourth cooling circuit. Similarly, the cooling
strategy CS-2 is also rather inconvenient owing to the
overheating, especially in the fifth and sixth cooling
circuits. But in the case of the cooling strategy CS-3, the
distribution of the temperature is very close to the target
temperature distribution in all three cooling circuits
plotted in Figure 6. The cooling scenario SC-3 can,
therefore, be considered as the optimum cooling strategy
for the investigated situation.
5 CONCLUSION
The paper presents a proposal for a model-based
predictive control system for the continuous casting of
steel. The system is based on the dynamic solidification
model as the numerical sensor for predicting the future
thermal behavior of cast billets under specific casting
conditions. The control system utilizes the developed
very fast dynamic solidification model that runs on gra-
phics processing units, GPUs. The solidification model,
using the concepts of the scenario-based cooling
strategies and the effective casting speed was then used
for the optimum control of continuous casting. The con-
trol system was tested for the case of a temporary drop in
the casting speed and the results show that the proposed
system is a promising tool for solving these control
problems. The next research will aim at tuning up the
developed control system and testing it in various
production control situations in the continuous casting of
steel.
Acknowledgement
The presented research was supported by project
GACR P107/11/1566 of the Czech Science Foundation,
by the NETME Centre, ED0002/01/01, by the NETME
CENTRE PLUS (LO1202), and by the BUT project
FSI-J-13-1977 for young researchers. The principal
author, the former holder of the Brno PhD Talent
Financial Aid sponsored by the Brno City Municipality,
also gratefully acknowledges the financial support.
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Slika 6: Temperature povr{ine na vrhu gredic (manj{i polmer), ki so
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530 Materiali in tehnologije / Materials and technology 48 (2014) 4, 525–530
B. [U[TAR[I^ et al.: DSC/TG OF Al-BASED ALLOYED POWDERS FOR P/M APPLICATIONS
DSC/TG OF Al-BASED ALLOYED POWDERS FOR P/M
APPLICATIONS
DSC/TG PRAHOV NA OSNOVI Al PRIMERNIH ZA P/M UPORABO
Borivoj [u{tar{i~1, Jo`e Medved2, Sre~ko Glode`3, Marko [ori3, Albert Koro{ec4
1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2OMM, NTF, University of Ljubljana, A{ker~eva 12, 1000 Ljubljana, Slovenia
3University of Maribor, FNM, Koro{ka cesta 160, 2000 Maribor, Slovenia
4Talum, Tovarna aluminija d. d., Tovarni{ka cesta 10, 2325 Kidri~evo, Slovenia
borivoj.sustarsic@imt.si
Prejem rokopisa – received: 2013-11-13; sprejem za objavo – accepted for publication: 2014-04-16
Al-based alloyed powders, appropriate for the sintering procedure (powder metallurgy, P/M) contain the alloying elements with
a high solid solubility in Al, enabling reaction and liquid-phase sintering. They are surface oxidised because of a high affinity of
Al to oxygen. Besides, this type of powders contains a polymeric lubricant (wax), which reduces the friction on die walls during
automatic die compaction into the final compact shape of a product. This lubricant has to be removed slowly during the first
stage of sintering in order to prevent deformations and cracking of the product. Consequently, its sintering is very complex.
Generally, these powders are sintered in pure nitrogen with a low dew point. The optimum sintering conditions are generally
determined on the basis of light and scanning electron microscopy. The investigation can also be completed very successively
with differential scanning calorimetry and thermo gravimetry. The first one allows an insight into the endo- and exothermic
reactions, taking place during the heating and cooling of a compacted metal powder, and the second one allows an insight into
the processes, connected with the mass loss (a reduction, a lubricant removal, etc.) or mass increase (an oxidation). The
DSC/TG of three commercial Al-based alloyed powders was performed in the frame of our investigations. The results were
compared with the theoretical thermodynamic-based calculations and the optimum sintering conditions were proposed.
Keywords: Al-based alloyed powders, sintering, differential scanning calorimetry and thermo gravimetry (DSC/TG)
Legirani prahovi na osnovi Al, primerni za sinter postopek (P/M, metalurgija prahov), vsebujejo zlitinske elemente z veliko
topnostjo v trdnem Al, kar omogo~a reakcijsko sintranje v prisotnosti teko~e faze. Zaradi velike afinitete aluminija do kisika so
na povr{ini oksidirani. Poleg tega vsebujejo ti prahovi polimerno mazivo, ki zmanj{uje trenje na stenah orodja med avtomatskim
stiskanjem prahu v kon~no oblikovan izdelek. To mazivo moramo v prvi fazi procesa sintranja po~asi odstraniti, sicer bi lahko
pri{lo do nepopravljive deformacije ali celo pokanja izdelka. Zato je njihovo sintranje zelo zahtevno. Navadno se sintrajo v ~isti
du{ikovi atmosferi z nizko temperaturo rosi{~a. Poleg fizikalno-kemijske karakterizacije sintranih izdelkov s svetlobno in elek-
tronsko mikroskopijo je za dolo~itev optimalnih pogojev sintranja zelo uporabna diferencialna vrsti~na kalorimetrija, kombini-
rana s termogravimetrijo (DSC/TG). Prva omogo~a vpogled v endo- in eksotermne reakcije, ki potekajo med segrevanjem in
ohlajanjem kompaktiranega kovinskega prahu, druga pa s temi procesi povezano izgubo (redukcija, odstranitev maziva itd.) ali
prirastek (oksidacija) mase. V okviru na{ih raziskav smo izvedli DSC/TG treh komercialno dosegljivih prahov na osnovi Al,
ugotovljene reakcije smo primerjali z napovedmi teoreti~ne termodinamike in predlagali optimalne pogoje sintranja teh prahov.
Klju~ne besede: prahovi zlitine na osnovi Al, sintranje, diferencialna vrsti~na kalorimetrija s termogravimetrijo
1 INTRODUCTION
Al-based alloyed powders, appropriate for the sinter-
ing procedure (powder metallurgy, P/M) contain the
alloying elements (Cu, Zn, Mg, etc.) with a high solid
solubility in Al enabling reaction and liquid-phase sinter-
ing. Generally, these powders are surface oxidised be-
cause of a high affinity of Al to oxygen. Besides, these
types of powders contain mass fraction approximately w
= 1.5 % of a polymeric lubricant, which reduces the fric-
tion on die walls during automatic die compaction
(ADC) into the final compact shape of a product. This
lubricant has to be removed slowly during the first stage
of sintering in order to prevent deformations and
cracking of the product. Therefore, its sintering is very
complex. Generally, these types of powders are sintered
in pure nitrogen (N2, 5.9) with a low dew point (below
–40 °C). The optimum sintering conditions are common-
ly determined on the basis of light (LM) and scanning
electron microscopy combined with a microchemical
analysis based on the measurement of the dispersed kine-
tic energy of X-rays (an energy dispersive X-ray
spectrometer – SEM/EDS).1 The investigation can also
be completed very successively with heating microscopy,
as well as differential scanning calorimetry and thermo
gravimetry (DSC/TG). The DSC method allows an
insight into exothermic/endothermic reactions, and the
TG method allows an insight into the mass
increase/decrease occurring during the heating/cooling
of a compact, respectively.
DSC is an effective and widely used method of the
thermal analysis (TA) of metallic and other materials. It
is a modern, completely automated and highly improved
version of an older method known as differential thermal
analysis (DTA) where the temperature differences bet-
ween the investigated and the standard (neutral, usually
alumina) samples are measured. The temperature diffe-
rences are the consequence of the heat release/consume
of the exothermic/endothermic reactions associated with
different physical and chemical processes (melting, soli-
Materiali in tehnologije / Materials and technology 48 (2014) 4, 531–536 531
UDK 621.762.5:669.71 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)531(2014)
dification, evaporation, oxidation, reduction, solid-state
transformations, etc.) occurring during the heating and
cooling of the investigated sample. The actual difference
between the DTA and DSC methods is a more precise
determination of the released or consumed heat. With the
DSC method it is possible to determine the released or
consumed heat much better because the specific heats
and their temperature dependencies of the investigated
samples are considered. The new DSC devices also have
a significantly higher number of temperature sensors
(thermocouples) in a very small space of the measuring
cell and an improved calibration system of the thermal
buoyancy, enabling a much better measuring of tempera-
ture gradient dT/dt (°C/s). The measuring cell is usually
placed on a very precise balance. This also enables ther-
mogravimetry (TG) and a simultaneous tracking of the
mass changes due to different reactions. Modern DSC/
TG devices enable experimental work with a controlled
heating/cooling rate in different stationary or flow atmo-
spheres. They enable the heating up to very high tem-
peratures (for the metals, generally, up to 1600 °C; the
maximum of 2400 °C is also possible). Mass changes are
the results of the changes in the investigated sample (a
wax removal, an evaporation, etc.) or its reaction with
the selected atmosphere (hydrogen, oxygen, water va-
pour, etc.).
The DSC/TG analyses of three commercial2 Al-based
alloyed powders were performed in the frame of our
investigations. The results were compared with the theo-
retical thermodynamic-based calculations (ThermoCalc)3
and the optimum sintering conditions were proposed.
Our experiments were performed with a Netzsch STA
(simultaneous thermal analysis) device, Germany.4,5
2 EXPERIMENTAL WORK
The actual chemical compositions of the selected
commercial Al powders are given in Table 1. From this
table one can notice that the first AlCuSiMg-based alloy
(A) is of type 2xxx (2014), the second alloy (B) is a
special hypereutectoid Al-Si-based alloy with a high Si
content and additions of Cu and Mg. The third alloy (C)
is an AlZnMgCu-based alloy of 7xxx family type (7075).
The powders were compacted on an ADC mechanical
press (Dorst, Germany, 60 kN) with a 450 MPa pressure
into standard tensile-test specimens (DIN ISO 2740).6
Thirty-five (35) pieces of each alloy were prepared for
further experiments and investigations. From three (3)
characteristic specimens (with the average green density)
samples of approximately 5 g were cut off for the DSC/
TG experiments. As already mentioned, the experiments
were performed with a Netzsch STA 449C Jupiter5
device installed at the Laboratory for Thermodynamics
of Materials, Department for Metallurgy and Materials,
University of Ljubljana. This equipment enables a simul-
taneous performance of differential scanning calorimetry
and thermogravimetry with the selected heating/cooling
rate and atmosphere conditions. In our case the samples
were heated/cooled at the constant rate of 5 °C/min in a
stationary atmosphere of Ar (the purity of 5.9), as well as
in a flow atmosphere (10 mL/h) of pure nitrogen (5.9).
The samples were heated up to the maximum tempera-
ture of 650 °C and then cooled down. The calibration
(determination of the base line) of the device as well as
the evacuation of the cell were performed every time be-
fore the start of an analysis.
Table 1: Actual average chemical compositions of the selected pow-
ders in mass fractions (w/%)
Tabela 1: Povpre~na kemijska sestava izbranih prahov v masnih dele-
`ih (w/%)
Designation Cu Si Mg Zn Al Remarks
Chemical
composition w/%
Alloy A 4.5 0.62 0.48 – bal 0.08 % Fe
Alloy B 2.7 15.0 0.58 – bal
Alloy C 1.6 – 2.40 5.8 bal 0.29 % Sn
The average bulk chemical compositions of the
powders were determined with a classical Agilent 720
ICP-OES (ion coupled plasma - optical emission spec-
troscopy) instrument with a limit of detection (< 0.001 %
of an individual element). However, the microchemical
compositions of individual powder particles were deter-
mined with SEM/EDS (a JEOL FE HR JSM-6500F and
Oxford EDS INCA Energy 450, X-Sight LN2 detector).
The compacted tensile-test specimens were also sin-
tered at the selected sintering conditions. The obtained
mechanical properties are given in Table 2. They are in
accordance with the expectations and the powder-produ-
cer specifications.2
The standard metallographic samples of the powders
and sintered materials were prepared for the microstruc-
tural and microchemical investigations under LM and
SEM/EDS. In this article, the description is focused
B. [U[TAR[I^ et al.: DSC/TG OF Al-BASED ALLOYED POWDERS FOR P/M APPLICATIONS
532 Materiali in tehnologije / Materials and technology 48 (2014) 4, 531–536
Table 2: Average mechanical properties of the materials after the ADC and sintering of the tensile-test specimens made of the selected Al-based
powders
Tabela 2: Povpre~ne mehanske lastnosti materialov, dobljene po stiskanju in sintranju nateznih preizku{ancev iz izbranih Al-prahov
Alloy
Green
density
Sintered
density HardnessHB2,5/31,25
Tensile
strength
Yield
strength
Young’s
modulus
Elongation
%
g/cm3 MPa
A 2.62 2.60 65 202 156 3834 2.23
B 2.52 2.62 104 239 219 4399 0.70
C 2.61 2.73 102 325 250 4102 3.90
mainly on DSC/TG investigations and analyses. The
description of the other investigations can be found
elsewhere.1
3 RESULTS AND DISCUSSION
This article will only focus on some general charac-
teristics and the most important findings. An additional
description of the results of the investigations can be
found elsewhere.1 Figure 1 shows the typical measuring
protocol and the results of the DSC analysis obtained for
both experimental conditions during the investigation of
alloy A. The heating/cooling program is recorded with a
fine dotted line. One can notice that the resulting DSC
curve for the analysis performed with stationary Ar is
almost the same as the one performed with a flow of
nitrogen. Only small differences can be noticed. They
can be ascribed to the fluctuations in the sample com-
position, solubility of gas in the melt, possible formation
of nitrides/oxides and variations of the used experimental
conditions. The sole large exothermic (exo) and equiva-
lent endothermic (endo) peaks can be ascribed to the
melting and solidification of the alloy, respectively. The
exo peak connected with the solidification is sharper for
the nitrogen-flow experiment. Other changes are less
distinctive and can be observed only in the magnified
mode. The cooling parts of the DSC curves have no
visible additional peaks, which means that, during the
melting, the homogeneous alloy was completely formed,
without any significant precipitation of the secondary
phases during the cooling of the sample.
Figure 2 shows the heating part of the DSC curve in
the magnified mode obtained with the DSC analysis of
alloy A. In the temperature region between 100 °C and
500 °C there is a number of small less recognizable endo
peaks. The first one at 138 °C (i.e., 141 °C; the data for
the experiments in stationary Ar is given in parentheses)
can be associated with the beginning of the melting of
the polymer wax, but the later peaks at ((261), 322 (328),
(348) and 440 (442)) °C are the results of the evapora-
tion of the multicomponent lubricant system. A constant
slight ascent of the DSC curve in this temperature region
can also be noticed. This is the result of the exothermic
nature of the wax removal (burning). Simultaneously, at
the temperatures above 400 °C, above the solvus line of
the complex alloying system, one can also expect the
beginning of the formation of the final solid solution
(-Al) because of the increased diffusion of the alloying
elements into the aluminium. The SEM/EDS microche-
mical analyses of loose powders showed1 that not all the
individual (pre-alloyed) powder particles are of the same
chemical compositions, but they are of different
compositions forming the final mixtures with the average
compositions given in Table 1. For example, SEM/EDS
analyses1 showed that alloy A consisted of the powder
particles of pure Al, alloys Al-10Si-Mg (but Mg was not
detected) and Cu-5Al. Alloy B consisted of the powder
particles made of the Al-3Si, Al-27Si-1Mg-6Cu and
Si-Al-based alloys, and, finally, alloy C consisted of the
particles of pure Al and the particles of the Al-Zn-Mg-
Cu alloy (approximately w(Al) = 79 %, w(Mg) = 5 %,
w(Cu) = 4 %, w(Zn) = 12 %). Thus, the endo peaks at
approximately (452, 486, 482 (484) °C and 506 (509))
°C can be associated with the diffusion of the alloying
elements and the formation of the -Al solid solution.
EDS analyses also showed that alloy B had the most
oxidised particles (a high Si content) and alloy C was the
least oxidised of all.
The theoretical calculation of the thermodynamic
equilibrium with ThermoCalc3 predicts the formation of
the first melt at 525 °C (Figure 3). On the experimental
DSC curve, it is visible at 534 °C. However, the real
large endo peak caused by the melting starts at 567.8
(573.2) °C and finishes at approximately 640 °C when
the melting of the alloy is completed. ThermoCalc
predicts the two-phase region (-Al+L) between 525 °C
and 638 °C (Figure 3). This is in a relatively good agree-
ment with the experimental results of the DSC analysis.
From these results one can find that the optimum
liquid-phase sintering temperature is somewhere in the
middle of the -Al+L region. This is in a good agree-
ment with the powder producer that recommends, for
B. [U[TAR[I^ et al.: DSC/TG OF Al-BASED ALLOYED POWDERS FOR P/M APPLICATIONS
Materiali in tehnologije / Materials and technology 48 (2014) 4, 531–536 533
Figure 2: Heating part of the DSC curve of alloy A, protective atmo-
sphere of N2 5.9, flow of 10 mL/h, 1 × vacuum, Tmax = 650 °C, 5 °C/min
Slika 2: DSC segrevalna krivulja zlitine A, za{~. atm. N2 5,9, pretok
10 mL/h, 1 × vakuum, Tmaks. = 650 °C, 5 °C/min
Figure 1: Measuring protocol of the DSC analysis of alloy A (heat-
ing/cooling is shown with a fine dotted line), a comparison of the DSC
curves obtained during the experiments performed in stationary Ar
(dashed line) and in a flow of N2 (full line)
Slika 1: Merilni protokol DSC-analize zlitine A (segrevanje/ohlajanje
– fina pik~asta ~rta); primerjava DSC-krivulj, izdelanih v
stacionarnem argonu (~rtkana ~rta) in pretoku du{ika (polna ~rta)
alloy A, a sintering temperature between 590 °C and 600
°C. In this case most of the intermetallic phase is already
dissolved in the solid solution of Al and a small amount
of the liquid needed is also present.
The thermodynamic analysis with ThermoCalc also
shows (Figure 3) that, in the equilibrium, alloy A con-
tains solid crystals of the homogeneous solid solution of
-Al from room temperature up to 638 °C and mainly
intermetallic phase Al2Cu (the 
 phase) up to 501 °C. As
already mentioned, the first liquid appears at 525 °C.
The theory also predicts possible formations of phases
Al5Cu2Mg8Si6 (up to 500 °C), Al7Cu2M (M = Fe, up to
565 °C), 	AlFeSi (up to 223 °C) and Si up to 396 °C.
From this analysis one can conclude that alloy A is a
typical precipitation-hardening alloy. Therefore, an im-
provement in the mechanical properties of this alloy can
be achieved with a combination of homogenization
annealing at the temperature of a complete solid solution
(for example, 500 °C, 20 min), fast cooling and natural
(T4) or artificial ageing (T6, for example, 150 °C, 15
min). The optimization of the ageing parameters can be
performed with the help of theoretical and experimental
CCT (continuous-cooling temperature) diagrams. The
optimum process parameters (temperature/time) of
ageing lead to an alloy with a fine uniform dispersion of
the precipitates of the intermetallic phases in the -Al
solid solution. In the case of high cooling rates (103 °C/h)
only a formation of very fine GP (Guiner-Preston) zones
and, eventually, a fine 
’ (Theta prime) phase can be
expected. Lower cooling rates are undesirable because of
the formation of other larger intermetallic phases such as
S’ or 
 phases.7,8
The experimental cooling DSC curve (Figure 4)
shows only one large exo peak with its beginning at
637.9 (637.2) °C, which is the result of the sample soli-
dification during cooling. The solidification is finished
between approximately 580 °C and 590 °C. Later, only
one, almost invisible exo peak appears at approximately
570.1 (563.5) °C. This could be associated with the
precipitation of intermetallic phases. The precipitation of
phases Al2Cu and Mg2Si is possible during a relatively
slow cooling (5 °C/min). This could only be confirmed
B. [U[TAR[I^ et al.: DSC/TG OF Al-BASED ALLOYED POWDERS FOR P/M APPLICATIONS
534 Materiali in tehnologije / Materials and technology 48 (2014) 4, 531–536
Figure 4: Cooling DSC curve of alloy A, protective atmosphere of N2
5.9, flow of 10 mL/h, 1 × vacuum, Tmax = 650 °C, 5 °C/min
Slika 4: Ohlajevalna DSC-krivulja zlitine A, za{~itna atmosfera N2
5,9, pretok 10 mL/h, 1 × vakuum, Tmaks = 650 °C, 5 °C/min
Figure 3: a) Theoretical equilibrium thermodynamic phase stability of
alloy A and b) detail of the diagram up to a fraction of 0.1 mol, cal-
culated with ThermoCalc3
Slika 3: a) Teoreti~na ravnote`na termodinamska stabilnost faz zlitine
A in b) detajl v diagramu do 0,1 molskega dele`a, izra~unano z oro-
djem ThermoCalc3
Table 3: Comparison of the results of DSC analyses in the flow of N2 for the selected powders
Tabela 3: Primerjava rezultatov izvedenih DSC-analiz v toku N2 za izbrane prahove
Alloy
Dewaxing* Alloy formation* Melting* Solidification**
(temperature, °C) (temperature, °C) (temperature, °C) (temperature, °C)
Start Intermediate Finish Start Intermediate Finish Start Finish Start Finish
A 138 348 447 440 506.2 534.2 567.8 650 637.9 570.1
B 138 344 426 498 509.0 - 526.0 600 520.0 615.0
C 139 332 406 466 - - 542.0 655 542.0 638.0
* …… endo peaks, ** …… exo peaks
with XRDS (X-ray diffraction spectroscopy) of the
metallographic samples (not performed yet).
The results of the DSC experiments for alloys B and
C are compared with the ones for alloy A in Table 3.
The theoretical calculations of the thermodynamic equi-
librium performed with ThermoCalc were also obtained
for alloys B and C. The temperatures of the thermo-
dynamic stability of all the phases are given in Table 4.
Figure 5 shows a typical measuring protocol and the
results of the TG analysis obtained for both experimental
conditions during the investigation of alloy A. The heat-
ing/cooling program is recorded with a dashed line. One
can notice that the resulting TG curve for the analysis
performed with stationary Ar is almost the same as the
one performed with the flow of nitrogen. The cooling
part of the TG curve shows no changes and can, there-
fore, be omitted. Figure 6, therefore, shows only the
heating TG curve of alloy A obtained for the flow of
nitrogen.
From Figure 6 one can clearly see that the mass
decrease starts at 248 (247) °C because of the lubricant
removal (dewaxing). It is completely finished at 447
(446) °C. The total mass loss in this temperature region
is 1.42 (1.38) %. This result is in accordance with the pow-
der-producer specification (the wax content w = 1.5 %)
considering the fluctuation of the wax content inside the
volume of the compacted tensile-test specimen. Surpri-
singly, the mass of the sample again starts to increase
after dewaxing up to w = 0.53 (0.48) %. At the higher
temperatures, a reoxidation of the sample is possible and
most likely to happen. The first experiments were per-
formed in the pure stationary Ar, and the first assumption
was that either the entrapped vapour of the wax caused
the sample to reoxidise, or a small leakage of the cell
occurred. Therefore, the new DSC/TG experiments were
performed in the flow of nitrogen. As one can see, the
results of the analyses are very similar (Figure 6) and
the final conclusion is that the reoxidation is a conse-
quence of the entrapped molecules of air in the green
compact because the released vapours of the wax are
removed from the cell together with the flow of nitrogen.
The powder producer recommends the following
dewaxing procedure during the sintering process: the
flow of N2 in the temperature region between 380 °C and
420 °C for 20 min. However, our TG experiments show
that the dewaxing starts much earlier (at approximately
245 °C) and finishes later (at approximately 450 °C).
The process of dewaxing must be, therefore, optimized
by adequately slowing down the heating procedure in
this temperature region. The results of the TG experi-
ments involving alloys B and C are compared with the
ones for alloy A in Table 5. Dewaxing of the samples
made of the green compacts of alloys B and C starts at
even lower temperatures (below 200 °C) in the case of
the TG performed with stationary Ar. In spite of this,
dewaxing of the green compacts made of these materials
can be performed in the same way as for alloy A. An
interesting result of the TG analyses is that alloy (pow-
B. [U[TAR[I^ et al.: DSC/TG OF Al-BASED ALLOYED POWDERS FOR P/M APPLICATIONS
Materiali in tehnologije / Materials and technology 48 (2014) 4, 531–536 535
Table 4: Theoretical temperature thermodynamic phase stability of the selected alloying systems, calculated with ThermoCalc3
Tabela 4: Teoreti~na temperaturna termodinamska stabilnost faz v izbranih zlitinskih sistemih, izra~unana z orodjem ThermoCalc3
D
es
ig
na
ti
on
te
m
pe
ra
tu
re
,
°C

-A
l
A
l 2
C
u
A
l 5
C
u 2
M
g 8
S
i 6
A
l 7
C
u 2
M
S
i
A
lF
eS
i_
B
et
a
M
g 2
X
_C
1
S
_A
l 2
C
uM
g
M
gZ
n 2
T
_A
lC
uM
gZ
n

-A
l
+
L
L
Alloy A < 637.8 < 501.0 < 500 223.2–565.0 < 395.8 < 223 – – – – 525.0– 637.8 > 525.0
Alloy B < 562.6 < 447.8 < 526 – < 619.7 – – – – – 532.6–562.6 > 532.6
Alloy C < 630.0 – – – – – < 511.1 < 442.7 < 409.7 > 250.1 516.5–630.0 > 516.5
Figure 6: Heating TG curve of alloy A, protective atmosphere of N2
5.9, flow of 10 mL/h, 1 × vacuum, Tmax = 650 °C, 5 °C/min
Slika 6: Segrevalna TG-krivulja zlitine A (sprememba mase vzorca
med segrevanjem), za{~. atm. N2 5,9, pretok 10 mL/h, 1 × vakuum,
Tmaks = 650 °C, 5 °C/min
Figure 5: Measuring protocol of the TG analysis of alloy A (heating/
cooling is presented as a fine dotted line), comparison of the TG
curves obtained during the experiment performed in stationary Ar
(dashed line) and the flow of N2 (full line)
Slika 5: Merilni protokol TG-analize zlitine A (segrevanje/ohlajanje –
pik~asta ~rta), primerjava TG-krivulj, dobljenih med preizkusom v
stacionarnem pretoku Ar (~rtkana ~rta) in pretoku N2 (polna ~rta)
der) C with the lowest original oxidation was reoxidised
the most during the TG test and that alloy B with the
highest original oxidation was reoxidised the least during
the TG test. One can speculate that, under the given con-
ditions (time/temperature/atmosphere), each Al-based
alloy has a specific oxidation potential if exposed to
oxygen.
SEM/EDS analyses1 showed that powder particles are
surface oxidised and that the selected powders are mix-
tures of different alloys. During ADC some molecules of
air are also entrapped. The exact wax composition is not
known and it is, therefore, difficult to determine the na-
ture of its melting, evaporation and oxidation. Generally,
it is the producer’s know-how. But, the waxes appro-
priate for ADC generally consist of multicomponent
systems. The performed DSC/TG analyses can give only
the basic but important information about the optimiza-
tion of dewaxing and sintering processes, respectively.
4 CONCLUSIONS
In the frame of the present work the theoretical ther-
modynamic analyses and microstructure characterization
of the selected Al-alloyed powders are completed with
DSC/TG analyses. These provide a more precise insight
into the events occurring during the heating and cooling
of the green compacts made of the selected Al powders,
as well as the optimization of the sintering process. The
investigations show that DSC/TG experiments must be
performed very carefully and must be, as much as possi-
ble, similar to the performed sintering procedure. In the
first heating phase, the mass of the samples gradually
decreases because of dewaxing. This is associated with
the endo (melting and evaporation) and exo (burning)
reactions detected on DSC heating curves. Unexpectedly,
in all the cases the increase in the sample mass starts
above approximately 450 °C. The most probable reason
for this is the reoxidation of the samples with the mole-
cules of air entrapped in the green compact. DSC ana-
lyses have also shown when exactly, during the heating,
the solution of the alloying elements is formed and the
homogeneous solid solution of -Al, the formation of the
first melt, takes place, and when the general melting of
the sample starts and finishes. During cooling, DSC ana-
lyses have shown the start and finish of the solidification,
also indicating when the precipitation of the secondary
intermetallic phases occurs at a given cooling rate. These
analyses were completed with the theoretical thermo-
dynamic calculations that allow a better understanding of
the microstructure evolution, the optimization of heat
treatment (precipitation hardening) and an improvement
in the mechanical properties.
Acknowledgement
The authors wish to thank the Slovenian Research
Agency for the financial support in the frame of project
No. L2-4283.
5 REFERENCES
1 B. [u{tar{i~, I. Paulin, M. Godec, S. Glode`, M. [ori, J. Fla{ker, A.
Koro{ec, S. Kores, G. Abramovi~, Morphological and microstructural
features of Al-based alloyed powders for powder-metallurgy appli-
cations, Mater. Tehnol., 48 (2014) 3, 439–450
2 ECKA Granules, http://www.ecka-granules.com/en/ecka-granules/
products/
3 Thermocalc software, http://www.thermocalc.com/Software.htm
4 http://www.netzsch-thermal-analysis.com/en/home.html
5 Netzsch: STA 449 F1: Simultaneous Thermal Analysis, Method,
Technique, Applications, http://private.netzschcdn.com/uploads/
6 International Standard ISO 2740-1986 (E), Sintered metal materials
–Tensile test pieces, 1986-10-01, (connected with ISO 6892, Metallic
materials- Tensile testing)
7 J. F. Chinella, Z. Guo, Computational Thermodynamics Character-
ization of 7075, 7039, and 7020 Aluminum Alloys Using JMatPro,
ARL, MD, USA, 2011
8 Z. Guo, N. Saunders, A. P. Miodownik, J. P. Schillé, Prediction of
room temperature mechanical properties in aluminium castings, Pro-
ceedings of the 7th Pacific Rim International Conference on Model-
ing of Casting and Solidification Processes, Dalian, China, 2007
B. [U[TAR[I^ et al.: DSC/TG OF Al-BASED ALLOYED POWDERS FOR P/M APPLICATIONS
536 Materiali in tehnologije / Materials and technology 48 (2014) 4, 531–536
Table 5: Comparison of the results of the TG analyses of the selected powders
Tabela 5: Primerjava rezultatov izvedenih TG-analiz izbranih prahov
Designation
Dewaxing Reoxidation
RemarksStart Finish Mass decrease Start Finish** Mass increase
°C w/% °C w/%
Alloy A 248 (247)* 447 (446) 1.40 (1.38) 447 (446) 650 0.53 (0.48) –
Alloy B 249 (194) 426 (477) 1.53 (1.93) 426 (477) 650 0.25 (0.14) Early start and late finish
of dewaxing in ArAlloy C 249 (198) 426 (463) 1.38 (1.73) 426 (463) 650 1.38 (1.92)
* …. Results of the experiments performed in stationary Ar are given in parenthesis
** …. All experiments finished at 650 °C
M. TORKAR et al.: MICROSTRUCTURE CHARACTERISTICS OF THE MODEL SPRING STEEL 51CrV4
MICROSTRUCTURE CHARACTERISTICS OF THE
MODEL SPRING STEEL 51CrV4
ZNA^ILNOSTI MIKROSTRUKTURE MODELNEGA VZMETNEGA
JEKLA 51CrV4
Matja` Torkar1, Franc Tehovnik1, Bo{tjan Arh1, Monika Jenko1, Bo`idar [arler2,
@arko Raji}3
1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2Laboratory for Multiphase Processes, University of Nova Gorica, Slovenia
3CIMOS, d. d., Koper, Slovenia
matjaz.torkar@imt.si
Prejem rokopisa – received: 2014-01-28; sprejem za objavo – accepted for publication: 2014-04-03
There are many data on solidification processes in the literature, but there are no relevant comparisons performed for the
formation of an as-cast microstructure, the formation of gas porosity, and the influence of different cooling rates on the
solidification of the molten, non-killed spring steel 51CrV4. During a relatively rapid solidification the majority of the dissolved
gases remain entrapped and due to a solid shell being formed they cannot escape from the steel. The result is the formation of
gas porosity and a shrinkage cavity. Here we present the results of an investigation of the appearance and distribution of gas
bubbles and a shrinkage cavity as well as the characterization of an as-cast microstructure in ingots of model spring steel, cast
and cooled in different ways.
Keywords: spring steel, solidification, gas porosity, shrinkage cavity, SDAS, segregations
V literaturi je mnogo podatkov o procesu strjevanja, vendar pa ni podatkov o primerjavi med nastankom strjevalne strukture in
nastankom poroznosti med strjevanjem nepomirjenega vzmetnega jekla 51CrV4. Pri relativno hitrem strjevanju ve~ina
raztopljenih plinov ostane ujetih v jeklu in zaradi nastanka skorje ne morejo pobegniti iz njega. Rezultat tega je nastanek plinske
poroznosti in lunkerja. Predstavljeni so rezultati preiskav ter pojava in razporeditve plinskih mehur~kov in lunkerja, kot tudi
karakterizacija mikrostrukture v ingotih vzmetnega jekla, ulitih in ohlajenih na razli~ne na~ine.
Klju~ne besede: vzmetno jeklo, strjevanje, plinska poroznost, lunker, SDAS, izceje
1 INTRODUCTION
The typical stages of the solidification process in
alloys can be described with the following stages:
nucleation of a solid (stable nucleus); growth, which can
be cellular or dendritic; solute distribution; and further
growth, which proceeds by the movement of the inter-
face and makes up the final as-solidified structure. The
accumulation of the solute and the heat ahead of the
interface can lead to conditions in which the liquid in
front of the solidification front is supercooled. Under
specific conditions the solidification becomes dendritic.
A characteristic tree-like structure of crystals grow-
ing along an energetically favourable crystallographic
direction is called a dendrite in metallurgy. The form and
the size of dendrites have a large influence on material
properties.1
The requirement for dendrite formation is that the
molten material is undercooled, below the freezing point
of the solid. After the nucleation a spherical solid nu-
cleus starts to grow as a sphere, later it becomes unstable
and the solid shape depends on the preferred growth
directions of the crystal. The anisotropy of the surface
energy of the solid-liquid interface influences the growth
direction. Besides that, the solidification rate influences
the secondary dendrite arm spacing (SDAS), which can
also be calculated using equation (1).2
The SDAS () is a function of the cooling rate and
can be written as:
 = ⋅ −B T n (1)
where B and n are experimental constant parameters
with values of 319.4 and 0.378, respectively,3 and T is
the cooling rate (°C/s) for this type of steel.
The SDAS () can also be determined from a micro-
graph.4
During their grow dendrites bump into one another,
until their growth is restricted by the other dendrites
around them. At that point, the dendrite has reached its
maximum size and it represents a grain. The space that
exists between the grains is referred to as a grain
boundary. If molten steel is cooled slowly, the result is
the formation of a large grain size. However, when the
steel is cooled quickly, the number of dendrites increases
and a smaller grain size is the result.5 All these para-
meters influence the microstructure development and the
properties of the spring steel.6
Solidification proceeds at various rates. For that
reason the microstructure is not homogeneous and vari-
ations in the composition appear as segregation. Segre-
gation occurs because the diffusion in the solid phase is
Materiali in tehnologije / Materials and technology 48 (2014) 4, 537–543 537
UDK 669.14:536.421.4 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)537(2014)
too slow to achieve a uniform equilibrium structure, as
predicted in the phase diagram. Segregation is classified,
according to its scale, as macro-segregation or micro-se-
gregation. This macro-segregation occurs on the scale of
the grains or the entire casting and can be observed visu-
ally. It arises from a large-scale fluid caused by forced
natural and solutal convection. The formation of segre-
gation requires the transport of a solute-poor liquid and
solid phases during the solidification over distances
much larger than the dendrite arm spacing. The interden-
dritic flow of the liquid due to solidification shrinkage
and changes in the liquid density are an unavoidable
cause. These density changes can be caused by tempe-
rature changes or by changes in the composition of the
liquid.7–10
It was revealed that the cooling rate also has an
influence on the nucleation temperature and the width of
the solidification interval.11
Micro-segregation usually arises between the dendri-
te arms. Due changes in the chemical composition, the
strength and the ductility are lower in the transverse di-
rection compared to the longitudinal axis of the dendri-
tes.12
An important piece of data on the solidification struc-
ture is the coefficient of segregation k. It is determined as
the rate of maximum and minimum concentration among
the secondary dendrite arms (k = Cmax/Cmin) 13 or on the
scale of the crystal grains.
Shrinkage cavities and voids are formed during the
solidification due the volumetric shrinkage of the metal.
This type of defect occurs because most metals contract
by 2–6 % in terms of volume when the liquid transforms
to a solid.12 During hot rolling most of the voids are
welded without any influence on the mechanical proper-
ties. In the case of gas bubbles (hydrogen), especially if
they are close to the surface, they can form blisters on
the surface during the cold rolling of sheets.14
Large shrinkage cavities can be minimized by the
careful control of the heat transfer during the solidifica-
tion process. For example, attaching a liquid metal
reservoir, or riser, to a sand casting or to an ingot head,
encourages shrinkage to occur in the riser, which feeds
metal into the casting.
The molten metal dissolves and contains gases. The
gases originate from the material, from the atmosphere
or from reactions between the molten metal and the
mould material. Since a liquid metal has a much higher
solubility for gases than a solid metal, the gases are
expelled during the solidification of steel. If they cannot
escape they may form various defects in the material,
e.g., porosity in the metal casting15.
The maximum solubility of nitrogen in liquid iron is
approximately 450 · 10–6, and less than 10 · 10–6 at ambi-
ent temperature (Figure 1).16 The presence of the alloy-
ing elements in the liquid iron or steel affects the
solubility of nitrogen. An important feature is that the
presence of dissolved sulfur and oxygen limit the absorp-
tion of nitrogen because they are surface-active elements.
This is exploited during steelmaking to avoid excessive
nitrogen pickup, particularly during tapping.16
Shrinkage porosity is represented by small voids in a
casting due to solidification shrinkage. Gas porosity is
caused by gas bubbles that are evolved during the solidi-
fication and become trapped to form small, smooth,
round voids or pinholes inside the casting.17
The bubbles can also be formed by a chemical reac-
tion between the dissolved oxygen and the carbon in
steel, to create carbon-monoxide bubbles. The carbon-
oxygen reaction in Al-free steel can lead to blowholes,
but also to pinholes in the cast material. Carbon and
oxygen become enriched during the solidification in the
interdendritic melt (as all elements), and the pressures of
the carbon monoxide and carbon dioxide increase
according to the reactions: [C] + [O] = {CO} and [C] +
2[O] = {CO2}.
After the CO and CO2 bubbles are initiated they start
to grow with the advancing solid shell and form elon-
gated blowholes.
During rapid solidification a majority of the dis-
solved gas remains entrapped, as due to the solid shell
they cannot escape from the steel and this causes gas
bubbles below the solidified shell during the solidifica-
tion process. In the literature no comparisons are made
from among the formation of gas porosity and the pro-
perties of the as-cast microstructure during the solidifi-
cation of the molten spring steel 51CrV4 (1.8159) at
different cooling rates. The standardized chemical com-
position of this type of steel is in mass fractions w(C) =
0.47–0.55 %, w(Si, max ) = 0.4 %, w(Mn) = 0.7–1.1 %,
w(P, max) = 0.025 % , w(S, max)) = 0.025 % , w(Cr) =
0.9–1.2 %, w(V) = 0.1–0.25 %.
The aim of the present investigation was to study the
influence of the cooling rate of the not killed model
spring steel on the appearance and distribution of the gas
bubbles and the shrinkage cavity as well as the characte-
rization of the as-cast microstructure.
M. TORKAR et al.: MICROSTRUCTURE CHARACTERISTICS OF THE MODEL SPRING STEEL 51CrV4
538 Materiali in tehnologije / Materials and technology 48 (2014) 4, 537–543
Figure 1: Solubility of nitrogen in iron for the temperatures 600–2000
°C16
Slika 1: Topnost du{ika v `elezu pri temperaturah 600–2000 °C16
2 EXPERIMENTAL
The base material was prepared by the remelting of
C, Si, Mn, Cr, V spring steel 51CrV4 in a Heraeus induc-
tion melting furnace. The final chemical composition of
the experimental ingots was: w(C) = 0.53 %, w(Si) =
0.31 %, w(Mn) = 0,96 %, w(Cr) = 1.11 %, w(Mo) = 0.06
%, w(Ni) = 0.17 % and w(V) = 0.16 %. The melt was not
deoxidized with Al prior to casting. The temperature of
the melt in the melting furnace was 1580 °C. The melt
was poured from the furnace into the preheated ceramic
pot and then the melt with a temperature of 1530 °C was
cast into an iron mould and sand form, to obtain different
solidification cooling rates. The cross-section of the
moulds and the sand cavity was 60 mm × 60 mm.
Special iron moulds were taken for the casting. The
moulds consisted of two parts, separated by the diagonal
cross-section, enabling the rapid opening and release of
the ingot with a solidified shell. The thickness of the
mould wall was 25 mm. The moulds were protected with
a zirconium-oxide-based coating and heated prior to cast-
ing at 150 °C. The first sample, cast into an iron mould,
was taken immediately after casting from the mould and
cooled down in water. The second sample was cast into
iron mould and cooled down in the mould. The third
sample was cast and cooled down in a sand form. Three
ingots were produced from the same melt, using three
different cooling rates.
To characterize the soundness and the microstructure
of the as-cast ingots, a 1-cm-thick plate was cut with a
water-jet cutter, in a longitudinal direction, from all three
ingots. Samples for metallography were cut from the
plate in the middle of the ingots. The samples were pre-
pared by a standard metallographic method and etched in
Nital for observation in a light microscope.
Also, the individual micrographs were taken in 5 mm
steps from the surface to the center of the ingots. The
microstructure was observed using a Nikon Microphot
FXA light microscope. The secondary dendrite arm
spacing was determined from the etched samples. The
microstructure was also observed with a JSM-6500F
scanning electron microscope (SEM) equipped with a
field-emission source of electrons and analysed with an
INCA ENERGY Oxford Instruments Energy-Dispersive
Spectroscopy (EDS).
3 RESULTS WITH DISCUSSION
The ingots solidified with three cooling rates were
manufactured using the casting protocols as follows: a)
cast in an iron mould and cooled in water, b) cast and
cooled in an iron mould, c) cast and cooled in a sand
form.
The ingot cooled in water was cut in a longitudinal
direction, as presented in Figure 2. The cut was not very
clear as the internal gas bubbles and shrinkage cavity
(Figure 3) disturbed the high-pressure water-jet cutting
process. A similar effect was also observed for the other
two ingots, as presented in Figures 4 and 5.
The dendrites in the microstructure were observed
only in the ingot cooled in water. The microstructure was
martensitic. Normal crystal grains and no dendritic mi-
crostructure were observed in the other two ingots,
where the microstructure was pearlite with a different
inter lamellar spacing, dependent on the cooling rate. In
the samples cooled in the mould or in sand some ferrite
on the grain boundaries was observed near the surface of
the ingot.
From the comparison of the longitudinal cross-sec-
tions of the ingots it is evident that there is a different
formation and position of the gas bubbles and shrinkage
porosity, both dependent on the cooling rate and the
movement of the solidification front. In the ingot cooled
M. TORKAR et al.: MICROSTRUCTURE CHARACTERISTICS OF THE MODEL SPRING STEEL 51CrV4
Materiali in tehnologije / Materials and technology 48 (2014) 4, 537–543 539
Figure 2: Longitudinal section of ingot cooled in water and cut with a
high-pressure water jet cutter. The presence of gas bubbles disturbed
the cutting process. The other two ingots were cut in the same way.
Slika 2: Vzdol`ni prerez ingota, ohlajenega v vodi in razrezanega na
rezalniku z visokotla~nim vodnim curkom. Tudi druga dva ingota sta
bila razrezana na enak na~in.
Figure 3: Longitudinal section of ingot, cooled in water. Observed are
the gas bubbles and the closed shrinkage cavity in the head and the
longitudinal central crack due to rapid cooling after solidification.
Evident is the thickness of the solidified shell prior to cooling in the
water. All the gas bubbles are below this shell.
Slika 3: Vzdol`ni prerez ingota, ohlajenega v vodi, s plinskimi me-
hur~ki in zaprto poroznostjo v glavi ter vzdol`na razpoka v sredini
zaradi hitrega ohlajanja po strjevanju. Vidna je debelina strjene skorje
pred ohlajanjem v vodi. Vsi plinski mehur~ki so pod to skorjo.
in water larger gas bubbles and a large shrinkage cavity
was observed, as well as a long crack in the middle of
the ingot’s longitudinal direction. Evident is also the
thickness of the rapidly solidified shell that pushed the
bubbles away (Figure 3)
The relatively sound material, concerning gas
bubbles, is in the ingot cooled in the iron mould. Some
gas bubbles are in the central part of the ingot, closer to
the head. The reason is the solidification front that
moved from the sides and from the bottom toward the
head of ingot, where the gas bubbles and the shrinkage
porosity are collected (Figure 4).
Almost no shrinkage porosity is observed in the ingot
cast and cooled in the sand form (Figure 5). The gas
bubbles (pinholes) are present along the whole ingot and
closer to the surface, as the movement of the solidifica-
tion front was slower. Some individual pinholes are also
present in the central part of ingot and more numerous
gas bubbles are closer to the head of the ingot.
In general, the gas bubbles and shrinkage porosity for
the present steel are more expressed than the steel melt
that was not killed with Al or other desoxidant. For that
reason the gas content in the melt was higher and as a
consequence also the number of bubbles was higher.
It is evident that the dissolved gases (Figure 1) did
not have enough time to separate completely from the
melt, due to the relatively high cooling rate, and the
small cross-section at all three ingots. The concentration
of gas bubbles depends on the cooling rate. A large
closed shrinkage cavity in the head and larger gas
bubbles below the surface were observed in the ingot
with the fastest solidification rate. In the ingot cooled in
the mould the bubbles and the shrinkage cavity moved
toward the head of the ingot.
For the ingot cooled in the sand form most of the gas
bubbles were observed closer to the surface and no
shrinkage cavity was observed in the head of the ingot.
This is also evidence that the melt was not killed with
aluminium or silicon.
The samples for metallography were prepared from
plates cut with a water jet from ingots. Light microscopy
revealed that the ingot with the fastest cooling in water
has a dendritic microstructure with martensite (Figure 6)
and visible interdendritic segregations.
The other two ingots, with the slower cooling rate,
have a perlitic microstructure with crystal grains of diffe-
rent size and also different interlamelar spacings of ce-
mentite lamelas in pearlite (Figure 7). The slower cool-
ing rate enables a pearlite transformation in the material.
Due to the relatively slow solidification and cooling
in the iron mould and in the sand form, the segregations
are clearly visible for the ingot cooled in water. The den-
drite arm spacing (Table 1) was determined either by
measurements of the segregation distance or calculated
using equation (1). The cooling rate was also calculated
using equation (1).
Table 1: Secondary dendrite arm spacing , measured and determined
with equation (1), for three cooling rates of the ingots
Tabela 1: Oddaljenost sekundarnih dendritnih vej , izmerjena in
dolo~ena z ena~bo (1), za tri hitrosti ohlajanja ingotov
Parameter
Cooling of ingot
Water
(μm)
Iron mold
(μm)
Sand form
(μm)
Measured SDAS, /μm 50 150 170
Calculated SDAS, /μm 52 182 210
Calculated T/(°C/s) 121.7 4.4 3.0
 – secondary dendrite arms spacings (SDAS)
T – cooling rate
M. TORKAR et al.: MICROSTRUCTURE CHARACTERISTICS OF THE MODEL SPRING STEEL 51CrV4
540 Materiali in tehnologije / Materials and technology 48 (2014) 4, 537–543
Figure 4: Longitudinal section of ingot cooled in iron mould. Gas
bubbles and small closed shrinkage cavity moved toward the head.
The solidification front moves toward the head, so the gas bubbles are
concentrated below the head, where the last melt solidified.
Slika 4: Vzdol`ni prerez ingota, ohlajenega v `elezni kokili. Plinski
mehur~ki in majhna zaprta poroznost so pomaknjeni bli`e h glavi.
Fronta strjevanja se je pomikala proti glavi, zato so plinski mehur~ki
koncentrirani pod glavo ingota, kjer se je strdila zadnja talina.
Figure 5: Longitudinal section of ingot, cast and cooled in sand form.
Due to the lower solidification rate, the growth of the shell was slower
and the bubbles are closer to the surface. No shrinkage cavity is
observed.
Slika 5: Vzdol`ni prerez ingota, ulitega in ohlajenega v pesku. Zaradi
po~asnej{ega ohlajanja je bila po~asnej{a tudi rast skorje in iz taline
izlo~eni mehur~ki so bli`je povr{ini. Lunkerja ni opaziti.
M. TORKAR et al.: MICROSTRUCTURE CHARACTERISTICS OF THE MODEL SPRING STEEL 51CrV4
Materiali in tehnologije / Materials and technology 48 (2014) 4, 537–543 541
Distan. from
the surface
a) Cast into iron mold and cooled in
water b) Cast and cooled in iron mold c) Cast and cooled in sand form
0 mm
5 mm
10 mm
15 mm
20 mm
25 mm
Figure 6: Comparison of as-cast microstructure of ingots at given distances from the surface. Light microscope (magnification 100-times).
Slika 6: Primerjava mikrostrukture ulitih ingotov pri dani razdalji od povr{ine. Svetlobni mikroskop (pove~ava 100-kratna).
From the results given in Table 1 it is clear that the
highest cooling rate was in the ingot cooled in water. As
can be seen from Figures 6a and 7a the obtained mi-
crostructure under the light microscope was martensite
for the ingot cooled in water. On the other hand, the
cooling rates of other two ingots, solidified in the iron
mould and in the sand, are smaller, which is also con-
firmed by the microstructure (Figures 6b, 6c, 7b and
7c). The microstructures of the ingot solidified in the
iron mould and in the sand are pearlite. The reason is in
the formation of a gap in the iron mould that prevents
faster cooling. The lowest cooling rate is, as expected, in
the sand form. The ratio of the cooling rates, compared
to the cooling rate in the sand are: 40.56 : 1.46 : 1.00. It
means the cooling rate of the ingot in the water was
40.56 times faster compared with the cooling rate of the
ingot in the sand form. The cooling rate also influences
the SDAS. The secondary dendrite arm spacing, compa-
red to the cooling rates in water, are 1.00 : 3.5 : 4.03. The
SDAS is 4.03 times larger in the ingot that was cast and
cooled in the sand form, compared to the ingot that was
cooled in water.
The measurements of the Cr segregations in all three
ingots revealed larger differences in the Cr and V con-
centrations only in the material cooled in water (Figure
8 and Table 2) and less expressed segregations of Cr in
ingots cooled in the mould or in the sand (Tables 3 and
4), because of the lower cooling rate.
A less expressed segregation of V was observed only
in an ingot cooled in water. In the other two ingots the
segregation of V was not detected. For that reason only a
comparison of the Cr segregations in all three ingots is
presented (Tables 2, 3 and 4).
Table 2: Segregations of Cr in the ingot cooled in water
Tabela 2: Izceje Cr v ingotu, ohlajenem v vodi
w(Cr)/% k = Cmax/Cmin
Centre of Dendrite 1 1.42
Between dendrite 1/2 1.83 1.28
Centre of Dendrite 2 1.00
Between dendrite 2/3 1.34 1.34
Centre of dendrite 3 1.42
Between dendrite 3/4 1.75 1.23
Table 3: Segregations of Cr in the ingot cooled in the mould
Tabela 3: Izceje Cr v ingotu, ohlajenem v kokili
w(Cr)/% k = Cmax/Cmin
Centre of Dendrite 1 1.13
Between dendrite 1/2 1.20 1.06
Centre of Dendrite 2 1.20
Between dendrite 2/3 1.51 1.25
Centre of dendrite 3 1.27
Between dendrite 3/4 1.33 1.04
Table 4: Segregations of Cr in the ingot cooled in sand
Tabela 4: Izceje kroma v ingotu, ohlajenem v pesku
w(Cr)/% k = Cmax/Cmin
Centre of Dendrite 1 1.20
Between dendrite 1/2 1.22 1.01
Centre of Dendrite 2 1.20
Between dendrite 2/3 1.28 1.06
Centre of dendrite 3 1.22
Between dendrite 3/4 1.26 1.03
The experiments revealed that the increased cooling
rate decreases the secondary dendrite arm spacing,
M. TORKAR et al.: MICROSTRUCTURE CHARACTERISTICS OF THE MODEL SPRING STEEL 51CrV4
542 Materiali in tehnologije / Materials and technology 48 (2014) 4, 537–543
Figure 8: Segregation of chromium and vanadium along the line for
ingot cooled in water (SEM)
Slika 8: Izceja kroma in vanadija vzdol` ~rte v ingotu, ohlajenem v
vodi (SEM)
Figure 7: Microstructure of ingots: a) cooled in water, b) cooled in iron mould and c) cooled in sand, (SEM)
Slika 7: Mikrostruktura ingotov: a) ohlajanje v vodi, b) ohlajanje v `elezni kokili in c) ohlajanje v formi iz peska (SEM)
increases the intensity of the segregation of Cr and
influences the formation and distribution of gas bubbles
and the shrinkage cavity in as-cast material.
For the further processing of the steel the gas bubbles
may not be so detrimental because they can be welded
during a hot deformation.
To decrease the quantity of gas bubbles, the steel
melt needs to be killed more intensively, prior to casting.
4 CONCLUSIONS
Based on the investigations on as-cast 60 mm × 60
mm not-killed spring steel 51CrV4 ingots, the following
conclusions can be drawn.
• A comparison of the as-cast microstructures of ex-
perimental cast ingots revealed the dendritic structure
only in an ingot cooled in water.
• The ingot cooled in water had a martensitic micro-
structure.
• The ingots cooled in an iron mould or in a sand form
have a similar microstructure with pearlite grains,
which were coarser for the cooling in the sand form.
No dendrites were observed. The slower cooling rate
enabled the formation of pearlite with ferrite on some
grain boundaries.
• The increased cooling rate decreases the secondary
dendrite arm spacing and also influences the forma-
tion, as well as the layout, of the gas bubbles and the
shrinkage cavity in as-cast spring steel.
• A slower cooling rate caused the formation of gas
bubbles closer to the surface of the as-cast ingot.
• In the ingot cooled in water larger gas bubbles were
observed, distributed across the whole ingot. A large
shrinkage cavity was present, as well as longitudinal
cracks in the middle of the ingot.
• The increased cooling rate also increases the segre-
gation of chromium.
• To decrease the quantity of gas bubbles, the steel
needs to be aluminium or silicon killed more inten-
sively, prior to the casting.
Acknowledgement
Authors wish to thank to Slovenian Research Agency
for the financial support in the frame of projects No.
L2-3651 and P2-0132.
5 REFERENCES
1 W. Kurz, D. J. Fisher, Fundamentals of Solidification, 4th ed., Trans
Tech Publ. Ltd, Switzerland 1998
2 T. W. Clyne, W. Kurz: Metall. Trans. A, 12 (1981), 965–71
3 http://steeluniversity.org
4 J. Campbell, Castings, Butterworth-Heinemann, 2003, 270
5 H. Fredriksson Hasse, U. Akerlind, Solidification and crystallization
processing in metals and alloys, John Willey and sons, 2012
6 B. Podgornik, V. Leskov{ek, M. Godec, B. Sen~i~, Microstructure
refinement and its effect on properties of spring steel, Materials
Science and Engineering A, 599 (2014), 81–86
7 M. C. Flemings, G. E. Nereo, Trans. TMS-AIME, 239 (1967),
1449–61
8 R. Mehrabian, M. Keane, M. C. Flemings, Metall. Trans., 1 (1970),
1209–20
9 T. Fuji, D. R. Poirier, M. C. Flemings, Metall. Trans. B, 10 (1979),
331–39
10 Y. M. Won, B. G. Thomas, Simple Model of Microsegregation dur-
ing Solidification of Steels, Metallurgical and Materials Transactions
A, 32 (2001), 1755
11 D. Steiner Petrovi~, M. Pirnat, G. Klan~nik, P. Mrvar, J. Medved,
The effect of cooling rate on the solidification and microstructure
evolution in duplex stainless steel, J. Therm. Anal. Calorim., 109
(2012) 3, 1185–1191
12 K. O. Yu (ed.), Modeling for Casting and Solidification Processing,
1st ed., CRC Press, 2001, 12
13 M. Torkar, B. [u{tar{i~, Preiskava vodno atomiziranega prahu iz
zlitine Nimonic 80 A, Kovine zlitine tehnologije, 26 (1992) 1–2,
100–103
14 H. J. Grabke, Supersaturation of iron with nitrogen, hydrogen or
carbon and the consequences, Mater. Tehnol., 38 (2004) 5, 211–222
15 P. S. Wei, C. C. Huang, Z. P. Wang, K. Y. Chen, C. H. Lin, Growths
of bubble/pore sizes in solid during solidification – an in situ measu-
rement and analysis, Journal of Crystal Growth, 270 (2004) 3–4,
662–673
16 http://www.keytometals.com
17 R. A. Stoeh, Modeling the influence of fluid flow on the develop-
ment of porosity and microstructure in castings, Canadian
Metallurgical Quarterly, 37 (1998) 3–4, 179–184
M. TORKAR et al.: MICROSTRUCTURE CHARACTERISTICS OF THE MODEL SPRING STEEL 51CrV4
Materiali in tehnologije / Materials and technology 48 (2014) 4, 537–543 543

F. VODOPIVEC et al.: EFFECT OF CREEP STRAIN ON CREEP RATE IN THE TEMPERATURE RANGE 550–640 °C
EFFECT OF CREEP STRAIN ON CREEP RATE IN THE
TEMPERATURE RANGE 550–640 °C
VPLIV DEFORMACIJE NA HITROST LEZENJA PRI
TEMPERATURAH OD 550 °C DO 640 °C
Borut @u`ek, Franc Vodopivec, Monika Jenko, Bojan Podgornik
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
borut.zuzek@imt.si, franc.vodopivec@imt.si
Prejem rokopisa – received: 2014-03-06; sprejem za objavo – accepted for publication: 2014-04-03
Experimental and calculated creep rates were examined for a high-chromium, creep-resistant steel samples quenched and
tempered for 2 h and 400 h at 800 °C. Creep testing and annealing were carried out at 550 °C to 640 °C. It was found that the
difference in the creep rates due to the particle coarsening and dissolution was independent of the temperature, while the ratio of
the experimental and calculated creep rates increased with the temperature of the creep tests. The results suggest that an increase
in the creep rate is related to the process of vacancy climb in parallel with the increase in the diffusion rate of the iron in the
ferrite.
Keywords: creep-resistant steel, M23C6 particle coarsening and dissolution, experimental and calculated creep rates, effect of
temperature
Hitrost lezenja je bila eksperimentalno dolo~ena in izra~unana za jeklo, odporno proti lezenju, z visoko vsebnostjo kroma pri
vzorcih, ki so bili kaljeni in popu{~ani 2 h oz. 400 h pri 800 °C. Stourni preizkusi lezenja in `arjenja so bili izvr{eni pri 550 °C
do 640 °C. Razlika v hitrosti lezenja zaradi raztapljanja oziroma rasti izlo~kov je bila neodvisna od temperature, medtem ko je
razlika med eksperimentalno in izra~unano hitrostjo lezenja ve~ja pri vi{ji temperaturi. Rezultati meritev in izra~unov ka`ejo, da
je pove~anje hitrosti lezenja povezano z ve~jo hitrostjo plezanja vrzeli, ki je vzporedno s pove~anjem hitrosti difuzije `eleza v
feritu.
Klju~ne besede: jeklo, odporno proti lezenju, rast in raztapljanje izlo~kov M23C6, eksperimentalna in izra~unana hitrost lezenja,
vpliv temperature
1 INTRODUCTION
The microstructure of creep-resistant steels consists
of a ferrite matrix with different elements, mostly chro-
mium in solid solution, and particles, mostly carbides
and carbonitrides, with sizes in the range from a few nm
to about 1 · 103 nm. Those particles are unshearable
obstacles to the gliding of the matrix mobile
dislocations. For a particle size (dp) greater than the
spacing (mutual distance) of the mobile dislocations, the
particles intercept the mobile dislocations with a
probability p  1, while for smaller particles the
intercept probability is p  1.1 So far, it has not been
determined whether the unshearability is related to a
minimum particle size, i.e., the number of carbide
molecules in the particle. For example, the lattice
parameter (lp) of the cubic carbide Cr23C6 is lp = 1.016
nm,2 and the number of molecules in the particles is Npm
 dp3/lp. The number increases rapidly with the particle
size: it is Npm = 8 for particles with dp = 2 nm and Nm =
125 for dp = 5 nm. With isothermal tempering, the ave-
rage particle size increases and their number decreases,
as a result of the duality of the particle-coarsening
process: the growth of coarser particles, the shrinking of
particles with an intermediate size and even the
dissolution of smaller particles.1 This means that by
tempering, the average size of the particles increases,
and it increases more rapidly during a creep test, as the
iron self-diffusion and the chromium diffusion in ferrite
are enhanced by tensile stress.3,4
For creep-resistant steels, the improvement of Horn-
bogen’s creep equation5,6 for a uniform distribution of
particles was proposed, with the substitution of the
particle spacing  by ( – d) as a measure of the
dislocation climb length increase by overcoming of
particles inclusion of a constant representing the increase
of the stress exponent n from n = 2 to n = 3.65:

( )

  =
⋅ ⋅ ⋅ ⋅ ⋅ −
⋅ ⋅
k b D d
k T G
n2
b
(1)
where  is the acting stress ( = 170 MPa),  is the den-
sity of the mobile dislocations ( = 0.978 1014 m–2), D is
the diffusion coefficient, d is the average particle size, p
is the average particle spacing, b is the Burgers vector
(b = 2.5 · 10–10 m), kb is the Boltzmann constant (kb =
1.38 · 10–23 J K–1), T is the temperature, G is the shear
modulus7 and k = 3.65/2 = 4.78 · 103.
By substituting the molar content of carbide M23C6
with the molar content of the chromium in solution in
ferrite as a parameter in the LSW (Livshitz-Slyozov-
Wagner) equation, a reasonable agreement was obtained
for the calculated (Eq. 2) and experimental coarsening
rate of the M23C6 particles at 800 °C. Also, Eq. 3 was
deduced for the calculation of the coarsening rate at
lower temperature:8
Materiali in tehnologije / Materials and technology 48 (2014) 4, 545–548 545
UDK 669.14.018.44:539.37:548.4 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 48(4)545(2014)
d d
S D
k T
tt
b
3
0
3
8
9
− =
⋅ ⋅ ⋅ ⋅
⋅ ⋅
⎛
⎝
⎜
⎞
⎠
⎟ ⋅
 
(2)
k k
D
D TTc c
Cr, T
Cr,1073
=
⋅
⋅
⎛
⎝
⎜
⎞
⎠
⎟
,1073
1073
(3)
where dt is the particle size at the tempering time t, d0 is
the initial particles size, D is the chromium diffusion
rate, T is the tempering temperature, S is the atomic con-
tent of chromium in solid solution in the ferrite,  is the
carbide particle matrix interfacial energy ( = 0.37
J m–2),  is the volume of diffusing atoms, D is the
chromium diffusion rate (D = D0 exp(–Q/RT)) with D0
= 3.7 · 10–3 m2 s–1 and Q is the activation energy (Q =
267 kJ mol–1),9 kb is the Boltzmann constant and kc1073
is the experimental coarsening rate at 800 °C (kc1073 =
2.94 · 10–27 m3 s–1).
In this work, the effect of the creep strain on the
change of particle size and spacing is investigated with a
constant creep stress and time.
2 CALCULATIONS
The creep rate and creep fracture of creep-resistant
steels depend of a large number of variables,10 because
with the creep temperature, the initial microstructure of
the ferrite and the uniform distribution of particles is
changed due to particle coarsening. With this coarsening,
several processes occur, i.e., the growth of the particles’
average size and spacing, the dissolution of small
particles, the decrease in the number of particles and of
grain-boundary stringers.10 The change in the particles
size and spacing is calculated for 100 h of tempering at
550–640 °C in steps of 30 °C using Eqs. 2 and 3.
Assuming the particles to be spheres, the number of
particles Np was calculated by applying the series:1
f d d d d N dn
n
= ⋅ + + + + =∑π π
6 61
3
2
3
3
3 3
1
3... p a
and N
f
dp a
=
6
3π
(4)
where f is the volume of particles, Np is the number of
particles, d1, d2, d3,…. dn is the decreasing size of the
particles and da is the average particles size. The volume
of particles in the investigated steel was f = 0.047 10 and
it was determined that the particles were carbide M23C6
(Cr18Fe3Mo2C6).11,12
The particles dissolution velocity was deduced as:1
d d kd td
3
0
3− = t
d d
kd
d
c
=
−3 0
3
v
d
td
d
d
=
⎛
⎝
⎜
⎞
⎠
⎟
3 1 3/
(5)
where dd is the size of the dissolving particles, kc is the
isothermal coarsening rate (the term in parenthesis in
Eq. 2), td is the dissolution time, d0 = 2 nm and, vd is the
dissolution velocity.
The volume of the particles increases with dd3, and
for this reason the parameter d0 = 2 nm was omitted from
the calculations of the dissolution velocity.
The initial average particle size dia was deduced for
the tempering of specimens quenched in oil as:1
(dia)
3 = kcT t (6)
where kc is the coarsening rate at 1073 K (kc800 °C = 2.94
· 1027 m3 s–1) and t is the tempering time at 1073 K, and
dia = 148 nm was deduced.8 The volume of carbide par-
ticles Cr18Fe3Mo2C6 f = 0.047 was deduced from the
content of chromium in the investigated steel.13 The cal-
culated average particle size agrees well with the avera-
ge size d = 140 nm assessed from micrographs.13
The increase of the particle size with a tempering
time of 100 h and the creep test temperature is deduced
as:
daT = (kcT/kc1073) d1073 (7)
where daT is the increase of the average particle size
with 100 h of tempering at temperature T = (823, 853,
883 and 913) K, kcT is the coarsening rate deduced from
Eq. 3, kc1073 is the coarsening rate at 1073 K (kc1073 =
2.94 · 10–27 m3 s–1) and d1073 is the increase of parti-
cles with an initial size of 148 nm after 100 h tempering
at 1073 K. The number of particles N1p = 2.11 · 1019 m–3
was then deduced from Eq. 4 for f = 0.047.14
Then, the average particle spacing  for the particles
coarsening at the test temperature was calculated as:6
 = 4 d/f1/3 (8)
With tempering rsp. creep test temperature, a signi-
ficant number of particles with a size in the lower part of
size distribution is dissolved.1 The size of the dissolved
particles ddm was first deduced from Eq. 5, and then the
part of particle with size dd  ddm determined for M23C6
of average size da = 157 nm.1 Following that the number
of undissolved particles after 100 h of tempering at the
creep temperature was deduced to be NpT = Nip (10–2Nnd).
The changes of the average particle size du and
spacing u for the undissolved particles with 100 h at the
creep test temperature were calculated as:
d
f
N T
u
p
=
⋅ ⋅
⋅
⎛
⎝
⎜
⎞
⎠
⎟6 10
27
1 3
π
/
and  u u
p
= ⋅
⋅⎛
⎝
⎜
⎞
⎠
⎟k
N T
6 10 27
1 3/
(9)
with ku = 1.585 constant characteristic for the stochastic
distribution of cube particle. With an equal average
particle size and f = 0.047 the difference of in the
particle spacing deduced from Eq. 8 and 9 is about 1 %.
With the change of the creep temperature in Eq. 1,
the parameters T, DFe, , d and G are changed, and it is
possible to calculate the change of creep rate with the
known iron self-diffusion rate DFet and chromium DCrt
diffusion rate by tempering, particles size and spacing
and shear modulus. In the calculation it should be kept in
mind that by creep test iron self-diffusion rate is
increased and the chromium diffusion rate, determining
the change of particles size and spacing, changes as well.
The iron self-diffusion and chromium diffusion rates
were calculated from the data in8 and the shear modulus
deduced from.9
F. VODOPIVEC et al.: EFFECT OF CREEP STRAIN ON CREEP RATE IN THE TEMPERATURE RANGE 550–640 °C
546 Materiali in tehnologije / Materials and technology 48 (2014) 4, 545–548
3 RESULTS
In3,4 it is stated that the number of vacancies in ferrite
is increased by tensile stress and the iron and chromium
diffusion rate are increased. It is reasonable to presume
that with an increased content of vacancies, the climb
velocity is also increased. Eq. 1 indicates that with a
constant creep stress, the increase of the creep rate is
proportional to changes of the iron diffusion rate, of the
parameter ( – d) and of the shear modulus G.7 Assum-
ing that the effect of temperature and creep is equal for
iron and chromium diffusion, the change of particle size
and spacing was calculated.
The results of the calculations listed in Table 1 indi-
cate that the increase of coarsening rate due to the
growth of the average particle size is much lower than
that due to the dissolution of small particles. The greatest
increase of creep rate due to particle growth was  =
0.3 · 10–8 s–1, about 0.4 %, which was obtained at the
highest creep test temperature of 913 K (640 °C). The
increase of the creep rate due to the dissolution of small
particles was significantly greater and it increased with
the creep test temperature. At 913 K (640 °C) it was 
= 4.3 · 10–8 s–1 and by 6.1 %, about 15 times, greater than
that due to the growth of the average particle size.
The difference of the experimental creep rate and that
due to the particle dissolution increased with the test
temperature and was  = 0.47 · 10–8 at 823 K and  =
645.4 · 10–8 at 913 K.
In Figure 1 the experimental creep rate, as well as
the iron diffusion rate by tempering and the apparent iron
diffusion rate by creep are depicted as the dependence of
creep test temperature. As expected for diffusion-con-
trolled processes, the log value of the ordinate is propor-
tional to the abscissa (1/T). However, the dependences
are not parallel, as with creep tests the changes of the
parameter ( – d) and shear modulus G in Eq. 1 are inde-
pendent of the iron diffusion rate.
For T  823 K and creep stress 170 MPa, the effects
of an increase of the temperature on the experimental
creep rate and the iron creep diffusion rate are:
lg exp = –7.677 + 2.11 10
4 1/T and
lg DFecreep= –19.277 + 2.07 10
4 1/T (10)
In Eq. 1 the constant k = 4.78 · 103 = 3.5/2 = 4.78 ·
103. The equation can be written for the similar steel as
well as coarsening and dissolution rate of carbide par-
F. VODOPIVEC et al.: EFFECT OF CREEP STRAIN ON CREEP RATE IN THE TEMPERATURE RANGE 550–640 °C
Materiali in tehnologije / Materials and technology 48 (2014) 4, 545–548 547
Table 1: Basic parameters and results of the calculations for the temperature 823–913 K (550–640 °C)
Tabela 1: Osnovni parametri in rezultati izra~unov za temperature 823–913 K (550–640 °C)
Temperature (K) 823 853 883 913
Temperature (°C) 550 580 610 640
Shear modulus (MPa · 103) 54.7 53.7 52.3 51.7
Coars. part size, dc/nm 148 148 148.2 148.7
Average spacing of dc particles (c) 523 523 524 526
Particles coarsening rate (m3 s–1) 2.52 · 10–31 1.28 · 10–30 4.42 · 10–30 1.40 · 10–29
DFetemp /(m
2 s–1) 3.68 · 10–20 1.33 · 10–19 4.41 · 10–19 1.35 · 10–18
DCrtemp/(m
2 s–1) 4.72 · 10–20 1.85 · 10–19 6.64 · 10–19 2.18 · 10–18
DFecreep/(m
2 s–1) 3.74 · 10–20 2.61 · 10–19 1.80 · 10–18 1.14 · 10–17
Kccreep /(m
3 s–1) 2.52 · 10–31 3.29 · 10–31 1.80 · 10–29 5.65 · 10–29
Creep rate by di and i, i/(10
–8 s–1) 1.99 7.08 23.80 70.20
Creep rate of growth of da, at/(10
–8 s–1) 1.99 7.09 24.00 70.50
Size of dissolved particles, dd/nm 6.26 8.36 11.90 17.30
Part of undissolved particles, dd/(10
–2 Ni) 90.16 88.09 85.12 82.22
Number of undissolved particles, NpT/(10
19 m–3) 1.28 1.25 1.21 1.17
Average size of undissolved particles (nm) 154.00 155.00 157.00 159.00
Average spacing of undissolved particles, ud/nm 569.00 574.00 580.00 587.00
Creep rate due to t increase, t/(10
–8 s–1) 2.07 7.44 25.30 74.60
Experimental creep rate, exp/(10
–8 s–1) 2.50 15.50 110.00 720.00
Ratio 19/18 1.21 2.08 4.34 9.65
Ratio 9/7 1.02 2.04 4.27 8.44
Ratio 18/11 1.05 1.05 1.05 1.05
Figure 1: Creep rate calculated for undissolved particles (cud), expe-
rimental creep rate (crep), iron diffusion rate by tempering (DFetemp)
and apparent iron diffusion rate by creep test (DFecreep) versus the
inverse of the creep test temperature
Slika 1: Hitrost lezenja, izra~unana za neraztopljene izlo~ke (cud),
eksperimentalna hitrost lezenja (crep), hitrost difuzije `eleza pri
`arjenju (DFetemp) in navidezna hitrost difuzije `eleza pri lezenju
(DFecreep) v odvisnosti od recipro~ne vrednosti temperature
ticles as  k Dn1 Fecreep with k b k TG1
2= ( / ) b and n =
3.65 and the effect of change of creep stress on iron
diffusion and creep rate deduced.
The ratio of the experimental creep rate rsp. creep
rate calculated considering the coarsening of particles
and the dissolution of particles versus the creep tempe-
rature shown in Figure 2 is independent of the creep
temperature. It indicates that the increase of temperature
has an equal effect on the dissolution and coarsening of
the particles, as both depend on the diffusion rate of iron
and chromium.
The growth of iron creep diffusion rate (DFetemp) with
increasing temperature is lower than the change of the
creep rate and the difference is much higher than 1.15
due to the creep stress deduced from data in.4 In the
particle disjunctive matrix the gliding stress is decreased
strongly, while the climbing stress is diminished much
less. As creep strain consists of the glide and the climb
of dislocations, it is reasonable to assume that the climb
velocity is greater with the greater diffusion rate of
vacancy in matrix. Based on available data, it seems
reasonable to assume that DFecreep in Table 1 and Figure
1 represent5 the increase of a parallel increase of climb
velocity, rsp. creep rate related to the content of vacan-
cies.
4 CONCLUSIONS
• The coarsening and dissolution rates for M23C6
particles in a high-chromium, creep-resistant steel
was calculated for the temperature interval 550 °C to
640 °C;
• The creep rate was then calculated for the investi-
gated range of temperature using the equation with
particle size and spacing as parameters and
determined experimentally with static tests 100 h by
creep stress 170 MPa for equal temperatures;
• In the considered temperature interval, the creep rate
was increased for above one order of magnitude
stronger with small particles dissolution than with
coarser particles growth;
• The iron apparent diffusion rate increased strongly by
increasing creep tests temperature;
• Based on the obtained results, empirical relations are
deduced for the increase of the experimental creep
rate and the apparent iron diffusion rate in ferrite by
creep tests with increased temperature.
• It is suggested that the increase of the experimental
creep rate with creep temperature are related to a
greater iron diffusion rate and a greater climb
velocity due to the greater content of vacancies as
well as the climb stress in the particles disjunctive
matrix.
Acknowledgement
The authors are indebted to the Slovenian Research
Agency (ARRS) for the support of the investigation.
5 REFERENCES
1 F. Vodopivec, B. @u`ek, F. Kafexhiu, Change of M23C6 particles size
and spacing by tempering a high chromium creep resistant steel in
range 800 to 550 °C, Steel Research Int., doi: 10.1002/srin.
201400013
2 A. L. Bowman, G. P. Arnold, E. K. Storman, N. G. Nereson, Acta
Cryst. B, 28 (1972), 3102–3103
3 I. V. Valikova, A. V. Nazarov, Simulation of Characteristics Deter-
mining Pressure Effects on Self Diffusion in BBC and FCC Metals,
The Physics of Metals and Metallography, 109 (2010), 220–226
4 J. W. Jang, J. Kwon, B. J. Lee, Effect of stress on self diffusion in
bcc Fe, An atomistic simulation study, Scripta Mat., 63 (2010),
39–42
5 F. Vodopivec, B. @u`ek, M. Jenko, D. A. Skobir-Balanti~, M. Godec,
Mat. Sci. Tech., 29 (2013), 451–455
6 E. Hornbogen, Einfluss von Teilchen einer zweiten Phase auf das
Zeitverhalten, In: W. Dahl, W. Pitch, Festigkeits- und Bruchverhalten
bei höheren Temperaturen, Verl. Stahleisen, Düsseldorf 1980, 31–52
7 Y. F. Yin, R. G. Faulkner, Physical and elastic behaviour of creep
resistant steels, In: F. Abe, T. U. Kern, R. Viswanathan (eds.), Creep
resistant steels, Woodhead Publ. LTD., Cambridge, England 2008,
217–240
8 F. Vodopivec, D. Steiner - Petrovi~, B. @u`ek, M. Jenko, Coarsening
rate of M23C6 particles in a high chromium creep resistant steel, Steel
Res. Int., 84 (2013), 1110–1114
9 H. Oikawa, Y. Iijima, Diffusion behaviour of creep-resistant steels,
In: F. Abe, T. U. Kern, R. Wiswanathan (eds.), Creep resistant steels,
Woohead Publ. Lim., Cambridge, England 2008, 241–264
10 H. K. D. H. Bhadeshia, Mat. Sci. Techn., 24 (2008), 128–130
11 D. A. Skobir, F. Vodopivec, M. Jenko, S. Spai}, B. Markoli, Z.
Metallkd., 95 (2004) 11, 1020–1024
12 F. Vodopivec, D. A. Skobir - Balanti~, M. Jenko, B. @u`ek, M.
Godec, Mater. Tehnol., 46 (2012) 6, 633–636
13 F. Vodopivec, J. Vojvodi~-Tuma, B. [u{tar{i~, R. Celin, M. Jenko,
Mat. Sci. Tech., 27 (2011), 939–942
14 B. @u`ek, Influence of the distribution of carbide particles on the
activation energy for the creep deformation of steels, Ph. D. Thesis,
J. Stefan International Postgraduate school, Ljubljana, 2013
F. VODOPIVEC et al.: EFFECT OF CREEP STRAIN ON CREEP RATE IN THE TEMPERATURE RANGE 550–640 °C
548 Materiali in tehnologije / Materials and technology 48 (2014) 4, 545–548
Figure 2: Effect of creep test temperature on the ratio of the experi-
mental creep rate (exp) versus creep rates calculated considering the
particles coarsening (pc ) and dissolution (pd) by tempering at creep
test temperature and ratio of creep rates calculated for dissolution and
growth of average particles size (  pd pc/ )
Slika 2: Vpliv temperature preizkusa lezenja na razmerje eksperimen-
talna hitrost lezenja (exp) proti hitrosti lezenja, izra~unani za rast (pc )
in raztapljanje (pd) izlo~kov pri `arjenju pri temperaturi preizkusa, in
razmerje izra~unanih hitrosti lezenja za raztapljanje izlo~kov in rast
povpre~ne velikosti izlo~kov(  pd pc/ )
H. SRBOVÁ et al.: IDENTIFICATION OF THE INITIAL FAILURE AND DAMAGE OF SUBSTITUENTS ...
IDENTIFICATION OF THE INITIAL FAILURE AND
DAMAGE OF SUBSTITUENTS OF A UNIDIRECTIONAL
FIBER-REINFORCED COMPOSITE USING A
MICROMODEL
UPORABA MIKROMODELA ZA UGOTAVLJANJE ZA^ETNIH
NAPAK IN PO[KODB SESTAVIN KOMPOZITA, ENOSMERNO
OJA^ANEGA Z VLAKNI
Hana Srbová, Tomá{ Kroupa, Robert Zem~ík
University of West Bohemia in Pilsen, Department of Mechanics, Univerzitní 22, 306 14 Plzeò, Czech Republic
hsrbova@kme.zcu.cz
Prejem rokopisa – received: 2012-08-31; sprejem za objavo – accepted for publication: 2013-10-11
The main goal of this investigation was to test the capabilities of simple failure and damage criteria for the material phases of a
unidirectional long-fiber carbon-epoxy composite material. This goal requires the identification of the material parameters in
order to simulate simple tensile tests performed on thin coupons with various fiber orientations. Furthermore, a failure criterion
for the fibers and a damage criterion for the matrix are proposed and the material parameters are identified in order to capture
the moment of the first failure or damage. Finite-element analyses are performed on a unit cell with applied periodical boundary
conditions. The identifications were performed using Optislang 3.2.0 and Python 2.4. The finite-element analyses were
performed using Abaqus 6.11-1.
Keywords: unidirectional fiber composite, identification, matrix work-hardening function, unit cell, micromodel, non-linear
Glavni namen predstavljenega dela je preizkus sposobnosti enostavnih meril za napake in po{kodbe pri sestavinah
epoksikompozita, enosmerno oja~anega z dolgimi ogljikovimi vlakni. Ta cilj zahteva poznanje parametrov materiala za
simulacijo enostavnega nateznega preizkusa, izvedenega na tankih plo{~icah z razli~no orientacijo vlaken. Predlagana so merila
za poru{itev vlaken, za po{kodbe osnove in ugotovljeni so parametri materiala za ugotovitev trenutka prve poru{itve ali
po{kodbe. Analiza kon~nih elementov na osnovni celici je bila opravljena z uporabo periodi~nih robnih pogojev. Identifikacija
je bila izvr{ena z uporabo Optislang 3.2.0 in Python 2.4. Analize kon~nih elementov so bile izvr{ene z Abaqus 6.11-1.
Klju~ne besede: kompozit, enosmerno oja~an z vlakni, identifikacija, funkcija deformacijskega utrjevanja osnove, osnovna
celica, mikromodel, nelinearen
1 INTRODUCTION
Composite materials are widely used due to their
advantageous stiffness-to-weight and strength-to-weight
ratios. The composites’ material properties depend on the
volume ratio of the constituents and their properties.
There are several approaches to modeling a compo-
site material’s behavior. The material model can be
modeled on the micro-, meso- and macro-scale levels.
The macromechanical approach requires, compared to
the micromechanical, less detailed and time-consuming
modeling. In1 the non-linear behavior of composites was
modeled using a macroscopic approach. Nevertheless,
the micromechanical approach gives more detailed infor-
mation about the constituents,2 for example, the damage
and failure of the matrix and fiber. It is also possible to
capture the nonlinear behavior of a composite by
assuming the constituents to be nonlinear materials. A
composite material may vary its properties due to matrix
cracking and fiber breaking during loading.
A knowledge of damage mechanisms on the micro
level of the material is important during the design pro-
cesses of structures. This fact is more important when it
comes to a structure made of heterogeneous materials,
such as composites or when the task is to avoid the
damage or failure of the structure.
2 EXPERIMENT
Tensile tests of thin coupons made of unidirectional
longfiber carbon-epoxy composite SE84LV-HSC-450-
400-35 were performed on a ZWICK/ROELL Z050 test
machine. The coupons were cut from one large plate
using a water jet.3
Materiali in tehnologije / Materials and technology 48 (2014) 4, 549–553 549
UDK 66.017:519.61/.64 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(4)549(2014)
Figure 1: Geometry of composite coupons (mm) 2
Slika 1: Geometrija plo{~ice kompozita (mm) 2
The fiber direction forms angles of 0°, 15°, 30°, 45°,
60°, 75° and 90° with the direction of the loading force
(Figure 1). There were 10 specimens tested for each
angle 
. The cracked specimens are shown in Figure 2.
The specimens loaded along the fiber direction were
fractured by fiber failure. All the specimens loaded at a
different angle are fractured by matrix failure. The
resulting force-displacement diagrams are shown in
Figure 3.
2.1 Micromodel
The geometry of the finite-element model (micro-
model) with a periodically repeated volume (unit cell) of
the unidirectional composite material was created in the
CAD program Siemens NX. The unit cell is meshed
using ten-node tetrahedral elements. A perfect honey-
comb distribution of fibers and a fiber volume ratio of
55 % were assumed (Table 1). Furthermore, the finite
strain theory was used. The geometry of the unit cell was
exported to Abaqus/CAE 6.11-1 for the finite-element
analysis.
The global coordinate system (xyz) is given by the
force direction (x) and the direction perpendicular to the
composite surface (z). The local coordinate system (123)
is defined by the unit-cell edges, where the axis direc-
tions correspond to the fiber direction (1) and the direc-
tions perpendicular to it (Figure 4).
Assuming uniaxial stress across the whole specimen,
the behavior of the material can be simulated by loading
the unit cell with a normal stress corresponding to an
external force F:
 x
F
A
= (1)
where A is the cross-section of the specimen.
The effect of the loading force is transformed to the
local coordinate system using the transformation:





 
 
 


 




⎡
⎣
⎢
⎢
⎢
⎤
⎦
⎥
⎥
⎥
= −
cos sin sin cos
sin cos s
2 2
2 2
2
2 in cos
sin cos sin cos cos sin

 


 
 
 
 
 



− −
⎡
⎣
⎢
⎢
⎢
⎤
⎦
⎥
⎥
⎥
⎡
2 2
x
⎣
⎢
⎢
⎢
⎤
⎦
⎥
⎥
⎥
(2)
where 
 is the angle of rotation between the local and
global coordinate systems.3
The results from the finite-element analysis (strains)
are transformed back to the global coordinate system
using the transformation:




 
 
 


 

x
y
xy
⎡
⎣
⎢
⎢
⎢
⎤
⎦
⎥
⎥
⎥
=
−cos sin sin cos
sin cos sin
2 2
2 2 
 


 
 
 
 
 







cos
sin cos sin cos cos sin2 2 2 2−
⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥

⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥
(3)
H. SRBOVÁ et al.: IDENTIFICATION OF THE INITIAL FAILURE AND DAMAGE OF SUBSTITUENTS ...
550 Materiali in tehnologije / Materials and technology 48 (2014) 4, 549–553
Figure 4: Rotated coordinate systems and loading of the unit cell
Slika 4: Rotirani koordinatni sistemi in obremenitve osnovne celice
Table 1: Geometry ratios of the unit cell
Tabela 1: Geometrijska razmerja osnovne celice
Fiber radius r
Short side length 1.28 r
Long side length 2.22 r
Figure 3: Measured force-displacement diagrams (grey) and averaged
values (black)2
Slika 3: Diagrami sila – raztezek (sivo) in povpre~ne vrednosti (~rno)2
Figure 2: Cracked specimens with aluminum tabs on a specimen with
a 0° orientation2
Slika 2: Razpokani vzorci z aluminijevimi zavihki na vzorcu z
orientacijo 0° 2
The unit cell must also respect the periodical boun-
dary conditions:
Δ
Δ = −
Δ = −
u u u
v v v
w w w
= −B A
B A
B A
(4)
where u, v and w are the translation differences of
a pair of opposing nodes in directions 1, 2 and 3,
respectively. These differences must remain constant for
all the pairs of corresponding nodes on opposite sides.
The periodical boundary conditions were implemented
in Abaqus/CAE.
2.2 Material model of fiber
The fibers were modeled as transversely isotropic
elastic material respecting the non-linear behavior in the
axis direction. The dependence of the longitudinal
Young’s modulus of the fibers on the strain in the axial
direction of the fibers is4:
E E g11 11 1
0
111( ) ( ) = + (5)
where g is the coefficient describing the measure of
non-linearity and E1
0 is the initial Young’s modulus of
the fiber in the longitudinal direction.
It is assumed that fiber failure occurs when the failure
index in the axial direction, defined as:
f
X1
1=
 f
T
f (6)
where 1
f is stress in direction 1 in the fibers and X T
f is
the tensile strength of the fibers, reaches the value of 1.
2.3 Material model of matrix
The matrix was modeled as an isotropic elasto-plastic
material. The elastic parameters were the Young’s modu-
lus Em and the Poisson’s ratio μm. The Von Mises plasti-
city was used with the isotropic hardening and work-
hardening function proposed in the form:
 






y y
y
n n
T
T
= +
⎛
⎝
⎜
⎞
⎠
⎟
⎡
⎣
⎢
⎤
⎦
⎥
0
0
1
p
p
(7)
where the yield stress  y depends on the equivalent
plastic strain p and on additional material parameters
where  
y
is the initial yield stress, T0 is the tangent of
the hardening curve for  p = 0,  
y
is asymptotic stress
and n is a shape parameter.
The damage initiation criterion inspired by the crite-
rion for the ductile damage of metals was used, where
the fracture is assumed to be caused by the coalescence
of voids in a matrix and the damage initiates when the
following condition (damage factor) is fulfilled:


 
D
p
D
p
d
= =∫
( )
1 (8)
where function:
  D
p
D
p= ( ) (9)
is the equivalent plastic strain at the onset of the damage
and it is a function of the stress triaxiality factor:



= h
eqv
(10)
H. SRBOVÁ et al.: IDENTIFICATION OF THE INITIAL FAILURE AND DAMAGE OF SUBSTITUENTS ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 549–553 551
Figure 6: Force-displacement diagrams for 15° to 90° specimens
(black – identified, grey – experiments)
Slika 6: Diagrami sila – raztezek za vzorce od 15° do 90° (~rno –
identificirano, sivo – eksperimenti)
Figure 5: Force-displacement diagram for 0° specimen (black –
identified, grey – experiments)
Slika 5: Diagram sila – raztezek za vzorec 0° (~rno – identificirano,
sivo – eksperimenti)
where h is hydrostatic stress and eqv is the Von Mises
stress.5
2.4 Identification process
An optimization process, similar to the one used in2,
was performed using Python scripts, the optimization
software OptiSlang and the finite-element system Aba-
qus. The identification consisted of two steps. The first
step was to find the best combination of all the elastic
and plastic material parameters by minimizing the diffe-
rence between the experimental and numerical force-
displacement diagrams. Moreover, the parameter 

which represents the inaccuracy when cutting the
samples from the plate was identified. The second step
was to identify the tensile strength of fibers and the
function D
p
by minimizing the difference between the
value of D and 1 at the moment of fracture in the
experiment. The dependency of the equivalent plastic
strain on the onset of damage D
p
was proposed as a
piecewise linear function with five control vertices.
3 DISCUSSION OF RESULTS
The identified parameters and the parameters given
by the manufacturer are summarized in Table 2. The
identified parameters correspond well with the ones
identified in2 where the analysis was performed in
MSC.Marc software using eight-node brick elements.
The force-displacement diagram for the specimen 0°
shows a convex shape caused by stiffening of the fibers
(Figure 5). The force-displacement diagrams of the rest
of the specimens have concave shapes caused by the
matrix non-linearity (Figure 6). The identified shape of
the work-hardening function is shown in Figure 7. The
ability of the model to predict the initial damage of the
matrix is visualized in Figure 8. The identified depen-
dency of D
p
on the triaxiality stress factor is shown in
Figure 9. The distribution of the damage factor for the
matrix in the 45° specimen in the moment when the first
damage to the matrix occurs is shown in Figure 10.
H. SRBOVÁ et al.: IDENTIFICATION OF THE INITIAL FAILURE AND DAMAGE OF SUBSTITUENTS ...
552 Materiali in tehnologije / Materials and technology 48 (2014) 4, 549–553
Figure 10: Distribution of damage factor (factor is not calculated for
fibers, deformation is magnified 10 times)
Slika 10: Razporeditev faktorja po{kodbe (faktor ni izra~unan za
vlakna, deformacija je pove~ana za 10-krat)
Figure 9: Dependency of the equivalent plastic strain at the onset of
damage on the triaxiality stress factor
Slika 9: Odvisnost ekvivalenta plasti~nega raztezka na za~etku
po{kodbe od triosnega faktorja napetosti
Figure 8: Maximum damage factors for specimens from 15° to 90°
(black – identified, grey – experiments)
Slika 8: Faktorji maksimalne po{kodbe pri vzorcih od 15° do 90°
(~rno – identificirano, sivo – eksperimenti)
Figure 7: Work-hardening curve
Slika 7: Krivulja deformacijskega utrjevanja
Table 2: Summarized parameters (i = identified, m = manufacturer)
Tabela 2: Povzeti parametri (i = ugotovljeni, m = proizvajalec)
Parameter Units Value note
E1
0 GPa 188.73 i
g – 18.76 i
E2 = E3 GPa 15.00 m
μ12 = μ13 – 0.30 m
μ23 – 0.40 m
G12 = G13 GPa 5.00 m
G23 GPa 4.00 m
XTf GPa 3.88 i
Em GPa 3.61 i
μm – 0.40 m
0
y MPa 29.24 i
T0 GPa 8.49 i
A
y MPa 93.54 i
n – 1.35 i

 ° –0.98 i
4 CONCLUSION
Finite-element analyses of simple tensile tests on
unidirectional composite specimens with various fiber
orientations were performed. The material parameters of
the orthotropic non-linear elastic material of the fibers
and of the non-linear elasto-plastic material of the matrix
were identified. The simple failure (fiber) and the
damage (matrix) criteria were successfully tested and
good agreement with the experiments was achieved.
Future work will focus on an approximation of the
dependency of the equivalent plastic strain at the onset of
the damage on the triaxiality stress factor using a smooth
function and modeling of the subsequent damage process
using a cohesive-theory-based approach.
Acknowledgement
The work has been supported by the projects GA
P101/11/0288 and European project NTIS – New Tech-
nologies for Information Society No. CZ.1.05/1.1.00/
02.0090.
5 REFERENCES
1 T. Kroupa, V. La{, R. Zem~ík, Improved Non-Linear Stress–Strain
Relation for Carbon–Epoxy Composites and Identification of Mate-
rial Parameters, Journal of Composite Materials, 45 (2011) 9, 1045–
1057
2 H. Srbová, T. Kroupa, R. Zem~ík, Identification of the Material Para-
meters of a Unidirectional Fiber Composite Using a Micromodel,
Mater. Tehnol., 46 (2012) 5, 431–434
3 D. Roylance, Transformation of Stresses and Strains, Department of
Materials Science and Engineering, Massachusetts Institute of Tech-
nology, Cambridge, 2001
4 I. M. Djordjevi}, D. R. Sekuli}, M. M. Stevanovi}, Non-Linear Ela-
stic Behavior of Carbon Fibers of Different Structural and
Mechanical Characteristic, Journal of the Serbian Chemical Society,
72 (2007) 5, 513–521
5 S. S. Bhadauria, M. S. Hora, K. K. Pathak, Effect of Stress Triaxi-
ality on Yielding of Anisotropic Materials under Plane Stress Con-
dition, Journal of Solid Mechanics, 1 (2009) 3, 226–232
H. SRBOVÁ et al.: IDENTIFICATION OF THE INITIAL FAILURE AND DAMAGE OF SUBSTITUENTS ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 549–553 553

B. MA[EK et al.: TRANSFORMATION-INDUCED PLASTICITY IN STEEL FOR HOT STAMPING
TRANSFORMATION-INDUCED PLASTICITY IN STEEL
FOR HOT STAMPING
VPLIV PRETVORBE NA PLASTI^NOST JEKLA ZA VRO^E
[TANCANJE
Bohuslav Ma{ek1, Ctibor [tádler1, Hana Jirková1, Peter Feuser2, Mike Selig3
1Research Centre of Forming Technology – FORTECH, University of West Bohemia in Pilsen, Univerzitní 22, 306 14 Pilsen, Czech Republic
2Daimler AG, Bela-Barenyi-Str. 1, 71063 Sindelfingen, Germany
3AutoForm Development GmbH, Technoparkstrasse 1, 8005 Zürich, Switzerland
ctibor@email.cz
Prejem rokopisa – received: 2012-10-01; sprejem za objavo – accepted for publication: 2013-10-07
The hot-stamping process for manufacturing car-body components was patented in 1977. Its main advantages include the
precision of the product shape, the reduced spring-back effect and the resulting high strength of steel parts upon hardening.
Boron- and manganese-alloyed steels have been specially developed for this process. The 22MnB5 grade is a typical represen-
tative of this group with its strength up to 1500 MPa.
For the desired mechanical properties to be achieved, the final microstructure should consist primarily of martensite without any
substantial amounts of other phases. Further development of press hardening, therefore, requires all the phenomena associated
with the phase transformations during the cooling between dies to be well mapped. Significant parameters, in this respect,
include phase-transformation temperatures and the data on a number of additional phenomena, including transformation-
induced plasticity. Transformation-induced plasticity is manifested when a part held between dies undergoes a phase
transformation, typically, when being under stress. In the course of a lattice rearrangement, this stress causes the atoms to
occupy more favourable positions in terms of energy. At the macro-level, this can be detected as a change in the dimensions,
which, at the same time, significantly relieves the stress, therefore, eliminating the spring back. Despite a profound importance
of this phenomenon for the press-hardening process, it has not been explored in detail up to now.
For these reasons, a 22MnB5 steel grade was employed in this project. The impacts of the tensile and compressive stresses
occurring during phase transformations, upon the changes in the expansion of the sheet specimens, were explored using the
schedules that simulated real-world hot-stamping and press-hardening processes.
Keywords: press hardening, hot stamping, spring-back effect, 22MnB5, transformation-induced plasticity
Postopek vro~ega {tancanja za izdelavo komponent karoserije avtomobilov je bil patentiran leta 1977. Njegove glavne prednosti
so natan~nost oblike izdelka, zmanj{an vzmetni u~inek in visoka trdnost jeklenih delov po toplotni obdelavi. Posebno za ta
postopek so bila razvita jekla, legirana z borom in manganom. Zna~ilen predstavnik te skupine je jeklo 22MnB5 s trdnostjo do
1500 MPa.
Za doseganje `elenih mehanskih lastnosti mora biti kon~na mikrostruktura martenzitna, brez ve~jih dele`ev drugih faz.
Nadaljnji razvoj utrjevanja pri stiskanju zahteva, da se obvladajo vsi pojavi, ki so povezani s faznimi premenami med
ohlajanjem v orodju. S tega vidika so pomembni parametri, ki vklju~ujejo temperature faznih premen in podatki o {tevilnih
drugih pojavih, vklju~no z vplivom fazne premene na plasti~nost. Vpliv fazne premene na plasti~nost se poka`e, ko se v delu
med orodjema zgodi fazna premena, navadno pod tlakom. Med preureditvijo mre`e napetosti povzro~ijo, da atomi zasedejo
energijsko bolj ugodne polo`aje. Na makronivoju se to ka`e kot sprememba dimenzij, ki isto~asno zmanj{ajo napetosti in
odpravijo u~inek vzmeti. Kljub pomembnosti ta pojav utrjevanja pri stiskanju do sedaj {e ni bil podrobno raziskan.
Zato je bilo v tej raziskavi uporabljeno jeklo 22MnB5. Raziskani so bili u~inki nateznih in tla~nih napetosti, ki se pojavijo med
fazno premeno pri spremembi {irjenja vzorcev plo~evine, s simulacijo realnega vro~ega {tancanja in procesa utrjevanja s
stiskanjem.
Klju~ne besede: utrjevanje s stiskanjem, vro~e {tancanje, vzmetni u~inek, 22MnB5, vpliv pretvorbe na plasti~nost
1 INTRODUCTION
Transformation-induced plasticity is an all-too-often
overlooked phenomenon in manufacturing today. One of
the reasons for this is the fact that its impact on the out-
come of the entire manufacturing process is not readily
visible. Yet, there are technological problems in the
industry that may be caused by transformation-induced
plasticity and other issues that could be resolved by
making use of the very same transformation-induced pla-
sticity phenomenon.1,2 An example of the important role
of transformation-induced plasticity is the production of
sizable forgings, namely, their heat treatment.3 Only
recently have these phenomena received more attention
and the findings are reflected in FEM simulations. Their
descriptions have become more detailed, having an effect
on the current material models. Transformation-induced
plasticity has an important role in the process of hot
sheet forming with the subsequent press hardening,
known as hot stamping. In the hot-stamping process,
transformation-induced plasticity combined with other
effects reduces the spring back. Consequently, the final
pressed parts achieve greater shape precision than the
high-strength deep-cold-drawn sheet products.4 Since
this phenomenon typically occurs in the production of
complex-shape deep-drawn sheet products, it is difficult
to assess it accurately. Without an exact description of its
effect, no production-relevant and sufficiently accurate
material-based model can be constructed.5 This was the
background of the present experiment.
Materiali in tehnologije / Materials and technology 48 (2014) 4, 555–557 555
UDK 621.98 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(4)555(2014)
2 EXPERIMENTAL WORK
Modified tension-test pieces with a 10 mm width, 50
mm length and a thickness of 1.5 mm were used in this
study. The schedules were simulations of real-world
hot-stamping and press-hardening processes. The mate-
rial was a standard 22MnB5 steel grade (Table 1).
Table 1: Chemical composition of the 22MnB5 grade in mass frac-
tions (w/%)
Tabela 1: Kemijska sestava jekla 22MnB5 v masnih dele`ih (w/%)
C Cr Mn Si B P S Al
0.22 0.2 1.25 0.2 0.0029 0.025 0.015 0.006
The main objective of the research was to explore the
intensity of the transformation-induced plasticity pheno-
menon during the cooling process. The hot-stamping
process was simulated by heating the specimens to 950
°C. The specimens were kept free from the longitudinal
stress during the heating. They were held at the above
temperature for 3 min and then cooled down to 770 °C in
10 s. This step was a simulation of cooling the trimmed
feedstock during handling. Between 770 °C and 600 °C,
the stress within the specimen increased linearly from 0
to the prescribed level, representing the stress formation
in a real-world process. This stress level was then main-
tained during the subsequent constant-rate cooling. The
cooling rates were specified on the basis of the CCT
diagrams constructed under the conditions identical to
real-world processes. The cooling rates of 25 °C/s and
40 °C/s were selected as the characteristic values result-
ing in various microstructural states. The first cooling
rate led to a formation of a mixed microstructure of
bainite, martensite and some ferrite. The higher cooling
rate led to a transformation into martensite. Conse-
quently, the differences between the final dimensions of
the final disparate microstructures could be compared.
The changes in the length of the test specimen were
monitored in the course of the experiment, while the
longitudinal stress in the part was controlled.
The amount of stress was defined on the basis of the
measurements of the stress-strain characteristics at the
constant temperature of 600 °C, carried out at the
beginning of this research. It was less than one quarter of
the stress that caused plastic deformation. In the sub-
sequent tests, this amount was gradually increased to
twice and three times its value. Hence, the stress was
kept in the elastic region at all times to prevent any
substantial contribution to the elongation of the test
specimen, caused by the ordinary plastic deformation
and creep during the testing. These stress levels were as
follows: (31, 62 and 93) MPa. The ratio of the extension
caused by the elastic deformation under the highest load
to the sum of the overall contraction and the deformation
due to transformation-induced plasticity was no more
than a mere fraction of one percent (Figure 1).
3 RESULTS AND DISCUSSION
In the region of transformation, the experiment
revealed a non-linear extension of the material along the
applied stress direction (Figure 1). This is a definite evi-
dence of the profound effects of transformation-induced
plasticity. Transformation-induced plasticity takes effect
at the moment of lattice reconfiguration, when the atoms
take their new positions. The new positions are con-
sistent with the transition into a lower-energy, more
favourable state. This is the main cause of the permanent
deformation, manifested as the non-linearity on the
specimen elongation curve.
The experiment showed that the length of the test
specimen decreased by about 0.5 mm upon the cooling
from 600 °C to 70 °C without any longitudinal stress.
This corresponds to the 0 % transformation plasticity.
When the longitudinal stress of 31 MPa was applied, the
contraction was approximately 0.35 mm. The contraction
of the test specimen was, in this case, by 0.15 mm lower
than in the previous case, due to the transformation pla-
sticity. Once the longitudinal stress doubled, the contrac-
B. MA[EK et al.: TRANSFORMATION-INDUCED PLASTICITY IN STEEL FOR HOT STAMPING
556 Materiali in tehnologije / Materials and technology 48 (2014) 4, 555–557
Figure 1: Effect of stress on the length of the test specimen
Slika 1: Vpliv napetosti na dol`ino preizku{anca
tion was 0.13 mm. At the triple stress value, the contrac-
tion caused by the cooling was outweighed by the
transformation-induced plasticity effect, causing the
specimen to extend by an amount between 0.22 mm and
0.27 mm.
In this experiment, the effect of diffusion should be
taken into account as well, as it contributes to the process
even though its extent varies. However, its impact cannot
be explored separately from the other effects. The expe-
riment showed that at the lower cooling rates – when the
times available for diffusion were longer – the elongation
increased, though slightly. The impact of creep increases
with the stress level.
At the stresses above 70 MPa, the effect of transfor-
mation-induced plasticity can compensate for the con-
traction, thus eliminating residual stresses (Figure 1).
The relative extension between 600 °C and 70 °C
rose from 0 %, which was detected during the cooling
without any stress applied, to approximately 0.55 % at
31 MPa. It then increased almost linearly with the
increasing longitudinal stress to twice as high a value. At
62 MPa, it reached 1 % (Figure 2). When the stress was
increased further by one third of its value, the relative
extension was doubled, reaching more than 2 %.
4 CONCLUSION
The experiment conducted on the sheet specimens
made from the 22MnB5 steel involving the cooling rates
of 25 °C/s and 40 °C/s and various applied stress levels
revealed the intensity of the transformation-induced
plasticity effects during hot-stamping and press-harden-
ing processes. At the temperatures between 600 °C and
70 °C, the test specimens contracted. The contraction
was compensated for by an expansion due to the longi-
tudinal tensile stress and transformation-induced plasti-
city. With the longitudinal stress increasing from 0 to 93
MPa, the relative extension of the test specimen becomes
non-linear, increasing from 0 % up to 2 %. The results
show that at the stresses above 70 MPa, transformation-
induced plasticity and additional related phenomena can
compensate for the contraction due to cooling, thus
reducing the spring back.
The present experiment was not intended to exactly
match the real-world contraction behaviour of the press-
hardening process. Its purpose was to provide an insight
into and the data on the transformation-induced plasticity
under the specific cooling conditions of the hot-stamping
process.
Acknowledgement
This paper includes the results created within project
SGS-2012-040, New Methodology for Description of
Transformation Processes in Real Materials. The project
is subsidised from the specific resources of the state
budget for research and development.
5 REFERENCES
1 B. Norrbottens Jaernverk, Patent GB 1490535, Manufacturing of a
hardened steel article, 11 February 1977
2 H. Karbasian, A. E. Tekkaya, A review on hot stamping, Journal of
Materials Processing Technology, 210 (2010), 2103–2118
3 M. Taschauer, V. Wieser, S. Schramhauser, T. Hatzenbichler, B.
Buchmayer, Die Wärmebehandlung der Stähle unter Einbeziehung
des Selbstanlassens, Proc. of the XXXI. Verformungskundliches
Kolloquium, Planneralm, 2012, 91–96
4 A. Turetta, S. Bruschi, A. Ghiotti, Investigation of 22MnB5 forma-
bility in hot stamping operations, Journal of Materials Processing
Technology, 177 (2006), 396–400
5 S. Ertürk, M. Sester, M. Selig, P. Feuser, K. Roll, A Thermo-Mecha-
nical-Metallurgical FE Approach for Simulation of Tailored Tem-
pering, Proc. of the 3rd International Conference – Hot sheet metal
forming of high-performance steel, Kassel, Germany, 2011
B. MA[EK et al.: TRANSFORMATION-INDUCED PLASTICITY IN STEEL FOR HOT STAMPING
Materiali in tehnologije / Materials and technology 48 (2014) 4, 555–557 557
Figure 2: Relative elongation versus longitudinal stress
Slika 2: Relativni raztezek v odvisnosti od vzdol`ne napetosti

M. KRGOVI] et al.: INFLUENCE OF TEMPERATURE AND BINDER CONTENT ON THE PROPERTIES ...
INFLUENCE OF TEMPERATURE AND BINDER
CONTENT ON THE PROPERTIES OF A SINTERED
PRODUCT BASED ON RED MUD
VPLIV TEMPERATURE IN KOLI^INE VEZIVA NA LASTNOSTI
PROIZVODA IZ SINTRANEGA RDE^EGA BLATA
Milun Krgovi}1, Ivana Bo{kovi}1, Radomir Zejak2, Milo{ Kne`evi}2
1University of Montenegro, Faculty of Metallurgy and Technology, D`ord`a Va{ingtona bb, 81000 Podgorica, Montenegro
2University of Montenegro, Faculty of Civil Engineering, D`ord`a Va{ingtona bb, 81000 Podgorica, Montenegro
milun@ac.me
Prejem rokopisa – received: 2013-01-29; sprejem za objavo – accepted for publication: 2013-10-30
In this paper we present an investigation into the properties of a sintered product based on red mud with respect to the sintering
temperature and the binder content (illite-kaolinite clay). In the Bayer process for obtaining alumina, red mud is a by-product
(solid waste from the leaching of bauxite). Using red mud as a raw material mixture component to obtain sintered products, a
cheap raw material for construction and ceramics would be obtained, and the environmental problems resulting from the
frequent deposits of red mud in pools would be solved. Based on the overall results of our examinations in this study, it can be
concluded that the change of basic parameters, sintering temperature and binder content (illite-kaolinite clay), can provide
satisfactory characteristics for the sintered product in terms of mechanical properties, volume shrinkage during sintering and
total porosity.
Keywords: red mud, clay, porosity, compressive strength
^lanek obravnava preiskavo lastnosti sintranih proizvodov na osnovi rde~ega blata v odvisnosti od temperature sintranja in
vsebnosti veziva (glina ilit-kaolinit). Rde~e blato (trden odpadek pri izpiranju boksita) je stranski proizvod pri Bayerjevem
procesu pridobivanja glinice. Rde~e blato kot surovina v me{anici za sintranje komponent je poceni material za gradnjo in
keramiko, re{uje pa tudi okoljske probleme, ki izvirajo iz rednega odlaganja rde~ega blata v bazene. Na podlagi rezultatov te
{tudije se lahko sklene, da spreminjanje osnovnih parametrov, temperature sintranja in vsebnosti veziva (glina ilit-kaolinit),
omogo~a zadovoljive lastnosti sintranega proizvoda glede mehanskih lastnosti, kr~enja prostornine med sintranjem in celotne
poroznosti.
Klju~ne besede: rde~e blato, glina, poroznost, tla~na trdnost
1 INTRODUCTION
Red mud contains the following basic components:
iron in the form of oxides, silicon in the form of so-
dium-aluminosilicate and titanium as sodium-titanate.
Caustic soda, NaOH, is also present in red mud, while
with the addition of lime to the process it also contains
CaO in either the free or bound forms.
Illite-kaolinite clays, apart from basic clay minerals,
illite and kaolinite, mostly contain -quartz, Fe2O3,
CaCO3, and are considered as refractory clays (stable up
to a temperature of 1350 °C).1 During the sintering of
this two-component mixture (red mud, illite-kaolinite
clay), depending on the sintering temperature, different
reactions in the solid state occur, as well as the
polymorphic transformation of quartz and the formation
of a liquid phase.2 The relationship between the parti-
cular elements of the microstructure during sintering, in
addition to the firing regime, is importantly influenced
by the mineral content of the raw materials.3,4 Liquid-
phase formation accelerates the reactions in the solid
state due to the diffusion-coefficient increase in such
systems by up to one thousand times.5 The formation of
new crystalline phases in the solid-state reactions during
the sintering process, apart from the above-mentioned
factors, is also influenced by the temperature of the
sintering.6,7
Investigations have shown that the relations between
the elements of the microstructure and the total porosity
of the sintered product are importantly influenced by the
mass-ratio changes between the red mud and the clay in
a two-component mixture.8,9
2 EXPERIMENTAL
The raw material mixture for the production of
samples was formed on the basis of red mud (Aluminium
Plant Podgorica) and illite-kaolinite clay (“Bijelo Polje”)
as binders. The content of the red mud in the raw
material mixture in mass fractions was (10, 20, 30, 40
and 50) %. The samples were formed by plastic shaping
in a mould corresponding to a parallelepiped with di-
mensions of 7.7 cm × 3.9 cm × 1.6 cm. The raw material
mixture components were characterized by a determi-
nation of mineral content and the chemical content,
while the particle size distribution was determined by
granulometric analyses. The density and humidity of the
raw material mixture components were also determined.
Materiali in tehnologije / Materials and technology 48 (2014) 4, 559–562 559
UDK 621.762.5:691:620.1 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(4)559(2014)
The sintering of the samples with different red mud
contents was performed at temperatures of 1100 °C and
1200 °C.5,8 The characterization of the sintered products
included the following: a determination of the linear and
volume shrinkage during sintering, and a determination
of the total porosity and compressive strength. The deter-
mination of the mineral content of the sintered products
was performed by X-ray analysis.
3 RESULTS AND DISCUSSION
The mineral composition of red mud (Figure 1)
determined by the X-ray analysis shows the presence of
the following minerals: hematite, goethite, boehmite,
diaspora, calcite, as well as the presence of soda. The re-
sults of the chemical analysis of red mud (Table 1) show
a high mass content of Fe2O3 (40.78 %) and because of
its negative influence on the quality of sintered product,1
the content of red mud in the raw material mixture does
not exceed 50 %. The clay used as a binder is an illite-
kaolinite type (Figure 2, Table 2), with the presence of
-quartz.
Table 1: The chemical content of red mud in mass fractions (w/%)
Tabela 1: Kemijska sestava rde~ega blata v masnih dele`ih (w/%)
Oxides Fe2O3 Al2O3 SiO2 TiO2 Na2O CaO lg. loss
40.78 17.91 11.28 10.20 6.90 1.50 9.36
The results of the granulometric analysis show that
the red mud has a larger fraction (average grain size 30 μm)
compared to the used clay (average grain size 17 μm).
The density of the clay determined by the pycnometer
method is 2.34 g cm–3, while the density of the red mud
is 2.76 g cm–3.
The results of the volume shrinkage during sintering
(Figure 3) show higher values of shrinkage at the
sintering temperature of 1200 °C compared to 1100 °C,
and the shrinkage decreases with an increase in the
content of red mud in the raw material mixture. This is
explained by the granulometric analysis of the raw
material mixture (red mud has a larger fraction), the
relations between the elements of the microstructure, as
well as by an increase of the total porosity of the sintered
product with an increase of the red mud content.
The total porosity (Figure 4) has lower values at the
sintering temperature of 1200 °C, compared to the
sintering temperature of 1100 °C, and increases with an
increase of the red mud content in the raw material
mixture. During the sintering process of the ceramic
particles all the well-known mass-transport mechanisms
can occur, depending on the sintering conditions.10 Based
on the already-existing data and results it can be
concluded that diffusion is the dominant mass-transport
mechanism under these conditions. Along with the
temperature rise, the reactions in the solid state are
generated by increased diffusion.
M. KRGOVI] et al.: INFLUENCE OF TEMPERATURE AND BINDER CONTENT ON THE PROPERTIES ...
560 Materiali in tehnologije / Materials and technology 48 (2014) 4, 559–562
Table 2: The chemical content of "Bijelo Polje" clay in mass fractions (w/%)
Tabela 2: Kemijska sestava gline "Bijelo Polje" v masnih dele`ih (w/%)
Oxides SiO2 Fe2O3 Al2O3 CaO MgO Na2O K2O SO3 lg. loss
72.68 5.7 10.98 0.48 0.7 0.31 1.12 – 8.03
Figure 3: Volume shrinkage of products during sintering: a) T = 1100
°C, b) T = 1200 °C
Slika 3: Kr~enje prostornine proizvoda med sintranjem: a) T = 1100
°C, b) T = 1200 °C
Figure 1: X-ray diffractogram of red mud
Slika 1: Rentgenski difraktogram rde~ega blata
Figure 2: X-ray diffractogram of "Bijelo Polje" clay
Slika 2: Rentgenski difraktogram gline "Bijelo Polje"
The highest values of the compressive strength
(Figure 5) are present for the samples with the lowest
content of red mud. The samples sintered at the
temperature of 1200 °C have significantly higher values
of compressive strength than those sintered at the
temperature of 1100 °C.
The X-ray analysis results (Figures 6, 7, 8 and 9)
show different relations for the crystalline phases at the
different sintering temperatures. The Fe2O3 remains
M. KRGOVI] et al.: INFLUENCE OF TEMPERATURE AND BINDER CONTENT ON THE PROPERTIES ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 559–562 561
Figure 9: X-ray diffractogram of sintered product (w = 40 % red mud;
T = 1200 °C)
Slika 9: Rentgenski difraktogram sintranega proizvoda (w = 40 % rde-
~ega blata; T = 1200 °C)
Figure 6: X-ray diffractogram of sintered product (w = 30 % red mud;
T = 1100 °C)
Slika 6: Rentgenski difraktogram sintranega proizvoda (w = 30 % rde-
~ega blata; T = 1100 °C)
Figure 4: Total porosity of sintered products: a) T = 1100 °C, b) T =
1200 °C
Slika 4: Skupna poroznost sintranih proizvodov: a) T = 1100 °C, b) T
= 1200 °C
Figure 5: Compession strenght of sintered products: a) T = 1100 °C,
b) T = 1200 °C
Slika 5: Tla~na trdnost sintranih proizvodov: a) T = 1100 °C, b) T =
1200 °C
Figure 8: X-ray diffractogram of sintered product (w = 40 % red mud;
T = 1100 °C)
Slika 8: Rentgenski difraktogram sintranega proizvoda (w = 40 % rde-
~ega blata; T = 1100 °C)
Figure 7: X-ray diffractogram of sintered product (w = 30 % red mud;
T = 1200 °C)
Slika 7: Rentgenski difraktogram sintranega proizvoda (w = 30 % rde-
~ega blata; T = 1200 °C)
mostly unbound (which is more evident at the lower
sintering temperature). This also influences the different
values of the total porosity and the microstructure
changes of the sintered product, which has an effect on
the values of the compressive strength.
4 CONCLUSION
On the basis of the performed examinations the in-
fluence of the sintering temperature and the binder con-
tent on the properties of a sintered product containing
red mud we can conclude:
• red mud as a component of a raw material mixture
can be used in ceramic technology (for bricks,
blocks, etc.),
• the most important factors that determine the pro-
perties of the sintered products based on red mud are:
the amount of red mud in the raw material mixture,
the granulometric content of the components of the
raw material mixture and the sintering temperature,
• at the higher sintering temperature (1200 °C) more
favourable results regarding the total porosity and the
compressive strength are obtained,
• the differences in the values of the total porosity, the
volume shrinkage during sintering and the com-
pressive strength, with an increase of the red mud in a
raw material mixture (up to w = 50 %) are not
significantly higher.
5 REFERENCES
1 Lj. Kosti} - Gvozdenovi}, R. Ninkovi}, Inorganic Technology, Facul-
ty of Technology and Metallurgy, University of Belgrade, Belgrade
1997
2 M. Krgovi}, Z. Ja}imovi}, R. Zejak, Tile & Brick Int., 17 (2001) 3,
178–181
3 M. Krgovi}, M. Ivanovi}, N. Blagojevi}, @. Ja}imovi}, R. Zejak, M.
Kne`evi}, The Influence of the Mineral Content and Chemical Com-
position of Illite-Kaolinite Clays on the Properties of the Properties
of the Sintered Product, Interceram, 55 (2006) 2, 104–106
4 M. Krgovic, R. Zejak, M. Ivanovic, M. Vukcevic, I. Boskovic, M.
Knezevic, B. Zlaticanin, Properties of the Sintered Product based on
Electrofilter Ash depending on the Mineral Content of Binder,
Research Journal of Chemistry and Environment, 15 (2011) 4, 52–56
5 S. \urkovi}, The Possibilities of using Electrofilter Ash TE "Pljev-
lja" as Raw Materials component Mixture for producting sintered
product, Master Thesis, University of Montenegro, Podgorica, 2008,
4–5
6 M. Krgovi}, N. Marstijepovi}, M. Ivanovi}, R. Zejak, M. Kne`evi},
S. \urkovi}, The Influence of Illite-Kaolinite Clays’ Mineral Content
on the Products’ Shrinkage during drying and firing, Mater. Tehnol.,
41 (2007) 4, 189–192
7 B. @ivanovi}, R. Vasi}, O. Janji}, Ceramic Tiles, Monography, Insti-
tute of Materials in Serbia, Belgrade 1985, 12–17
8 M. Krgovi}, M. Kne`evi}, M. Ivanovi}, I. Bo{kovi}, M.Vuk~evi}, R.
Zejak, B. Zlati~anin, S. \urkovi}, Influence of the Temperature on
the Properties of the Sintered product on the basis of Electrofilter ash
as a component of the Raw Material Mixture, The Eleventh Annual
Conference "YUCOMAT 2009", Herceg Novi, 2009, 150–151
9 M. Krgovi}, M. Kne`evi}, M. Ivanovi}, I. Bo{kovi}, M.Vuk~evi}, R.
Zejak, B. Zlati~anin, S. \urkovi}, The Properties of a Sintered Pro-
duct based on Electrofilter Ash, Mater. Tehnol., 43 (2009) 6,
327–331
10 M. Tecilazi}-Stevanovi}, Principles of Ceramic Technology, Faculty
of Technology and Metallurgy, University of Belgrade, Belgrade
1990, 141–152
M. KRGOVI] et al.: INFLUENCE OF TEMPERATURE AND BINDER CONTENT ON THE PROPERTIES ...
562 Materiali in tehnologije / Materials and technology 48 (2014) 4, 559–562
R. GANESH, K. CHANDRASEKARAN: INFLUENCE OF THE PRECIPITATION TEMPERATURE ON THE THRUST FORCE ...
INFLUENCE OF THE PRECIPITATION TEMPERATURE
ON THE THRUST FORCE AND TORQUE IN DRILLING
AN Al 2219-SiCp COMPOSITE
VPLIV TEMPERATURE IZLO^EVALNEGA @ARJENJA NA POTISNO
SILO IN NAVOR PRI VRTANJU KOMPOZITA Al 2219-SiC
Radhakrishnan Ganesh1, Kesavan Chandrasekaran2
1Department of Mechanical Engineering, Anna University, 600025 Chennai, Tamilnadu, India
2Department of Mechanical Engineering, R.M.K Engineering College, 601206 Kavaraipettai, Tamilnadu, India
ganesh_akmcmc@yahoo.co.in, kesavan.chandrasekaran@gmail.com
Prejem rokopisa – received: 2013-06-10; sprejem za objavo – accepted for publication: 2013-09-13
This paper aims at understanding the influence of the precipitation temperature on the drilling performance of an Al 2219-SiC
composite. This type of composite is currently in use in aerospace industries. Drilling is the most frequently employed secon-
dary machining owing to the requirements of the fabrication. This composite was subjected to the precipitation heat treatment,
hence calling for a study of the precipitation-temperature effect on the machining performance. The composite was fabricated
through the powder-metallurgy route. The drilling trials were conducted with a PCD drill of a 5 mm diameter and 118° point
angle. The experiments were conducted with different spindle speeds of (500, 1000 and 1500) r/min and feed rates of (10, 15
and 20) mm/min. The thrust force and torque were the performance indicators. The influences of the mass fraction of
reinforcement, the precipitation temperature, the feed rate and the spindle speed on the thrust force and torque during drilling
were analyzed. It was observed that the drilling performance was influenced significantly by the feed rate followed by the
precipitation temperature, the mass fraction and the particle size of reinforcement. The spindle speed did not influence much the
machining performance.
Keywords: powder metallurgy, polycrystalline diamond, precipitation temperature
Namen ~lanka je ugotoviti vpliv temperature izlo~evalnega `arjenja na zmogljivost vrtanja kompozita Al 2219-SiC. Ta vrsta
kompozita se uporablja v letalski industriji. Zaradi potreb pri izdelavi je vrtanje najbolj pogosto uporabljena sekundarna
operacija obdelave. Ta kompozit je bil toplotno obdelan z izlo~evalnim `arjenjem, zato je potrebna {tudija vpliva temperature
izlo~evalnega `arjenja na zmogljivosti pri obdelavi. Kompozit je bil izdelan po postopkih pra{ne metalurgije. Preizkusi vrtanja
so bili izvr{eni s PCD-svedrom premera 5 mm in s konico pod kotom 118°. Preizkusi so bili izvr{eni pri razli~nih hitrostih
vrtenja vretena (500, 1000 in 1500) r/min ter hitrostih podajanja (10, 15 in 20) mm/min. Pokazatelja zmogljivosti sta bila
potisna sila in navor. Analiziran je bil vpliv masnega dele`a delcev za utrjevanje, temperature izlo~evalnega `arjenja, hitrost
podajanja, potisna sila in navor. Opa`eno je, da na zmogljivost vrtanja najbolj vpliva hitrost podajanja, nato temperatura
izlo~evalnega `arjenja, masni dele` in velikost delcev SiC. Hitrost vrtenja vretena nima bistvenega vpliva na zmogljivost
obdelave.
Klju~ne besede: pra{na metalurgija, polikristalni diamant, temperatura izlo~evalnega `arjenja
1 INTRODUCTION
The ongoing research on materials over the past
decades has produced advanced materials with the
properties superior to conventional materials such as
metal-matrix composites (MMCs). Metal-matrix com-
posites (MMCs) are one of the widely known composite
groups due to their superior specific strength, stiffness,
wear and corrosion resistances. Silicon-carbide-parti-
culate (SiCp) reinforced aluminium-based MMCs are
among the most common and commercially available
MMCs due to their economical production.1 These low-
cost composites show their advantages in the applica-
tions requiring the ability to withstand a relatively high
temperature besides high specific strength, stiffness and
fatigue resistance, the typical features of the composite
materials in the applications such as drive shafts, cylin-
der block liners, pistons and automotives.2 However,
because of the hard reinforcement particles in the matrix,
MMCs cause very serious problems in machining.3 For
this reason efforts have been made to develop the near-
net-shape manufacturing of these products. However, for
joining and assembling, secondary machining processes
such as drilling, milling, etc., are the prime requirements.
Machining, especially drilling, of cast MMC parts is
done using diamond-coated tools4 causing a considerable
increase in the tool life. The published literature on the
machining of MMCs indicate that only polycrystalline-
diamond (PCD) tools allow good tool life as PCD is
harder than the reinforcement materials such as SiC or
Al2O3 and it mostly does not chemically react with the
workpiece material. Kilickap et al.4 examined the tool
wear and the machined-surface characteristics during
machining SiC-reinforced aluminium-matrix composites
with coated and uncoated tool materials for different
feeds and speeds. They concluded that a higher cutting
speed and a lower feed rate result in good surface
quality. While a higher drilling speed reduces the mecha-
nical, thermal shock affecting the machine, a low feed
rate maintains an adequate flank clearance to avoid
Materiali in tehnologije / Materials and technology 48 (2014) 4, 563–569 563
UDK 621.762:66.017 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(4)563(2014)
rubbing/sliding. Quan and Ye5 investigated the hardness
and residual stress of the SiC/Al composites in a surface
layer affected by the machining and found that cutting at
a larger removal rate increases the possibility for the ten-
sile residual stress in the machined surface layers of the
composites. A surface integrity study was done by
Suresh Kumar Reddy et al.6 during the end milling of
Al/SiC composites. It was concluded from the study that
the reinforcement in the composites enhances the machi-
nability in terms of a lower tendency to clog the cutting
tool than in the case of an unreinforced Al alloy. In
addition, the tendency of the reinforcement to spall off
can enhance the chip forcing tendency and the conse-
quent machining. Ciftci et al.7 proved that cBN is unsui-
table for machining the MMCs containing the SiC parti-
cles of the average size of 110 μm as there was a heavy
fracture of the cutting edge and nose. This could be due
to the defects caused when the cBN was inserted onto
the core substrate.8 The tribological properties and me-
chanical properties of the Al2O3-SiC reinforced Al com-
posites manufactured through the powder metallurgy
route were investigated by Unlu,9 Ramulu et al.10 con-
ducted a drilling of Al2O3–aluminium-based metal-ma-
trix composites, using different drills and found that
PCD drills outperformed all the other drills in terms of
the drilled-hole quality and the drilling forces encoun-
tered. Basavarajappa et al.11 studied the drilling of hybrid
metal-matrix composites on the basis of Taguchi
techniques and it was revealed that the dependent
variables are more significantly influenced by the feed
rate than by the speed. Brun and Lee12 investigated the
machinability of the volume fraction Al-40 % SiC com-
posite and concluded that only polycrystalline diamond
tools exhibit useful tool life. Quigley et al.13 studied the
factors affecting the machinability of the Al/SiC metal-
matrix composite using different tool materials and
found that the triple-coated carbide tool, having the top
layer made of TiN performed the best in terms of the
flank wear compared to the other materials. The litera-
ture is mostly concerned with the effect of the reinforce-
ment on machinability; the machine-specific tool/drill/
performance is also reported. Only limited data on the
significance of the precipitation temperature is available.
The present experimental study aims at investigating the
influence of the precipitation temperature, the mass
fraction of reinforcement, the size of reinforcement and
the machining parameters on the thrust force and torque
when drilling the aluminium-matrix composites
reinforced with SiC particles; the selected cutting tool
for drilling was a fine-grained PCD drill.14,15
2 EXPERIMENTATION
The details of the MMC used and the working con-
ditions are listed in Table 1.
The matrix and reinforcement material are blended
mechanically in a planetary ball mill in order to ensure a
uniform distribution of the reinforcement for all the mass
fractions and particle sizes. It is then cold compacted
uniaxially to form cylindrical specimens with a diameter
of 10 mm and length of 20 mm. The workpieces were
precipitated in a muffle furnace at three different
temperatures of 500 °C, 550 °C and 600 °C for 4 h, and
then, each time, rapidly quenched in normal water and
artificially aged for 10 h at 170 °C in the same furnace.
This precipitation heat treatment was done to study the
effect of the precipitation temperature on the machining
performance of MMCs. The experimental set up is
shown in Figure 1. The PCD drill used for the experi-
ment is shown in Figure 2. The performance indicators
such as the thrust force and torque were measured using
a two-component piezoelectric dynamometer and the
experiment was repeated three times for the same sample
R. GANESH, K. CHANDRASEKARAN: INFLUENCE OF THE PRECIPITATION TEMPERATURE ON THE THRUST FORCE ...
564 Materiali in tehnologije / Materials and technology 48 (2014) 4, 563–569
Figure 2: Photograph of a 5-mm-diameter PCD drill used for the
experimentation
Slika 2: Posnetek PCD-svedra s premerom 5 mm, uporabljenega pri
preizkusih
Table 1: Experimental detail
Tabela 1: Eksperimentalne podrobnosti
Workpiece
Al 2219-matrix composite reinforced with
SiC-cold-compacted-reinforcement particle
Size – 37 μm and 67 μm
mass fraction – 10 %, 15 % and 20 %
Drilling
machine
ARIX VMC 100 CNC drilling machine
Tool PCD drill – twist drill of  = 5 mm,118° point angle
Drilling
conditions
Speed – (500, 1000, 1500) r/min
Feed rate – (10, 15, 20) mm/min
Environment – Dry
Figure 1: Photograph of the experimental set up with the dynamo-
meter attachment
Slika 1: Posnetek sestave za preizkuse z dodanim dinamometrom
and the average values of the output parameters were
noted.
3 RESULTS AND DISCUSSION
3.1 Density and porosity
From Figure 3, showing 37 μm SiC, it is clear that
the structural density increases with the precipitation
temperature and also with the mass fraction of rein-
forcement. A trend change can be seen around 550 °C. In
the case of 67 μm SiC, a reduction in the density can
mostly be seen above 550 °C. The SiC density increases
up to 15 %, followed by a drop.
From Figure 4, showing 37 μm SiC, it is clear that
the porosity reduces with the precipitation temperature; a
visible reduction occurs above 550 °C. This comple-
ments the observation on density. In the case of 67 μm
SiC, the porosity tends to drop up to 550 °C, followed by
a rise (by 10–15 %). Above 550 °C it rises and falls by
20 %. This finding mostly complements the observation
on density.
3.2 Hardness and compressive strength
The observed values of the hardness and compressive
strength for different mass fractions and precipitation
temperatures are illustrated in Figures 5 and 6. For 37
μm SiC, the hardness of the MMC tends to rise up to 550
R. GANESH, K. CHANDRASEKARAN: INFLUENCE OF THE PRECIPITATION TEMPERATURE ON THE THRUST FORCE ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 563–569 565
Figure 5: Observed VHN for different precipitation temperatures and
mass fractions: a) 37 μm SiC, b) 67 μm SiC
Slika 5: Izmerjena VHN pri razli~nih temperaturah izlo~evalnega
`arjenja in masnih dele`ih: a) 37 μm SiC, b) 67 μm SiC
Figure 4: Observed porosity for different precipitation temperatures
and mass fractions: a) 37 μm SiC, b) 67 μm SiC
Slika 4: Izmerjena poroznost pri razli~nih temperaturah izlo~evalnega
`arjenja in masnih dele`ih: a) 37 μm SiC, b) 67 μm SiC
Figure 3: Observed density for different precipitation temperatures
and mass fractions: a) 37 μm SiC, b) 67 μm SiC
Slika 3: Izmerjena gostota pri razli~nih temperaturah izlo~evalnega
`arjenja in masnih dele`ih: a) 37 μm SiC, b) 67 μm SiC
°C, followed by a reduction. With the increasing mass
fraction of reinforcement, the hardness increases. The
compressive strength increases up to 550 °C, followed
by a drop. It also increases with the massfraction of
reinforcement. Considering the finer-sized SiC-parti-
culate reinforcement and precipitation temperature, the
density increases with the reduction in the porosity. With
a high temperature (above 550 °C) the rise in the density
is associated with the hardness or compressive strength.
With the rise in the temperature, the matrix densifies
heavily. This is associated with a reduction in the flow
strength and the consequent densification but also a
reduction in the hardness.
The tendency to agglomerate is also lower. With 37
μm SiC, unlike in the case of 67 μm SiC, the MMC exhi-
bits a reduction in the density, especially at the tempera-
ture above 550 °C. A high mass-fraction (15 % and
20 %) reinforcement also exhibits only a marginal diffe-
rence. The porosity of the MMC tends to drop and rise
beyond 550 °C. With the increasing reinforcement
percentage, a mixed influence on the porosity (especially
with the 15 % and 20 % reinforcements) is seen. With
the coarser particulates and the increasing precipitation
temperature, a great portion of the heat flux could have
been absorbed by SiC resulting in a lower matrix flow
and the consequent increase in the porosity. However, the
MMC exhibits a rise in the hardness beyond 550 °C due
to the agglomeration or pooling of reinforcement, the
lower matrix flow and increased hardness, also reflected
in the increased compressive strength.
3.3 Thrust force and torque
The drilling performance was evaluated in terms of
two indicators: the thrust force and torque. The drill
thrust force and torque were monitored for all the trials.
The effects of the drilling conditions, the precipitation
temperature and reinforcement percentage, on the dril-
ling indicators are illustrated in Figures 7 to 10. It is
R. GANESH, K. CHANDRASEKARAN: INFLUENCE OF THE PRECIPITATION TEMPERATURE ON THE THRUST FORCE ...
566 Materiali in tehnologije / Materials and technology 48 (2014) 4, 563–569
Figure 7: Observed thrust force during the drilling of 37 μm SiC reinforced composites for different mass fractions of reinforcement, preci-
pitation temperatures, cutting speeds and feeds
Slika 7: Izmerjena potisna sila pri vrtanju kompozita, utrjenega s 37 μm SiC, pri razli~nih masnih dele`ih, temperaturah izlo~evalnega `arjenja,
hitrostih rezanja in podajanja
Figure 6: Observed compressive strength for different precipitation
temperatures and mass fractions: a) 37 μm SiC, b) 67 μm SiC
Slika 6: Izmerjena tla~na trdnost pri razli~nih temperaturah izlo~e-
valnega `arjenja in masnih dele`ih: a) 37 μm SiC, b) 67 μm SiC
seen that the thrust force increases with the feed rate and
speed for all the composites. An increase in the feed rate
results in a reduction in the effective clearance on the
flank surface (the point region) and the consequent rise
in the sliding and thrust force. The observed rise in the
thrust force with the increasing drill speed could be
attributed to a possible thermal softening of the matrix
aluminium with the increasing temperature and the con-
sequent clogging of the drill. With regard to the precipi-
tation temperature, it is seen that with the increasing
precipitation temperature, the MMC encounters a
higher-order thrust force. It is also clear that with a small
percentage of reinforcement a distinct influence of a
higher precipitation temperature (600 °C) on the thrust
force can be observed. With the increasing percentage of
reinforcement, the influence of the precipitation tempe-
rature is more or less evenly spread. With the increasing
precipitation temperature, the microstructure becomes
more densified, resulting in the observed rise in the
thrust force.
Unlike in the case of the MMC with the small-sized
SiC particle (37 μm), the MMC reinforced with the
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Materiali in tehnologije / Materials and technology 48 (2014) 4, 563–569 567
Figure 8: Observed thrust force during the drilling of 67 μm SiC reinforced composites for different mass fractions of reinforcement,
precipitation temperatures, cutting speeds and feeds
Slika 8: Izmerjena potisna sila pri vrtanju kompozita, utrjenega s 67 μm SiC, pri razli~nih masnih dele`ih, temperaturah izlo~evalnega `arjenja,
hitrostih rezanja in podajanja
Figure 9: Observed torque during the drilling of 37 μm SiC reinforced composites for different mass fractions of reinforcement, precipitation
temperatures, cutting speeds and feeds
Slika 9: Izmerjen navor pri vrtanju kompozita, utrjenega s 37 μm SiC, pri razli~nih masnih dele`ih, temperaturah izlo~evalnega `arjenja, hitrostih
rezanja in podajanja
coarse SiC particle (67 μm) exhibits a distinctly different
trend. The MMC with 67 μm SiC mostly encounters a
smaller-order thrust force, barring a higher mass frac-
tion. Unlike in the case of the 37 μm SiC reinforcement,
the MMC reinforced with the coarser particle, en-
counters a visible reduction in the thrust force with a
high feed rate. Also, with the precipitation temperature, a
relatively higher-order thrust force is encountered.
Unlike in the case of the 37 μm SiC reinforcement, the
MMC with the coarser particulate reinforcement encoun-
ters a reduction in the thrust force with the increasing
drilling speed. With the coarser reinforcement, the MMC
experiences a pooling of reinforcement, resulting in a
higher heterogeneity of the structure. With the coarser
particulate, the heat partition induces a great heat
absorption by the SiC reinforcement, resulting in a more
porous and less densified structure. The observed rise in
the thrust force for the 20 % reinforcement is attributed
to the pooling of reinforcement/higher-order heteroge-
neity, causing an excessive wear of the drill and the con-
sequent rise in the thrust force.
With regard to the torque, for 37 μm SiC composites,
it increases with the feed rate from 10 mm/min to 15
mm/min and decreases by 20 mm/min for all the mass
fractions of reinforcement and precipitation tempera-
tures. On the other hand, for 67 μm SiC composites, it
shows an increasing trend for the 10 % and 15 % mass
fractions at all the cutting speeds and precipitation
temperatures, but also a decreasing trend for the higher
mass fraction of 20 %. The torque was also observed to
have an increasing trend with the feed rate for all the
mass fractions and precipitation temperatures when the
spindle speed was increased from 500 r/min to 1000
r/min; however, it dropped after 1000 r/min. The feed
rate is associated with the chip thickness. A high feed
rate results in a larger cross-sectional area of the
undeformed chip, therefore, increasing the value of the
thrust force and torque. It was also observed that the
thrust force was high for the composites precipitated at
600 °C and low for the ones precipitated at 500 °C. This
may be attributed to the fact that at the higher precipi-
tation temperature, the matrix material becomes soft and
increases the area of contact between the tool and the
workpiece. At the higher spindle speeds, there was a
slight drop in the torque because of a decrease in the
contact area between the tool and the workpiece. With
R. GANESH, K. CHANDRASEKARAN: INFLUENCE OF THE PRECIPITATION TEMPERATURE ON THE THRUST FORCE ...
568 Materiali in tehnologije / Materials and technology 48 (2014) 4, 563–569
Figure 10: Observed torque during the drilling of 67 μm SiC reinforced composites for different mass fractions of reinforcement, precipitation
temperatures, cutting speeds and feeds
Slika 10: Izmerjen navor pri vrtanju kompozita, utrjenega s 67 μm SiC, pri razli~nih masnih dele`ih, temperaturah izlo~evalnega `arjenja,
hitrostih rezanja in podajanja
Figure 11: Typical profiles of: a) torque, b) thrust force
Slika 11: Zna~ilen profil: a) navora, b) potisne sile
67 μm SiC composites, the trend for the thrust force and
torque was similar to that of 37 μm SiC composites, but
the thrust force drops at the higher feed rate of 20
mm/min for all the spindle speeds. This may be due to
the larger particle size of reinforcement. The larger
particle size results in a poor bonding between the matrix
and the reinforcement. At the higher mass fraction and
higher spindle speeds, the torque decreases with an
increase in the feed rate due to an insufficient bonding
between the matrix and reinforcement. In general, for
both 37 and 67 SiC reinforced composites, with an
increase in the spindle speed, the thrust force and the
torque, there is a fall-and-rise trend. This may be initially
due to the reduction in the material hardness with the
significant increase in the temperature, whereas later the
softening of the matrix increases the contact area of the
matrix/reinforcement interface, thereby increasing the
thrust force and torque. The typical profiles of the thrust
force and torque obtained with the data-acquisition
system connected to the dynamometer is illustrated in
Figure 11.
4 CONCLUSION
The thrust force and torque results of the experimen-
tal study of the influence of the precipitation temperature
during the drilling of the Al 2219-SiCp composite mate-
rial using a PCD drill indicated the following conclu-
sions:
The thrust force increases with the feed rate at all the
cutting speeds, mass fractions of reinforcement and
precipitation temperatures for 37 μm SiC composites, but
it shows an increasing and decreasing trend for 67 μm
SiC composites. As the mass fraction of reinforcement
increased, the thrust force also increased significantly.
The thrust force was highly dependent on both the
spindle speed and the feed rate, showing an increasing
trend in the same way, regardless of the mass fractions of
reinforcement and precipitation temperature. The size of
the reinforcement particle has an appreciable influence
on the thrust force under different experimental condi-
tions. The precipitation temperature also has a significant
influence as the lowest thrust force was observed at the
higher precipitation temperature of the specimens.
The torque follows the rise-and-fall trend of the feed
rate for 37 μm SiC composites, whereas the increasing
trend is observed for 67 μm SiC composites. The torque
was observed to be the lowest at the lower feed rate and
for the higher mass fraction regardless of the sizes of
reinforcement. The trend of the torque remains the same
for all the precipitation temperatures. The size of the
reinforcement particle has an appreciable influence on
the torque under different experimental conditions,
whereas the precipitation temperature does not have a
significant influence on the torque.
5 REFERENCES
1 F. Bedir, B. Ogel, Proceedings of the 11th International Conference
on machine design and production, Turkey, 2004
2 M. Sklad El-Gallab, Journal of Materials Processing Technology, 83
(1998), 151–158
3 S. Jahanmir, M. Ramulu, P. Koshy, Machining of Ceramics and
Composites, Marcel Dekker, New York 2000
4 E. Kilickap, O. Cakir, M. Aksoy, A. Inan, Journal of Materials Pro-
cessing Technology, 164–165 (2005), 862–867
5 Y. M. Quan, B. Y. Ye, Journal of Materials Processing Technology,
138 (2003), 464–467
6 N. Suresh Kumar Reddy, S. Kwang-Sup, M. Yang, Journal of
Materials Processing Technology, 201 (2008), 574–579
7 I. Ciftci, M. Turker, U. Seker, Wear, 257 (2004), 1041–46
8 B. Anand Ronald, L. Vijayaraghavan, R. Krishnamurthy, Materials
and Design, 30 (2009), 679–686
9 B. S. Unlu, Materials and Design, 29 (2008), 2002–2008
10 M. Ramulu, P. N. Rao, H. Kao, Journal of Materials Processing
Technology, 124 (2002), 244–254
11 S. Basavarajappa, G. Chandramohan, J. Paulo Darim, Journal of
Materials Processing Technology, 196 (2008), 332–338
12 M. K. Brun, M. Lee, Wear, 104 (1985), 21–29
13 O. Quigley, J. Monaghan, P. O. Reilly, Journal of Materials
Processing Technology, 43 (1994), 21–36
14 T. S. Mahesh Babu, M. S. Aldrin Sughin, N. Muthukrishnan, Pro-
cedia Engineering, 38 (2012), 2617–2624
15 C. Dhavamani, T. Alwarsamy, Procedia Engineering, 38 (2012),
1994–2004
R. GANESH, K. CHANDRASEKARAN: INFLUENCE OF THE PRECIPITATION TEMPERATURE ON THE THRUST FORCE ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 563–569 569

V. MANOJLOVI] et al.: OPTIMIZATION OF THE RECYCLING PROCESSES FOR MAGNESIUM ...
OPTIMIZATION OF THE RECYCLING PROCESSES FOR
MAGNESIUM FROM A HIGHLY CONTAMINATED
WASTE
OPTIMIRANJE POSTOPKA RECIKLIRANJA MAGNEZIJA IZ
MO^NO KONTAMINIRANIH ODPADKOV
Vaso Manojlovi}1, @eljko Kamberovi}2, Miroslav Soki}1, Zvonko Guli{ija1,
Vladislav Matkovi}1
1Institute for Technology of Nuclear and Other Mineral Raw Materials, Bulevar Fran{ d’Eperea 86, 11000 Belgrade, Serbia
2Faculty of Technology and Metallurgy, University of Belgrade, Karnegijeva 4, 11000 Belgrade, Serbia
v.manojlovic@itnms.ac.rs
Prejem rokopisa – received: 2013-07-30; sprejem za objavo – accepted for publication: 2013-10-04
With the recently increased use of magnesium alloys, especially in the automotive industry, a large quantity of generated waste,
based on magnesium alloys, is expected. Such a waste often contains many impurities such as oil, paint, moisture, non-metallic
fractions, oxides, Cu, Fe, etc. In this paper, two different methods for extracting magnesium from a highly contaminated waste
are presented: the recycling of magnesium with flux and with the vacuum-distillation process. In addition, we present the
process of pre-treating the contaminated magnesium waste that has proved to be a very important step in the recycling process
for both economic and environmental reasons. The processing of post-consumer and contaminated magnesium scrap is
described with a diagram, providing a sustainable and environmentally friendly procedure for the treatment of such wastes.
Keywords: recycling, magnesium waste, vacuum distillation, flux, contaminated, impurities
Z nara{~anjem uporabe magnezijevih zlitin, posebno v avtomobilski industriji, se pri~akujejo velike koli~ine odpadkov na
osnovi magnezijevih zlitin. Taki odpadki pogosto vsebujejo {tevilne ne~isto~e, kot so olja, barve, vlaga, nekovinske frakcije,
oksidi, Cu, Fe in drugo. V tem ~lanku sta predstavljeni dve razli~ni metodi za ekstrakcijo magnezija iz mo~no onesna`enih
odpadkov: recikliranje magnezija s talilom in z vakuumsko destilacijo. Dodatno je prikazan {e postopek predobdelave kontami-
niranih magnezijevih odpadkov, ki se je izkazal za pomembnega pri recikliranju, tako iz ekonomskih kot tudi okoljevarstvenih
razlogov. Obdelava iztro{enega in kontaminiranega odpadnega magnezija je opisana z diagramom, ki zagotavlja trajnosten in
okolju prijazen postopek obdelave takih odpadkov.
Klju~ne besede: recikliranje, magnezijevi odpadki, vakuumska destilacija, talilo, kontaminiran, ne~isto~e
1 INTRODUCTION
Magnesium alloys, being the lightest metals, are used
widely for many parts in the automotive industry.1 Such a
trend is mainly a consequence of increased fuel prices,
leading to a decrease in the total mass of a vehicle. Fur-
thermore, many environmental regulations require the
use of lightmass materials in transport to reduce harmful
emissions. The European guideline on car usage has en-
forced the trend of using parts of lightmass metals.2,3
Due to their poor corrosion resistance, magnesium
automotive components exposed to a corrosive atmo-
sphere must be coated (paint and lacquer). Oily and
coated magnesium car components, at the end of their
life cycle, make the recycling process very difficult be-
cause of the contamination of magnesium melts with a
lot of impurities. There are several characterizations of
magnesium scrap, but the most referenced is the one
using eight classes of scrap. Class 1 of magnesium scrap
represents high-grade clean scrap without any impurities;
Class 2 is clean scrap with aluminium or steel inserts;
Class 3 includes clean, dry and uncontaminated turnings
and swarfs; Class 4 includes flux-free residues such as
dross and sludge; Class 5 refers to painted coated scrap
with/without aluminium or steel inlays and with no
copper or brass impurities; Class 6 includes oily and/or
wet turnings, swarfs, etc.; Class 7 includes unclean and
contaminated metal scrap, e.g., the post-consumer scrap
that may contain silicon (Al-alloys, shot blasting), Cu-
contaminated alloys, iron inserts, Ni-coatings and non-
magnesium sweepings; Class 8 includes the flux-con-
taining residues from Mg recycling, high amounts of
oxides, chlorides, fluorides (Mg amount < 30 %) and
Fe.4–6
Only the first-class scrap is relatively easy to recycle.
The recycling difficulties increase with the increase in
the class number of magnesium scrap. At present, only a
small amount of magnesium-alloy scrap is recycled (only
one third), while the rest is burned or buried in the
ground.3,7,8 Magnesium-alloy chips represent a big recy-
cling problem, caused by a very large specific surface
covered with MgO. Reactions of the chips with humidity
lead to a considerable oxidation, resulting in a loss of a
lot of metal,9,10 On the other hand, the accompanying
exothermic reaction can cause fire (the ignition tempe-
rature is 450 °C) during which the chips burn at 1000 °C
and they are difficult to extinguish. Hydrogen is liberated
Materiali in tehnologije / Materials and technology 48 (2014) 4, 571–575 571
UDK 658.567.1:669.721.5.004.8 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(4)571(2014)
during this reaction causing small explosions and
entailing another safety problem.1,5
The elements such as Ni, Cu, Fe, and Si have a
crucial influence on decreasing the magnesium-alloy
corrosion resistance, so the amount of these impurities
must be limited to the ppm level. It is very difficult to
recycle "old" scrap (especially from a shredder plant),
because it always contains a certain amount of these
impurities.4,9,11–13 A removal of Ni and Cu from a
magnesium melt with additives is practically impossible
due to strong interactions in the Mg-Ni and Mg-Cu
systems.6,9 Only the distillation process is an effective
way to separate Ni and Cu from magnesium alloys.
The recycling of magnesium must be, above all,
economically reasonable. The average price of primary
magnesium metal (99.9 %) has been about 3.1 $/kg in
the past five years.14 Different recycling methods for
magnesium scrap have been developed. The most com-
mon method for recovering clean scrap (classes 1, 2 and
3) involves flux. The recycling of magnesium scrap clas-
ses 4, 5, 6, 7 and 8 is not practical due to the increased
impurities (Cu, Ni, Fe, Si) and the dirtiness of the
starting material that decrease the economic efficiency of
the recycling process. The most efficient process for the
treatment of these classes of waste is the vacuum-distil-
lation process.5,6,10
2 EXPERIMENTAL PART
The first part of the experiment covers the most com-
mon method of magnesium recovery, the recycling with
flux. A pre-treatment was conducted in order to improve
both the economy and the technology of the process. The
second part of the experiment refers to the most efficient
method for recycling contaminated magnesium scrap, the
vacuum-distillation process.
In the experiments with the flux treatment of the
waste, two types of the samples were used. The first type
of the samples mostly included the parts from the
automotive industry and the casting process without any
oil contaminations, with different thicknesses of the
oxide layer on the surface as well as different sizes and
shapes. The second type of the samples included oil-con-
taminated parts with oxides, dust and other inorganic
impurities. The dimensions and descriptions of the sam-
ples are shown in Table 1.
The pre-treatment of samples 1 and 2 was conducted
in the following four steps:
• mechanical treatment in a rotating drum;
• dust separation;
• chemical treatment in a solution of sulfuric acid;
• washing and drying.
The mechanical treatment was performed in a labo-
ratory rotating drum with a volume of 6 dm3 and rotation
of 150 r/min. In this drum, the light fraction containing
dust was separated from the massive fraction. The mass
of the solid samples in the drum was (2000 ± 10) g. The
chemical treatment of the massive fraction was conduc-
ted at the temperatures from 45–50 °C with an addition
of a solution with mass fractions, w = 5–10 % H2SO4.
After the chemical treatment, the metal fraction in the
lattice basket was washed with water. The drying process
was performed in a furnace at the temperature of about
100–200 °C for a period of 6–8 h.
The smelting of prepared Samples 1 and 2 was con-
ducted in an induction 25 kVA furnace with a graphite
crucible of 1 dm3 and with additional mechanical stirr-
ing. A commercial smelting-salt mixture with the follow-
ing composition was used: 76 % emgesal, 18 % dekamag
and 6 % CaF2. Since Sample 3 was contaminated with
oil, the degree of oil contamination had to be determined.
The oil removal was carried out with a prepared water
solution of water glass (sodium silicate, 18.3 %), dode-
V. MANOJLOVI] et al.: OPTIMIZATION OF THE RECYCLING PROCESSES FOR MAGNESIUM ...
572 Materiali in tehnologije / Materials and technology 48 (2014) 4, 571–575
Table 1: Characteristics and dimensions of the samples
Tabela 1: Zna~ilnosti in dimenzije vzorcev
Sample Characteristics Dimensions, (a · h) mm2
Without oil
(Classes 4 and 7)
1 Compact parts from the auto industry and parts of a gatingsystem. A thin oxide layer is noticeable. from 15 · 20 to 90 · 200
2 Spongy, loose parts; mostly casting slag. A 1–3 mm oxide layeris noticeable on the surface. from 20 · 30 to 30 · 100
Oil contaminated
(Class 6) 3
Irregular biconcave parts with oil, oxides, dust and other
inorganic impurities. from 10 · 2 to 60 · 5
Figure 1: Vacuum-distillation furnace
Slika 1: Vakuumska destilacijska pe~
cylbenzenesulfonate (DBS, 6.7 %) and sodium hydro-
xide (NaOH, 5.0 %).
For the processing of the contaminated magnesium
scrap with the vacuum-distillation process, highly con-
taminated magnesium scrap was used, belonging to
Classes 5, 6 and 7. The waste was briquetted and its con-
tent was determined with a chemical analysis. The
vacuum-distillation furnace is shown in Figure 1.
A batch of (400 ± 5) kg was melted in the electric-
resistance furnace with two chambers under the vacuum
of 0.5 bar. In the lower part of the furnace, the tempe-
rature was around 840–950 °C and in the upper part it
was from 650–680 °C. The temperature of the condenser
was about 135 °C. The times of the distillation process
for magnesium from the waste were from 24–48 h.
3 RESULTS AND DISCUSSION
The processing of the magnesium waste, of Classes
4, 5, 6 and 7, was conducted with the flux-recycling pro-
cess and the vacuum-distillation process. In the first part
of the experiment, the importance of the waste pre-treat-
ment was shown. The smelting process of Samples 1 and
2 was conducted with different pre-treatment combina-
tions in order to indicate how each step affects the
magnesium recovery rate. In addition, smelting salt was
added in different amounts. The results are shown in
Table 2.
Table 2: Magnesium recovery rate of Samples 1 and 2 in dependence
of different combinations of the pre-treatment process and smelting
salt
Tabela 2: Stopnja pridobitve magnezija v vzorcih 1 in 2 pri razli~nih
kombinacijah predobdelave in talilne soli
Sam-
ple Pre-treatment
Smelting
salt, g
Mg
recovery
rate, %
1
Without pre-treatment, direct
smelting – 71 ± 2.5
Mechanical treatment, 1 h, dust
separation 50 91 ± 0.5
Chemical treatment with 5%
H2SO4, drying at 100 °C
50 81 ± 1.5
Mechanical treatment, 1 h, dust
separation, chemical treatment
with 5 % H2SO4, drying at 100 °C
50 94 ± 0.8
2
Mechanical treatment, 1 h, dust
separation – 44 ± 1.5
Mechanical treatment, 1 h, dust
separation 60 47 ± 1.5
Chemical treatment with 5 %
H2SO4, drying at 100 °C
60 46 ± 1.5
Mechanical treatment, 1 h, dust
separation, chemical treatment
with 5 % H2SO4, drying at 100 °C
60 49 ± 1.5
By processing Sample 1 (belonging to Class 7), it
was shown that it is possible to achieve a utilization of
up to 94 % with a pre-treatment. Without the pre-treat-
ment, the utilization is only 71 %, which indicates the
importance of the pre-treatment step in the recycling
process of this type of magnesium scrap. The maximum
utilization after the processing of Sample 2 (the Class 4
scrap) is about 49 %, which is a consequence of a rela-
tively thick oxide layer of 3 mm and a sponge structure
of the waste, i.e., a relatively high surface/volume ratio.
The Sample 2 pre-treatment does not significantly in-
crease the recovery of magnesium from the waste due to
the thick oxide layer.
The time for removing the oil from Sample 3, with a
constant stirring, was 5 min at the temperature of 60 °C,
while the sample mass was (1000 ± 10) g. The average
degree of oil contamination with powder impurities for
Sample 3 was (17 ± 0.5) %. It was shown that the opti-
mum share of the water solution for an efficient removal
of oil contamination (about (96 ± 1.5) %) is 10 %. When
oil was removed from the surface, the scrap was smelted
in an induction furnace. The total metal-recovery rate for
Samples 4 was (93 ± 1.8) %. The results for Sample 3
(the Class 6 scrap) showed an effective way to remove
oil from the waste, with a high recovery rate.
A pre-treatment of magnesium waste has a signifi-
cant influence on:
• the utilization of magnesium from the waste,
• the reduction of slag amounts,
• a better economy,
• a better ecology and safety of the process.
It should be noted that same experiments with the
flux were done in industrial conditions in the factory
Bela stena in Boljevac, Serbia, with the batches of the
total mass of 157 t (Samples 1, 2 and 3). The total utili-
zation of magnesium from the waste, in industrial con-
ditions, was (75 ± 0.5) % (118 t of magnesium).
In the second part of the experiment, the vacuum-
distillation process was employed for the extraction of
magnesium from the contaminated waste (Classes 5, 6
and 7). The share of impurities was about 10 % (7 %
moisture and 3 % oil, paint, oxides) and the rest was the
magnesium alloy (mostly from the automotive industry).
The experiments were conducted in different ranges of
temperature and time in order to obtain the optimum
parameters of the process. The results are shown in
Table 3.
Table 3: Parameters for the vacuum-distillation process
Tabela 3: Parametri pri vakuumskem destilacijskem procesu
Sample
Temperatu-
res down/
up the fur-
nace, °C
Time of
distillation
process, h
Metal in
crown, kg
Residue,
kg
Metal
recovery
rate, %
v1 840/650 24 200.2 170.5 55.5
v2 840/650 30 201.1 169.7 55.7
v3 840/650 40 203.6 170.3 56.2
v4 840/650 48 206.9 167.9 57.0
v5 900/670 24 207.5 162.4 57.4
v6 950/680 24 212.1 163.7 58.2
The chemical compositions of distilled magnesium
crowns are shown in Table 4. Since zinc has a lower
V. MANOJLOVI] et al.: OPTIMIZATION OF THE RECYCLING PROCESSES FOR MAGNESIUM ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 571–575 573
vapor pressure than magnesium, the zinc in a crown
required higher temperatures and longer times, but, in
general, high-purity magnesium was obtained.
Table 4: Chemical composition of the magnesium crown
Tabela 4: Kemijska sestava magnezijeve krone
Sample
Chemical composition of the magnesium distillate
Al Zn Mn Fe Si Cu Mg
v1 0.02 0.04 0.005 0.001 0.003 0.001 99.93
v2 0.03 0.05 0.003 0.001 0.003 0.001 99.91
v3 0.03 0.05 0.005 0.001 0.002 0.002 99.91
v4 0.02 0.08 0.002 0.002 0.004 0.001 99.89
v5 0.03 0.14 0.004 0.001 0.003 0.0005 99.82
v6 0.04 0.34 0.005 0.002 0.004 0.001 99.61
The waste in the form of pressed briquettes was con-
taminated with oil, oxides, paint and dust. The oxides of
magnesium and aluminium (MgO and Al2O3) contained
in a batch cannot be converted to their elementary state;
instead, they form a foamed slag along with the impu-
rities and inter-metallic compounds. The distillation resi-
due, besides the foamed slag, consists of a metal fraction
that has a chemical composition of 35 % Al, 0.2 % Zn,
1.2 % Mn, 0.03 % Fe, 0.05 % Si, and the rest is mag-
nesium and impurities. During the process, the amount
of the magnesium in the metallic fraction of the batch
was decreased and the amounts of the slag and alumi-
nium in the batch were increased. Consequently, the
utilization of magnesium is at a low level due to the
difficulties in the movement of magnesium atoms
through the oxides and alloys in the batch at the end of
the process.
According to the results from Tables 3 and 4, it can
be concluded that the duration of the process does not
significantly affect the extraction of magnesium from the
waste. However, by increasing the temperature, while
keeping the same duration of the process, the magnesium
recovery was increased by 2–3 %, which is a low rate
with respect to the increment of the energy consumption
of 7–8 %. In addition, with an increase in the operating
temperature, the impurity amount in the distillate
increased due to an intense evaporation of the impurities
from the batch surface.
The residue from the vacuum-distillation process can
be used for a recovery of aluminium, but a large amount
of impurities must be taken into consideration. Although
the recovery of magnesium from the waste is at a low
level (about 58 %), pure magnesium was obtained (up to
99.93 %), justifying the economy of the process. For the
industrial use, it is necessary to construct less expensive
and more mobile equipment for the vacuum-distillation
process. In addition, further research should focus on the
combination of pre-treatment and vacuum-distillation
process, which should increase the extraction of magne-
sium from the waste.
To remove an organic coating/paint and oil/moisture
the pre-treatment step was used as a very effective
method. The removal of these contaminations is essential
for decreasing the dross formation and harmful gas
exhaust and, thereby, reducing the loss of magnesium
through the dross. Besides, if not removed they can
significantly influence the vacuum-equipment lifetime
when there is a direct contact. If there is a filter in the
vacuum equipment for a disposal of harmful gases it will
increase the price of the investment and make it difficult
to create a vacuum, so a pre-treatment is a much better
solution.
In the post-consumer scrap, the main impurities such
as Ni, Cu, Fe and Si are always expected due to the com-
plexity of the waste structure, particularly in the shredder
process, where many materials are involved. In the
vacuum-distillation process, impurities are easily sepa-
rated from magnesium due to a different vapor pressure,
which is a noticeable advantage of this process. In the
process with flux, it is advisable to produce magnesium
pre-alloys for the aluminium-alloy production rather than
neutralizing it with Zn or Mn, especially at higher levels
of Ni and Cu in the alloys (grater then 0.002 % for Ni
and 0.03 % for Cu).
A recommended scheme for processing post-consu-
mer and contaminated magnesium scrap is shown in
Figure 2.
In Figure 2, we can see that the by-products such as
oil, slag, dust, metallic oxides and non-metallic impuri-
ties are obtained. Dust, slug and non-metallic impurities
can be used as additions to asphalt and oil can be used as
a fuel. The Al-alloy from the vacuum-distillation process
can be used for steel desulphurization.
4 CONCLUSION
In this paper two different methods for recycling
magnesium waste were investigated: recycling with flux
and recycling with the vacuum-distillation process. Mag-
V. MANOJLOVI] et al.: OPTIMIZATION OF THE RECYCLING PROCESSES FOR MAGNESIUM ...
574 Materiali in tehnologije / Materials and technology 48 (2014) 4, 571–575
Figure 2: Processing of post-consumer and contaminated magnesium
scrap
Slika 2: Obdelava izrabljenih in kontaminiranih odpadkov magnezija
nesium waste with oil, moisture and other non-metallic
impurities was used in these processes. On the basis of
the presented results, we can conclude that a pre-treat-
ment has a significant role in the recycling process. The
main advantages of the pre-treatment step applied to
contaminated magnesium waste in the recycling process
are the following:
• a significant increase in the recovery of magnesium
from the waste (up to 94 %);
• an improvement in the economy of the process;
• a reduced share of the slag in the flux-recycling pro-
cess;
• a removal of the moisture and oil from the waste
enables a proper storage of the waste, without any
safety risks;
• elimination of environmental problems (an emission
of harmful pollutants into the air is prevented);
• an improved diffusion throughout the batch in the
vacuum-distillation process, increasing the intensity
of the process;
• prevention of contamination and vacuum-equipment
damage.
In addition, a recommended scheme for processing
contaminated magnesium waste was given, whereby
several semi-products were obtained. In the case of high-
er proportions of Ni and Cu contained in a magnesium
alloy, it is recommended to use these alloys for alumi-
nium pre-alloying. If high-purity magnesium has to be
obtained, the vacuum-distillation process is inevitable.
By-products such as oil, slag, dust, metallic oxides and
non-metallic impurities can be used as a fuel and as
building-material components.
Acknowledgement
The authors wish to acknowledge the financial
support from the Ministry of Education and Science of
the Republic of Serbia within projects TR34002 and
TR34033.
5 REFERENCES
1 M. J. Fereiría Gándara, Mater. Tehnol., 45 (2011) 6, 633–637
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& Materials Engineering, 15 (2009), 189–200
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nesium Scrap, Proc. of the symp. jointly sponsored by the Magne-
sium Committee of the Light Metals Division of TMS with the Inter.
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Technologies of Magnesium, The Government of Canada, Enhanced
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6 G. Hanko, G. Macher, Technologies for Efficient Mg-Scrap Recycl-
ing, Proc. of the TMS Annual Meeting, Magnesium technology,
Warrendale, 2003, 29–32
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its Alloys, Proc. of Magnesium Alloys and Their Applications,
Wolfsburg, Germany, 1998, 685–690
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9 G. Hanko, H. Antrekowitsch, P. Ebner, Journal of Metals, 54 (2002)
2, 51–54
10 A. Gesing, A. Dubreuil, Recycling of Post-Consumer Mg Scrap,
Proc. of 65th Annual World Magnesium Conference, Warsaw, 2008,
131–141
11 W. Stevens, F. Trembay, Fundamentals of UBC Decoating / De-
lacquering for Efficient Melting, Proc. of the Technical Sessions
Presented by the TMS Aluminum Committee at the 126th Annual
Meeting, Orlando, Florida, 1997, 709–713
12 W. Green, J. Hillis, Process for Producing High Purity Magnesium,
Patent number US4891065, 1990
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tals, 41 (2000), 72–75
14 Infomine [online], 5 Year Magnesium Prices and Price Charts.,
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V. MANOJLOVI] et al.: OPTIMIZATION OF THE RECYCLING PROCESSES FOR MAGNESIUM ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 571–575 575

A. MAVI, I. KORKUT: MACHINABILITY OF A Ti-6Al-4V ALLOY WITH CRYOGENICALLY TREATED ...
MACHINABILITY OF A Ti-6Al-4V ALLOY WITH
CRYOGENICALLY TREATED CEMENTED CARBIDE
TOOLS
OBDELOVALNOST ZLITINE Ti-6Al-4V S KRIOGENSKO
OBDELANIMI ORODJI IZ KARBIDNE TRDINE
Ahmet Mavi1, Ihsan Korkut2
1Gazi University, OSTIM Vocational School, Mechatronics Programs, Ankara, Turkey
2Gazi University, Faculty of Technology, Manufacturing Engineering, Ankara, Turkey
amavi@gazi.edu.tr
Prejem rokopisa – received: 2013-09-09; sprejem za objavo – accepted for publication: 2013-09-20
In this study, the effects of the treatment (applied to cutting parameters and cutting tool) on the cutting forces, surface roughness
and tool life were investigated. Part of the cutting tools was subjected to a cryogenically treatment at –145 °C for 24 h. The
effects of the cryogenically treated and untreated cemented carbides (uncoated) on the cutting forces, surface roughness and
wear behaviors were investigated. Machinability tests were carried out at four different cutting speeds (30, 45, 60, 75) m/min,
three different feed rates (0.20, 0.25 and 0.30) mm/r and the 1 mm cutting depth. Wear tests were made for four different chip
volumes (20, 40, 60 and 80) cm3, four different cutting speeds (30, 45, 60, 75) m/min, the 0.25 m/r feed rate and the 1 mm
cutting depth. At the end of the tests, the cryogenically treated inserts gave better results compared to the untreated tools with
respect to the wear behavior, cutting forces and surface roughness.
Keywords: Ti-6Al-4V, tool wear, cryogenic process, cutting forces, surface roughness
V tej {tudiji je bil preiskovan vpliv obdelave (uporabljene za rezalne parametere in rezalno orodje) na sile pri rezanju, hrapavost
povr{ine in zdr`ljivost orodja. Del orodij iz karbidne trdine je bil kriogensko obdelan 24 h pri –145 °C. Preiskovan je bil vpliv
kriogensko obdelanih in neobdelanih orodij iz karbidne trdine (brez prevleke) na sile rezanja, hrapavost povr{ine in vedenje pri
obrabi. Preizkusi obdelovalnosti so bili izvr{eni pri {tirih razli~nih hitrostih rezanja (30, 45, 60, 75) m/min, pri treh razli~nih
hitrostih podajanja (0,20, 0,25 in 0,30) mm/r in pri globini rezanja 1 mm. Preizkusi obrabe so bili izvr{eni pri {tirih volumnih
odrezkov (20, 40, 60 in 80) cm3, pri {tirih hitrostih rezanja (30, 45, 60, 75) m/min, pri podajanju 0,25 m/r in pri globini rezanja 1
mm. Preizkusi so pokazali, da dajejo s stali{~a obrabe, sil pri rezanju in hrapavosti povr{ine, kriogensko obdelani vlo`ki bolj{e
rezultate v primerjavi z neobdelanim orodjem.
Klju~ne besede: Ti-6Al-4V, obraba orodja, kriogenski postopek, sile rezanja, hrapavost povr{ine
1 INTRODUCTION
Titanium and its alloys have high strength, heat and
corrosion resistance. They are used in the medical, elec-
tronics and computer, aviation and space industries.1,2
These alloys preserve their properties even at high tem-
peratures during machining. Therefore, titanium belongs
to the group of hard-to-machine materials. Among these
alloys with different properties, Ti-6Al-4V has the lion’s
share with a 60 % usage in industrial applications.3,4
While at the higher temperatures during the treatment, a
titanium alloy can preserve its strength, the cutting tool
loses its strength due to the high temperature and
pressure. The reaction with the cutting-tool material at
high temperatures and the build-up edges (BUE) on the
cutting-tool tip significantly affect the cost and efficiency
of the treatment.5–7 Siekman indicated that the machining
of these alloys is always problematic no matter which
classical method is used.8 From this aspect, an investi-
gation of suitable machining conditions is important for
the machining of these alloys.
Cryogenic treatment is a supplementary process for
the heat treatment applied to increase the wear resistance
of the materials subjected to high wear. It is also known
as the below-zero cryogenic treatment. Contrary to the
coatings, it is a cheap, long-lasting process; it is carried
out once and it affects the whole piece. With this
method, a conversion of the residual austenite into
martensite in a conventionally heat-treated material, a
formation of thin carbide precipitates and a uniform
carbide distribution are obtained. In this way, serious
improvements in the mechanical properties of the mate-
rials such as the hardness and wear resistance are
achieved. Cryogenic treatment used to be applied to
molding materials, but nowadays, due to its application
to the cutting tools in the machining practice, significant
developments have been made with respect to tool wear,
tool life and cutting conditions. It was found that by
applying cryogenic treatment to certain tool materials,
improvements in the tool wear by 91 % to 817 % were
achieved.9–12
In this study, during the turning of an Ti-6Al-4V
alloy under dry cutting conditions with cryogenically
treated and untreated, uncoated cemented carbide tips,
the effects of different combinations of the cutting depth,
cutting speed and feed rate on the cutting forces, surface
roughness and tool wear were examined.
Materiali in tehnologije / Materials and technology 48 (2014) 4, 577–580 577
UDK 669.295:621.9 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(4)577(2014)
2 EXPERIMENTAL PROCEDURES
2.1 Materials
In the tests, the Ti6Al4V alloy having the AMS 4928
characteristics was used. The chemical constituents and
the physical properties of the material are given in
Tables 1 and 2, respectively.
Table 1: Chemical composition of Ti-6Al-4V (mass fractions, w/%)
Tabela 1: Kemijska sestava Ti-6Al-4V (masni dele`i, w/%)
N C H Fe O Al V
0.08 2.00 0.75 0.045 0.03 16.0–18.0 10.0–14.0
Table 2: Mechanical properties of Ti-6Al-4V
Tabela 2: Mehanske lastnosti Ti-6Al-4V
Tensile strength Yield strength Hardness,Rockwell C Elongation
900–1100 MPa 830 MPa 36 10 %
2.2 Machining tests
The tests were made using a JohnfordTC35 CNC
turning center with no cooling liquid at four different
cutting speeds (30, 45, 60, 75) m/min, three different
feed rates (0.20, 0.25, 0.30) mm/r and a constant cutting
depth (1 mm). The wear tests were realized for four
different chip volumes (20, 40, 60 and 80) cm3, four
different cutting speeds (30, 45, 60 and 75) m/min, a
0.25 m/r feed rate and a 1 mm cutting depth. In the tests,
the uncoated cemented carbide tips of the SANDVÝK
Coromant company were used. Half of these tips were
cryogenically treated for the purpose of making a
comparison. A graphic presentation of the steps of the
cryogenic treatment is given in Figure 1.
The surface-roughness measurements on the
machined surfaces of the samples were made with a
Mahr Perthometer, type M1, surface-roughness-measure-
ment device. The measurements were made parallel to
the working-piece axis and on three different surfaces by
rotating a workpiece 120° around its axis after each
measurement. The average surface-roughness (Ra) values
were determined by taking the arithmetical mean of the
three values obtained at the end of the tests. The cutting
force was measured with a Kistler 9257A three-
component piezoelectric dynamometer and the asso-
ciated 5019 B130 charge amplifiers connected to a PC
employing the Kistler Dynoware force-measurement
software.
The measurement of the tool flank wear obtained at
the end of the turning operations was made with a scann-
ing electron microscope (SEM).
3 RESULTS AND DISCUSSIONS
3.1 Flank wear
Tool wear is the result of the frictions and tempe-
rature taking place in the areas where the cutting-tool
material and the workpiece material are in contact.
Friction is the most important reason for the wear,
whereas the temperature decreases the resistance to wear
and, for this reason, it is the factor accelerating the
wear.13 One of the most common types of the wear
occurring in machining is flank wear. Flank wear occurs
because of the contact between the side surfaces of a
cutting tool and the workpiece. In Figure 2, the forma-
tion of the maximum flank wear at different cutting
speeds for normal and cryogenically treated, cemented
carbide tools is given. The flank wear increased with the
increasing cutting speed (Figure 2) because the contact
area of the cutting tool increased with an increase in the
cutting speed and because the temperature was raised
due to the increasing cutting speed leading to flank
wear.14 The flank wear was lower for the cryogenically
treated inserts compared to the normal tools. The reason
for this was the fact that the cryogenically treated inserts
had more uniform and more stable microstructures
(Figure 3). The effects of more uniform structures on the
wear behavior are also known.9,15–17 When the wear
values of the cutting tools are considered it is seen that
the wear of the cryogenically treated inserts is lower by
10–22 %.
The tool wears depending on different cutting speeds
are given in Figures 4 and 5. From these figures it is
A. MAVI, I. KORKUT: MACHINABILITY OF A Ti-6Al-4V ALLOY WITH CRYOGENICALLY TREATED ...
578 Materiali in tehnologije / Materials and technology 48 (2014) 4, 577–580
Figure 1: Schematic diagram of the cryogenic treatment
Slika 1: Shematski prikaz kriogenske obdelave
Figure 2: Maximum width of the flank wear at different cutting
speeds
Slika 2: Najve~ja {irina obrabe boka pri razli~nih hitrostih rezanja
clear that the wear is higher for the untreated cutting
tools compared to the cryogenically treated inserts. In
Figures 4 and 5, it is observed that, at all the cutting
speeds, the flank wear is lower when the chip volume is
20 cm3 and 40 cm3. However, depending on the increas-
ing chip volume, there is a rapid increase in the flank
wear. For both cutting tools (cryogenically treated and
untreated) when the chip volume is 80 cm3 and the
cutting speed is 75 m/min, the flank wear is maximum.
The main reason for this is the fact that the high tempe-
ratures at the tool-chip interface depend on the increas-
ing cutting speed and the chip volume.18
3.2 Cutting forces
In Figure 6 the changes in the main cutting forces of
the cryogenically treated and untreated inserts are given.
It is clear that a decrease in the main cutting force is due
to the increasing cutting speed. It is understood that this
tendency is the same for both tools, as the power used in
the machining is usually converted into heat on the
sliding plane, at the circumference of the cutting tip.
Most of the heat created on the sliding plane is removed
with the chip, but a certain portion is conducted to the
workpiece. This conducted heat decreases the hardness
of the workpiece. When the hardness decreases, the duc-
tility increases making the chip removal easier. However,
if this conduction of the heat continues to increase, the
changes in the cutting forces can occur18–20 due to a start
of the BUE tendency. In the cryogenically treated inserts,
the main cutting forces are lower compared to the
untreated inserts. This situation can be explained with a
low heat at the cryogenically treated insert tip and a
lower flank wear. The best result with respect to the main
cutting forces was 75 m/min for the cryogenically treated
inserts.
A. MAVI, I. KORKUT: MACHINABILITY OF A Ti-6Al-4V ALLOY WITH CRYOGENICALLY TREATED ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 577–580 579
Figure 6: Main cutting force (Fc) at different cutting speeds
Slika 6: Sila rezanja (Fc) pri razli~nih hitrostih rezanja
Figure 4: Maximum width of the flank wear at different chip volumes
for the untreated insert
Slika 4: Najve~ja {irina obrabe bokov pri razli~nih volumnih
ostru`kov pri neobdelanem vlo`ku
Figure 3: Microstructures of cutting inserts: a) untreated insert, b)
cryogenically treated insert
Slika 3: Mikrostruktura vlo`ka za rezanje: a) neobdelan vlo`ek, b)
kriogensko obdelan vlo`ek
Figure 5: Maximum width of the flank wear at different chip volumes
for the cryogenically treated insert
Slika 5: Najve~ja {irina obrabe bokov pri razli~nih volumnih
ostru`kov pri kriogensko obdelanem vlo`ku
3.3 Surface roughness
The surface quality and roughness obtained with each
machining method changed. The correct selection of the
surface roughness directly affects the product cost.
Because of this, the surface roughness is one of the most
important machining parameters. As it is seen from
Figure 7, with the increasing cutting speed the surface-
roughness values decrease at first, but when the cutting
speed exceeds 60 m/min they are affected adversely,
exhibiting a sudden increase. The improvement in the
surface roughness with the increasing cutting speed,
depending on the increasing temperature at higher
speeds, can be explained with an easy deformation of the
workpiece material at the cutting edge and the circum-
ference of the tip radius, and the yielding formation at
high temperatures.19,20 However, as it is observed from
the wear graphics in Figures 4 and 5 and from the SEM
image of the cutting tool in Figure 8, the reason for the
rapid increase in the surface roughness at the 60 m/min
cutting speed is caused by the formation of the wear of
the cutting tools.
4 CONCLUSION
Cryogenically treated inserts gave better results than
the other tools with respect to tool wear.
The wear of the cryogenically treated inserts was
lower than the wear of the other inserts by 10–22 %.
At all the cutting speeds, cryogenically treated inserts
gave better results with respect to the main cutting
forces.
When the cutting speed was 30 m/min and 45 m/min,
the surface roughness decreased, but when it increased to
higher levels the average surface roughness increased
too.
Cryogenically treated inserts gave better results with
respect to the main cutting forces and the average surface
roughness.
The cryogenic treatment made the microstructure of
the cutting tool more uniform and more stable.
Acknowledgements
The authors would like to acknowledge the Gazi
University Scientific Research Project (07/2011-58) for
the financial support throughout this study.
5 REFERENCES
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2 E. O. Ezugwu, Z. M. Wang, Journal of Materials Processing Tech-
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3 H. A. Kishawy, J. Wilcox, International Journal of Machine Tools &
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Materials Processing Technology, 100 (2000), 1–11
19 I. Çiftçi, Journal of the Faculty of Engineering and Architecture of
Gazi University, 20 (2005) 2, 205–209
20 E. M. Trent, Metal cutting, Butterworths Press, London 1989
A. MAVI, I. KORKUT: MACHINABILITY OF A Ti-6Al-4V ALLOY WITH CRYOGENICALLY TREATED ...
580 Materiali in tehnologije / Materials and technology 48 (2014) 4, 577–580
Figure 8: Wear image of cemented carbides
Slika 8: Posnetek obrabe karbidne trdine
Figure 7: Surface roughness (Ra) at different cutting speeds
Slika 7: Hrapavost povr{ine (Ra) pri razli~nih hitrostih rezanja
C. C. KUO: A COST-EFFECTIVE APPROACH TO THE RAPID FABRICATION OF FUNCTIONAL METAL PROTOTYPES
A COST-EFFECTIVE APPROACH TO THE RAPID
FABRICATION OF FUNCTIONAL METAL PROTOTYPES
STRO[KOVNO U^INKOVIT NA^IN ZA HITRO IZDELAVO
FUNKCIONALNIH KOVINSKIH PROTOTIPOV
Chil-Chyuan Kuo
Department of Mechanical Engineering, Ming Chi University of Technology, No. 84, Gungjuan Road, Taishan Dist. New Taipei City 24301,
Taiwan
jacksonk@mail.mcut.edu.tw
Prejem rokopisa – received: 2013-09-26; sprejem za objavo – accepted for publication: 2013-10-10
A quick response to the market is considered as one of the important factors to ensure a company’s competitiveness. New
products must be swiftly and cost-effectively developed, manufactured and introduced to the market. This work demonstrates a
technology for the rapid manufacturing of the metal prototypes for a laptop hinge using both rapid-prototyping and
rapid-tooling techniques. A small-batch production for the laptop hinge of about 20 pieces can be obtained economically via a
silicone rubber mold using vacuum casting, without applying any plastic injection molding. The fabricated metal prototypes
have excellent physical and mechanical properties, which can be used for functional tests. A cost reduction of 75 % and a time
saving of 82 % can be achieved. This technology has considerable economic benefits and prospects for broad market
applications.
Keywords: laptop hinge, metal prototype, silicone rubber mold, epoxy resin, surface roughness
Hiter odziv na zahteve trga je eden od klju~nih dejavnikov za konkuren~nost podjetja. Novi proizvodi morajo biti hitro in
stro{kovno u~inkovito razviti, izdelani in predstavljeni trgu. To delo predstavlja tehnologijo hitre izdelave kovinskih prototipov
te~aja ohi{ja prenosnega ra~unalnika s tehniko hitre izdelave prototipa in hitre izdelave orodja. Majhna proizvodna serija, okrog
20 kosov te~aja za prenosni ra~unalnik, se lahko izdela s formo iz silikonske gume in z ulivanjem v vakuumu brez uporabe
brizganja plastike. Izdelani prototipi iz kovine imajo odli~ne fizikalne in mehanske lastnosti ter so uporabni za funkcionalno
preizku{anje. Dose`e se lahko 75-odstotno zmanj{anje stro{kov in 82-odstoten prihranek ~asa. Ta tehnologija ima velike eko-
nomske prihranke in mo`nosti {iroke uporabe na trgih.
Klju~ne besede: te~aj ohi{ja osebnega ra~unalnika, kovinski prototip, forma iz silikonske gume, epoksi smola, hrapavost
povr{ine
1 INTRODUCTION
New market realities demand faster product develop-
ment and a reduced time to market. They also require
higher quality, cost reduction and greater efficiencies. To
reduce the product development time and reduce the cost
of manufacturing, rapid prototyping (RP)1–3 has been
developed, which offers the potential to completely
revolutionize the process of manufacture. However, the
features of the prototype do not usually meet the needs
of the end product with the required material. Rapid-
tooling (RT) technologies have been developed.4–7 RT is
regarded as a natural extension of RP, since it is a tech-
nology that uses RP technologies and applies them to the
manufacturing of tool inserts. Since the importance of
RT goes far beyond component performance testing, RT
is regarded as an important method for reducing the cost
and time to market in the development of a new product.
Several RT technologies are commonly available in indu-
stry now. For the purposes of classification, RT is divi-
ded into soft and hard tooling as well as indirect and
direct tooling. Soft tooling is easier to work with than
tooling steels, because these tools are created from mate-
rials such as epoxy-based composites with aluminum
particles, silicone rubber or low-melting-point alloys. It
is well known that RT is capable of replacing conven-
tional steel tooling, saving both costs and time in the
manufacturing process. Indirect soft tooling is used more
frequently in the development of new products, rather
than direct tooling, because it is fast, simple and cost-
effective.
It is a well-known fact that the silicone rubber mold
is employed frequently because it has flexible and elastic
characteristics, so that parts with a sophisticated geome-
try can be fabricated.8 In addition, epoxy tools are fre-
quently employed for intermediate tooling, because this
material has acceptable mechanical properties and a
high-temperature resistance.9 Although some technolo-
gies, such as selective laser melting,10 direct shell pro-
duction casting,11 direct metal laser sintering,12 laser-
engineered net shaping,13 shape-deposition manufactur-
ing,14 metal injection molding15 and high-speed machin-
ing16 can be used to produce metal prototypes, the cost of
the hardware is expensive. Thus, developing a low-cost
method to manufacture a metal prototype is required. In
this work, a cost-effective method for fabricating a func-
tional metal prototype using RP and RT technologies is
demonstrated to address this important issue. The surface
roughnesses of the metal prototypes were measured
using a white-light interferometry (WLI) technique. The
Materiali in tehnologije / Materials and technology 48 (2014) 4, 581–585 581
UDK 621.7 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(4)581(2014)
properties of the fabricated metal prototypes were also
discussed, as well as the advantages of this method.
2 EXPERIMENT
Figure 1 shows a schematic illustration of the pro-
cess flow for this work. The master pattern model for
this work is a laptop hinge, which was designed using
Pro/ENGINEER software. The master pattern model was
fabricated using an RP system (Z Printer 310; Z Corpo-
ration, Inc.). A silicone rubber mold was fabricated from
the master pattern model. Figure 2 shows a schematic
illustration of the fabrication process for a silicone
rubber mold in detail. The base compound (KE-1310ST;
ShinEtsu, Inc.) and the hardener (CAT-1310S; ShinEtsu,
Inc) were mixed thoroughly to fabricate the silicone rub-
ber mold. The amounts of base compound and hardener
were calculated by multiplying the desired volume of the
silicon rubber mold to be made by the density of silicone
rubber (1.07 g/cm3 at 23 °C). In general, the curing agent
and the silicone rubber in a mass ratio of 10 : 1 were
mixed thoroughly with a stirrer. The properties of the
silicone rubber mold, such as durability and mold life,
are significantly affected by the relative amounts of
curing agent and silicone rubber, because it will affect
the cross-linking of the silicone rubber.17,18 Thus, calcu-
lating the weight of the base and the curing agent
precisely is crucial before mixing. To reduce the level of
human error, a user-friendly, man-machine interface was
developed using the Visual Basic program. A vacuum
machine was used for generating a vacuum atmosphere,
which provides the function of the degassing process to
remove the air bubbles derived from the mixing process
of the silicone rubber and the hardener. Depending on
the extent of the air bubbles in the mixture, the degassing
process can range from 25 min to 60 min. After this
degassing process, the pressure inside the vacuum ma-
chine was changed by breaking the vacuum atmosphere.
Thus, a silicone rubber mold can be fabricated without
the defects caused by the air bubbles derived from the
mixing process.
It is a well-known fact that epoxy-resin composite
materials are frequently used for making rapid tooling.
Epoxy resin (C1; Devcon, Inc.) was used to fabricate
metal prototypes of the laptop hinge. The epoxy resin
used is usually aluminum filled to enhance the mecha-
nical or physical properties. The amount of epoxy resin
required was calculated by the size of the master pattern.
The hardener and base compound (70 % aluminum
powder and 30 % liquid epoxy resin) in a weight ratio of
1 : 112 were mixed thoroughly with a stirrer. The mixing
process of aluminum powder and epoxy resin is very
important because a uniform distribution of the alumi-
num particles in the epoxy resin has a great effect on the
mechanical property of the RT. Depending on the extent
of the air bubbles in the mixture, the degas process can
range from 30 min to 80 min. A stepwise post-curing
cycle for the fabricated metal prototypes was performed
in a convection oven (DH400; Deng Yag, Inc.) to ensure
the completion of the curing reaction of the composite
materials. Figure 3 shows the temperature profiles used
C. C. KUO: A COST-EFFECTIVE APPROACH TO THE RAPID FABRICATION OF FUNCTIONAL METAL PROTOTYPES
582 Materiali in tehnologije / Materials and technology 48 (2014) 4, 581–585
Figure 2: Schematic illustration of the fabrication process for a sili-
cone rubber mold using vacuum casting
Slika 2: Shematski prikaz postopka izdelave forme iz silikonske gume
z ulivanjem v vakuumu
Figure 1: Schematic illustration of the process flow for this work
Slika 1: Shematski prikaz poteka postopka
for post-curing the metal prototype of the laptop hinge.
This temperature profile consists of three phases. In the
first phase the green part is heated to 55 °C and remains
at a constant temperature for 2 h to reach a suitable tem-
perature for the cross-linking of epoxy resins. In the
second phase the temperature was increased from 55 °C
to 121 °C, and then remains at a constant temperature for
2 h to enhance the mechanical properties of the metal
prototypes. In the third phase the temperature was
increased from 121 °C to 204 °C and remains constant
for 2 h to cure the metal prototypes fully. The surface
roughnesses of the metal prototypes were measured
using a WLI (vertical resolution 0.1 nm, horizontal
resolution 0.5 μm, Chroma 7502). The sampling area
was chosen to be 50 μm × 50 μm. The arithmetic average
roughness value (Ra) was used to quantify the surface
finish. The Ra values represent the average displacement
of the peaks and valleys measured with respect to a mean
line.
3 RESULTS AND DISCUSSION
Figure 4 shows the master-pattern models of the
male and female parts of the laptop hinge. The fabricated
master pattern models required further processing,
because of their relatively poor surface roughness, which
results from the layered structure inherent in the building
method.19 In this work, hand finishing was used to
improve the stair-step surface texture inherently in the
fabricated master-pattern models.20 Figure 5 shows the
silicone rubber molds for the male and female parts of
the laptop hinge. The cured silicone rubber block was cut
along the parting line and the master pattern was re-
moved. It is clear that the surfaces of the silicone rubber
molds have the desired surface finish. In addition, the
silicone rubber mold enables the master-pattern models
to release without breakage or damage.21,22
To evaluate the validity of the fabricated silicone
rubber mold, some wax patterns were replicated from the
elastomeric mold. Figure 6 shows that the wax patterns
of the laptop hinge were replicated from the silicone
rubber mold because of the low interfacial free energy
and the chemically inert surface.23 It is clear that the
laptop hinges made of wax were successfully transferred
from the silicone rubber mold by vacuum casting. The
detailed appearance of the laptop hinges was faithfully
replicated. This means that the silicone rubber mold
fabricated can be used to replicate metal prototypes of
the designed laptop hinges. The use of epoxy resin is
often the fastest way to complete short runs of functional
prototypes manufactured by vacuum casting. Figure 7
shows the green part of a male laptop hinge fabricated
C. C. KUO: A COST-EFFECTIVE APPROACH TO THE RAPID FABRICATION OF FUNCTIONAL METAL PROTOTYPES
Materiali in tehnologije / Materials and technology 48 (2014) 4, 581–585 583
Figure 5: Silicone rubber molds for: a) male part and b) female part of
a laptop hinge
Slika 5: Forma iz silikonske gume za: a) mo{ki del in b) `enski del
te~aja za ohi{je prenosnega ra~unalnika
Figure 4: Master-pattern models of: a) male and b) female parts of a
laptop hinge
Slika 4: Modeli: a) mo{kega in b) `enskega dela te~aja za ohi{je
prenosnega ra~unalnika
Figure 3: Temperature profiles used for post-curing the metal proto-
type of a laptop hinge
Slika 3: Temperaturni profil, uporabljen za utrjevanje po strjevanju
kovinskega prototipa le`aja ohi{ja prenosnega ra~unalnika
using a silicone rubber mold. It can be observed that the
metal prototype of the male laptop hinge was fabricated
faithfully. The life of a silicone rubber mold depends to a
large extent on the surface finish of the fabricated laptop
hinges. In general, a silicone rubber mold will reproduce
up to 20–30 parts with a gradual deterioration of the
surface quality under normal operating conditions. After
the post-curing process, metal prototypes of the laptop
hinge were shown in Figure 8. Thus, a small-batch pro-
duction of the metal prototype can be obtained econo-
mically via a silicone rubber mold using vacuum casting
without applying any plastic injection molding. Accord-
ing to the specification provided by the epoxy-resin
composite supplier, the fabricated metal prototypes have
excellent physical and mechanical properties (tensile
strength ASTM D 638, compressive strength ASTM D
695, and cured hardness ASTM D 2240). A tensile
strength of 35439 kPa (5140 psi) can be achieved, which
can be used for a functional test. Figure 9 shows the
WLI image of a metal prototype. The average roughness
of the metal prototype is approximately 114 nm. This
result shows that the fabricated metal prototype has the
desired surface finish. As described above, this method is
a viable alternative for the fabrication of metal proto-
types.24 Table 1 shows the time needed to manufacture
metal prototypes of a laptop hinge. Table 2 shows the
cost needed to manufacture metal prototypes of a laptop
hinge. The estimated cost and time for manufacturing the
metal prototype of a laptop hinge was NT$ 683 and 19.5
h, which represents a cost reduction of 75 % and a time
saving of 82 % compared to the traditional manufac-
turing method (NT$ 2730 and 108 h)25 using a metal
injection molding process.26–28 This method provides a
C. C. KUO: A COST-EFFECTIVE APPROACH TO THE RAPID FABRICATION OF FUNCTIONAL METAL PROTOTYPES
584 Materiali in tehnologije / Materials and technology 48 (2014) 4, 581–585
Figure 7: Green part of a male laptop hinge fabricated using a silicone
rubber mold
Slika 7: Surovec mo{kega dela te~aja za ohi{je prenosnega ra~unal-
nika, izdelan v formi iz silikonske gume
Figure 6: Wax patterns of the laptop hinge were replicated from the
silicone rubber mold
Slika 6: Vo{~eni modeli te~aja za ohi{je prenosnega ra~unalnika,
izdelani z uporabo forme iz silikonske gume
Figure 9: WLI image of a metal prototype
Slika 9: WLI-posnetek kovinskega prototipa
Figure 8: Metal prototypes of the laptop hinge after the post-curing
process
Slika 8: Kovinski prototipi te~aja za ohi{je prenosnega ra~unalnika po
postoku utrjevanja
Table 1: The time needed to manufacture the metal prototype of a
laptop hinge
Tabela 1: ^as izdelave kovinskega prototipa te~aja za ohi{je
prenosnega ra~unalnika
Manufacturing
process
Rapid
prototyping of
a laptop hinge
Silicone
rubber
mold
Metal
prototype of a
laptop hinge
Time (h) 0.5 8 11
Table 2: The cost needed to manufacture the metal prototype of a
laptop hinge
Tabela 2: Stro{ki izdelave kovinskega prototipa te~aja za ohi{je
prenosnega ra~unalnika
Manufacturing
process
Rapid
prototyping of
a laptop hinge
Silicone
rubber
mold
Metal
prototype of a
laptop hinge
NT$ 141 444 98
low-cost method to fabricate the functional metal
prototype with the desired mechanical properties for the
functional test. A major contribution of this work is that
this method is a simple and cost-effective method for
fabricating metal prototypes of a laptop hinge with high
yield.
4 CONCLUSIONS
Manufacturing industry is currently under increasing
pressure due to international competition. New products
must be developed more quickly and cheaply, and then
manufactured and introduced to the market. A procedure
for fabricating metal prototypes with the desired surface
finish and accuracy has been demonstrated in this work.
The research results reported in this study have industrial
application values because this is a simple and effective
method in terms of cost and time reductions. This pro-
cess has considerable economic benefits, which can be
employed to fabricate metal prototypes for a functional
test in the first development phase of a new product. In
addition, this process has the potential to be extended to
other fields of applications with the development of
casting materials, such as ceramics, low-melting-point
alloys, or some functional polymers.
Acknowledgement
The authors gratefully acknowledge the financial
support of the National Science Council of Taiwan under
contracts nos. NSC 102-2221-E-131-012 and NSC
101-2221-E-131-007. The skillful technical assistance of
Mr. Zhen-Jia Xu of Ming Chi University of Technology
is highly appreciated.
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M. OPIELA: THERMOMECHANICAL TREATMENT OF Ti-Nb-V-B MICRO-ALLOYED STEEL FORGINGS
THERMOMECHANICAL TREATMENT OF Ti-Nb-V-B
MICRO-ALLOYED STEEL FORGINGS
TERMOMEHANSKA OBDELAVA ODKOVKOV IZ
MIKROLEGIRANEGA JEKLA Ti-Nb-V-B
Marek Opiela
Silesian University of Technology, Institute of Engineering Materials and Biomaterials, 18a Konarskiego Street, 44-100 Gliwice, Poland
marek.opiela@polsl.pl
Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2013-10-09
The work presents the research results of the influence of thermomechanical treatment through forging on the microstructure
and mechanical properties of newly elaborated micro-alloyed steel containing 0.28 % C, 1.41 % Mn, 0.028 % Ti, 0.027 % Nb
and 0.019 % V. The applied thermomechanical treatment makes it possible to obtain a fine-grained microstructure of austenite
during the hot working and the production of forged parts that acquire an advantageous set of mechanical properties and a
guaranteed crack resistance after controlled cooling from the finishing plastic-deformation temperature and the successive
tempering. The forgings produced with the method of thermomechanical treatment, consecutively subjected to tempering in the
temperature range from 550 °C to 650 °C, have YS0.2 from 994 MPa to 939 MPa, UTS from 1084 MPa to 993 MPa and KV–40
from 77 J to 83 J.
Keywords: micro-alloyed steels, thermomechanical treatment, forged parts, mechanical properties
Delo predstavlja rezultate vpliva termomehanske obdelave s kovanjem na mikrostrukturo in mehanske lastnosti novo izdelanega
mikrolegiranega jekla z 0,28 % C, 1,41 % Mn, 0,028 % Ti, 0,027 % Nb in 0,019 % V. Uporabljena termomehanska obdelava
omogo~a doseganje drobnozrnate mikrostrukture avstenita med vro~o predelavo in izdelavo odkovkov, ki dobijo dobre mehan-
ske lastnosti in zagotovljeno odpornost proti razpokam po kontroliranem ohlajanju iz kon~ne temperature plasti~ne deformacije
in po popu{~anju. Odkovki, izdelani po metodi termomehanske obdelave in nato popu{~ani v temperaturnem obmo~ju od
550 °C do 650 °C, imajo YS0.2 od 994 MPa do 939 MPa, UTS od 1084 MPa do 993 MPa in KV–40 od 77 J do 83 J.
Klju~ne besede: mikrolegirana jekla, termomehanska obdelava, odkovki, mehanske lastnosti
1 INTRODUCTION
A lowering of the production cost is the basic reason
for implementing economical technologies for the products
made of constructional alloy steels with the methods of
thermomechanical treatment. These methods consist of
plastic working in the conditions adjusted to the steel
chemical composition and micro-alloying with the con-
secutive direct cooling of the parts from the temperature
of the plastic-deformation finish or after a particular time
specified. This allows a reduction of the expensive heat
treatment of products involving the tempering.1–4
The HSLA-type (high-strength low-alloy) micro-
alloyed steels – containing up to 0.3 % C and 2 % Mn,
micro-additions with a high chemical affinity to N and C,
i.e., Nb, Ti and V in the amount of about 0.1 %, and
sometimes also a slightly increased concentration of N
and up to 0.005 % of B, increasing hardenability – are
particularly useful for the production of forged parts with
a fine-grained microstructure using the method of ther-
mo-mechanical processing. The interaction of micro-
additions in solid steel depends on their state influenced
by the conditions of the performed plastic working.
Micro-additions in the solid solution raise the tempe-
rature of the recrystallization of plastically deformed
austenite. Moreover, the segregation of micro-additions
at the grain boundaries decreases their mobility. Instead,
the micro-additions precipitating on dislocations in the
form of dispersive particles of MX interstitial phases,
slow down the course of dynamic recovery and possibly
also dynamic recrystallization during the plastic defor-
mation and additionally decrease the rate of thermal
recovery and static or metadynamic recrystallization,
limiting the grain growth of the recrystallized austenite
in the intervals between the successive stages of defor-
mation and after its completion. Dispersive particles of
MX-type phases hamper the grain-boundary movement
of the recrystallized austenite and, hence, it is possible to
produce fine-grained-microstructure products addition-
ally strengthened by carbonitrides.5–11 The contribution
of micro-additions to the precipitation strengthening can
be calculated using the Ashby-Orowan relationship:



p =
⎛
⎝
⎜ ⎞
⎠
⎟0538
2
1 2.
ln
/Gbf
b
(1)
where G is the shear modulus, b is the Burgers vector, f
is the particle-volume fraction and  is the mean diame-
ter of particles.
When engineering the calculations, there is also
another way of assessing the effect of the strengthening
from the micro-alloyed particles:12
p = Kp[M] (2)
Materiali in tehnologije / Materials and technology 48 (2014) 4, 587–591 587
UDK 669.14.018.298:620.17 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(4)587(2014)
where [M] is the micro-addition content in the mass
fraction and Kp is the constant.
Niobium is often used as a micro-addition in the
steels subjected to various processes of thermomecha-
nical treatment. Very dispersive Nb(C,N) precipitates
have a very important role in this case. Therefore, the
value of p is often calculated with the assumption that
the precipitation-strengthening effect depends mainly on
this element. It should be noted that the effectiveness of
Nb in the precipitation hardening is reduced by a Ti
micro-addition forming both its carbides and nitrides as
well as complex precipitates of (Ti,Nb)C and (Ti,Nb)
(C,N).13
The state of niobium (in a solution or as a precipi-
tate), determined by the reheating temperature, can affect
the recrystallization, the grain growth and the  trans-
formation of austenite.14,15 For example, the recrystal-
lization and grain growth of austenite are significantly
suppressed by the precipitation of NbC prior to the 
transformation. In addition, coarse NbC particles can be
the preferred sites for the ferrite nucleation. In particular,
the control of the austenite recovery and recrystallization
is an important part of the grain-refinement technique in
the modern thermo-mechanical controlling process.16 An
addition of Nb to steel is considered to have three
primary effects:17,18 (i) being an inhibitor of the austenite
grain coarsening during reheating, (ii) suppressing the
austenite recrystallization prior to the  transforma-
tion through the strain-induced precipitation of NbC and
(iii) causing the precipitation hardening from NbC in the
low-temperature transformation step of the thermo-me-
chanical process. The strongest contribution to the
strengthening is the refinement of the final microstruc-
ture (essentially ferrite grain size), which accounts for
80–90 % of y. The key role of the precipitates is to
provide the dispersion strengthening that is often
generated in micro-alloyed steels by the NbC, VC,
Nb(C,N) or V(C,N) particles, depending on whether N is
added or not, with less than 20 nm in size.19 Special
attention also has to be paid to the vanadium micro-
alloying addition. This element is easily added to liquid
steel and its solubility during reheating is very high. The
strengthening effect is enhanced as nitrogen is also
added to the solution.
The final microstructure and mechanical properties
strongly depend on the chemical composition, the con-
trolled hot-working parameters and the cooling condi-
tions of the steel. High strength, good ductility and good
formability are developed in steel products during the
manufacturing processes and, to achieve these goals,
properly balanced quantities of micro-alloying additions
and suitable thermomechanical processing schedules
have to be carefully applied.20,21
2 EXPERIMENTAL PROCEDURE
The study was performed on newly elaborated mi-
cro-alloyed steel, intended for the production of forged
machine parts with a high strength, using thermomecha-
nical treatment. The investigated steel contains 0.28 %
C, 1.41 % Mn, 0.29 % Si, 0.008 % P, 0.004 % S, 0.26 %
Cr, 0.11 % Ni, 0.22 % Mo, 0.025 % Al and Nb, Ti, V
and B in the amounts of (0.027, 0.028, 0.019 and 0.003) %,
respectively.
The investigated steel with the weight of 100 kg,
molten in a VSG-100 type laboratory vacuum induction
furnace and cast in an atmosphere of argon into quadratic
hot-topped ingots with the following dimensions: top –
160 per bottom – 140 mm × 640 mm. The ingots were
forged into 32 mm × 160 mm flat bars, with open-die
forging in a high-speed hydraulic press, using a force of
300 MN.
The conditions of the hot processing and cooling,
leading to the desired mechanical properties of the forg-
ings, were selected with the following procedures:22
• analyses of the kinetics of the MX-type interstitial-
phase precipitation in the solid state;
• an investigation of the process of hot working of steel
with the method of continuous compression of the
specimens at the rate of (1, 10 and 50) s–1 in a tem-
perature range from 1100 °C to 900 °C;
• an examination of the kinetics of strain hardening
(recrystallization) and softening of plastically de-
formed austenite in the mentioned conditions,
• an investigation of the kinetics of the phase transfor-
mations of undercooled austenite.
The obtained results were used to develop two
variants of forging with thermomechanical treatment of
160 mm × 32 mm flat bars into 14 mm thick flat bars, in
the temperature range of 1100 °C to 900 °C at the strain
rate of 3 s–1, applying 50 % of draft. The soaking tem-
perature for forging was 1150 °C and the time was 45
min. In the first variant, segments of flat bars were
hardened in water directly from the temperature of the
forging finish, while, in the second variant, flat bars were
isothermally held at the temperature of 900 °C for (10,
60 and 100) s after the forging finish and prior to harden-
ing in water. Directly after quenching, the obtained flat-
bar sections were tempered for 1 h at the temperatures of
550 °C and 650 °C.
For the metallographic investigations of the speci-
mens hardened after plastic deformation in the men-
tioned conditions and after high-temperature tempering,
a Leica MEF 4A light microscope was used. The thin
foils were examined with a TITAN80-300 FEI ultra-high
resolution transmission electron microscope at the
accelerating voltage of 300 kV.
Static tensile tests were performed with an INSTRON
1115 universal testing machine on the samples with a
diameter of 8 mm and a gauge length of 40 mm. The
impact testing at room temperature and at –40 °C was
carried out on a Charpy pendulum machine with the
initial energy of 300 J, using V-notch specimens with a
cross-section of 8 mm × 10 mm. Brinell hardness was
also determined for all the specimens.
M. OPIELA: THERMOMECHANICAL TREATMENT OF Ti-Nb-V-B MICRO-ALLOYED STEEL FORGINGS
588 Materiali in tehnologije / Materials and technology 48 (2014) 4, 587–591
3 RESULTS
Metallographic examinations of the samples of
Variant I of thermomechanical treatment, i.e., the ones
quenched in water immediately after plastic deformation,
revealed the microstructure of fine-lath martensite
(Figure 1). The segments produced according to Variant
I of thermomechanical treatment, subjected successively
to tempering at the temperature of 650 °C, have the
microstructure of the tempered martensite with the preci-
pitates of carbide particles (Figure 2).
The analyses of the microstructure of thin foils
revealed that the microstructure of steel, hardened
directly from the forging temperature, consists of lath
martensite and an uneven distribution of dislocation
density in strongly plastically deformed austenite. Fine
(Ti,Nb)C particles, located mainly at the boundaries of
martensite laths (Figure 3) were observed in the marten-
site with a diversified spatial orientation of individual
laths. Particles of similar chemical composition, morpho-
logy and size were also identified in the C-Mn-0.16Nb
steel.23 Austenite in the form of thin films between the
laths of martensite was observed. The Kurdjumov-Sachs
crystallographic relationship between the retained auste-
nite and martensite was confirmed.
A similar morphology of lath martensite was found
for the steel coming from the flat bar section obtained
with Variant II of thermomechanical forging. Laths of
different widths occur in the martensite packets. Some of
the martensite laths show a fragmentation caused by
high-angle subgrain boundaries formed into the statically
recovered austenite that hampers their growth. As in the
previous case, the (Ti,Nb)C and (Ti,Nb,V)C (Figure 4)
particles with the size from 40 nm to 80 nm were found
in the martensite.
The forgings produced in both variants of thermo-
mechanical processing and the forgings quenched con-
ventionally from the temperature of 900 °C have quite
different microstructures after the tempering in the tem-
perature range from 550 °C to 650 °C. The microstruc-
ture of the specimens taken from flat bar sections,
M. OPIELA: THERMOMECHANICAL TREATMENT OF Ti-Nb-V-B MICRO-ALLOYED STEEL FORGINGS
Materiali in tehnologije / Materials and technology 48 (2014) 4, 587–591 589
Figure 3: a) Carbide (Ti,Nb)C in the martensite, b) EDS spectrum; Variant I of thermomechanical treatment: 900 °C, water
Slika 3: a) Karbid (Ti,Nb)C v martenzitu, b) EDS-spekter; Varianta I termomehanske obdelave: 900 °C, voda
Figure 1: Martensitic microstructure; Variant I of thermomechanical
treatment: 900 °C, water
Slika 1: Martenzitna mikrostruktura; Varianta I termomehanske
obdelave: 900 °C, voda
Figure 2: High-tempered martensite; Variant I of thermomechanical
treatment: 900 °C, water; tempering temperature: 650 °C
Slika 2: Visoko popu{~eni martenzit; Varianta I termomehanske
obdelave: 900 °C, voda; temperatura popu{~anja: 650 °C
quenched directly after forging, after the tempering at the
temperature of 550 °C, consists of the tempered marten-
site with a precipitation of granular and lamellar Fe3C
particles, distributed inside the grains and at the lath
boundaries. The lamellar and granular precipitates,
formed in such conditions, fulfill the matrix spatial
dependences established by Bagariacki. An increase in
the tempering temperature leads to a coagulation of
cementite. The M3C lamellar precipitates retain the privi-
leged spatial orientation with ferrite, while the coagu-
lated granular particles of this phase reveal random
crystallographic orientation with respect to the matrix.
The samples obtained with the Variant-II forging of
thermomechanical treatment have a similar microstruc-
ture in the tempered state. In this case, in the microstruc-
ture of the steel tempered at 550 °C there are thin
lamellar precipitations of Fe3C at the lath boundaries,
while inside the laths of the recovered ferrite, dispersive
lamellar particles of the phase with the privileged spatial
orientation with the matrix are observed. Lamellar pre-
cipitates transform into granular Fe3C particles at the lath
boundaries and subgrain boundaries of the recovered
ferrite with an increase in the tempering temperature.
Moreover, the precipitates of (Ti,Nb)C with a morpho-
logy similar to that after thermomechanical processing
were found after tempering.
The steel microstructure in both hardened and tem-
pered state significantly affects the mechanical proper-
ties of the flat bar sections obtained with both thermo-
mechanical treatments. The results of the investigation of
the mechanical properties and the impact energy of the
Charpy-V samples, taken from the forgings are listed in
Table 1. The data presented in this table show that the
flat bar sections quenched in water directly from the
temperature of the forging finish (Variant I of thermo-
mechanical treatment), have the following properties
after the tempering at 550 °C: YS0.2 of about 973 MPa,
UTS of about 1057 MPa, TEl of about 13.5 % and RA of
about 51.5 %. The flat bar sections obtained from
Variant II of thermomechanical processing have higher
mechanical properties and a distinctly higher crack
resistance in the tempered state. However, the best set of
mechanical properties and crack resistance was found for
the forging isothermally held at the temperature of
M. OPIELA: THERMOMECHANICAL TREATMENT OF Ti-Nb-V-B MICRO-ALLOYED STEEL FORGINGS
590 Materiali in tehnologije / Materials and technology 48 (2014) 4, 587–591
Figure 4: a) Carbide (Ti,Nb,V)C in the martensite, b) EDS spectrum; Variant II of thermomechanical treatment: 900 °C, 60 s, water
Slika 4: a) Karbid (Ti,Nb,V)C v martenzitu, b) EDS-spekter; Varianta II termomehanske obdelave: 900 °C, 60 s, voda
Table 1: Results for the mechanical properties and impact-fracture energy of Charpy-V samples after thermomechanical forging and successive
tempering
Tabela 1: Mehanske lastnosti in udarna `ilavost Charpy-V vzorcev po termomehanskem kovanju in toplotni obdelavi
Variant
Treatment conditions
Tempering
temperature
°C
Mechanical properties Impact energy
Charge
heating/finish
forging tem-
perature, °C
Isothermal
holding
time, s
Cooling
medium
YS0.2
MPa
UTS
MPa
TEl
%
RA
%
KV
J
KV–40
J
I 1150/900
– water 550 973 1057 13.5 51.5 69.3 55.0
– water 650 909 976 14.0 52.0 81.7 68.6
II 1150/900
10 water 550 967 1063 13.6 49.6 91.6 71.3
10 water 650 892 958 14.5 52.7 99.0 79.7
60 water 550 994 1084 14.3 50.9 95.0 77.3
60 water 650 911 974 15.1 50.2 108.7 82.6
100 water 550 988 1077 13.0 49.6 95.7 75.7
100 water 650 939 993 14.8 51.3 101.3 80.0
900 °C for 60 s prior to hardening in water when, after
the tempering in the temperature range from 550 °C to
650 °C, the following properties were obtained: YS0.2
from 994 MPa to 911 MPa, UTS from 1084 MPa to 974
MPa, KV from 95 J to 109 J and KV–40 from 77 J to 83 J.
The examination of the influence of the applied
processing variant and the tempering temperature on the
hardness show that the highest hardness – of approxima-
tely 330 HBW – is obtained for the forgings of the
Variant-II thermomechanical treatment, including iso-
thermal holding at the forging-finish temperature for 60 s
prior to water quenching, and tempering at 550 °C. An
increase in the tempering temperature for a forging up to
650 °C results in a mild decrease in the hardness, to
about 300 HBW.
4 CONCLUSIONS
The implemented thermomechanical processing
allows a production of forged products with the follow-
ing properties achieved after controlled cooling from the
temperature of the plastic deformation finish and a sub-
sequent tempering in the temperature range from 550 °C
to 650 °C: YS0.2 from 994 MPa to 939 MPa, UTS from
1084 MPa to 993 MPa, TEl from 14.3 % to 15.1 %, RA
from 51.5 % to 52.7 %, KV from 96 J to 109 J and KV–40
from 77 J to 83 J. The high strength in such a state at a
high crack resistance, also at a decreased temperature, is
noteworthy.
The best set of mechanical properties and crack
resistance was obtained for the forging isothermally held
at 900 °C for 60 s prior to quenching in water, and sub-
sequently subjected to a tempering at 650 °C.
The relatively low hardness of the steel after a
high-temperature tempering should not cause difficulties
during the machining of forgings.
Microstructural observations of thin foils using trans-
mission electron microscopy revealed that the micro-
structure of the steel in the quenched state consisted of
lath martensite with a high density of dislocations and a
considerable amount of precipitates. In the examined
steel, the particles of (Ti,Nb)C and (Ti,Nb,V)C revealed
precipitation on dislocations during plastic deformation,
decreasing the rate of dynamic recovery, possibly dyna-
mic recrystallization and, after hot working, also
decreasing the rate of recovery and static or metadyna-
mic recrystallization. The identified particles with the
sizes between 40 nm and 80 nm significantly increase
the precipitation strengthening; however, they limit the
grain growth of the recrystallized austenite and stimulate
a formation of fine-grained microstructures.
The investigation makes it possible to develop an
industrial technology for the forgings with high mecha-
nical properties and a guaranteed crack resistance, also at
a decreased temperature, using thermomechanical treat-
ment.
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M. OPIELA: THERMOMECHANICAL TREATMENT OF Ti-Nb-V-B MICRO-ALLOYED STEEL FORGINGS
Materiali in tehnologije / Materials and technology 48 (2014) 4, 587–591 591

T. MANDYS et al.: PROGRESSIVE FAILURE ANALYSIS OF COMPOSITE SANDWICH BEAM ...
PROGRESSIVE FAILURE ANALYSIS OF COMPOSITE
SANDWICH BEAM IN CASE OF QUASISTATIC
LOADING
NAPREDNA ANALIZA PO[KODB SESTAVLJENEGA
KOMPOZITNEGA NOSILCA PRI KVAZISTATI^NEM
OBREMENJEVANJU
Tomá{ Mandys, Tomá{ Kroupa, Vladislav La{
University of West Bohemia in Pilsen, NTIS - New Technologies for Information Society, Univerzitní 22, 306 14 Plzeò, Czech Republic
tmandys@kme.zcu.cz, kroupa@kme.zcu.cz, las@kme.zcu.cz
Prejem rokopisa – received: 2013-10-02; sprejem za objavo – accepted for publication: 2013-10-28
This paper is focused on a progressive failure analysis of a sandwich composite panel considering a non-linear model of core
and skins using explicit analysis. The material properties were determined from tensile and compressive tests of the outer
composite skin and the foam core. A user-defined material model was used to describe the non-linear orthotropic elastic
behaviour of the composite skin. The material parameters of the outer composite skin were determined using the identification
process performed using a combination of optimization method and finite-element simulation. The Low-Density Foam material
model was used for the foam core. The obtained material data were validated using a three-point bending test of a sandwich
beam.
Keywords: sandwich panel, fiber-glass fabric, foam core, three point bending test, damage, optimization
^lanek obravnava napredno analizo po{kodb z uporabo eksplicitne analize sestavljene kompozitne plo{~e z upo{tevanjem
nelinearnega modela jedra in skorje. Lastnosti materiala so bile dolo~ene iz nateznih in tla~nih preizkusov zunanje kompozitne
skorje in penaste sredice. Uporabni{ko dolo~en model materiala je bil uporabljen za opis nelinearnega ortotropnega elasti~nega
vedenja kompozitne skorje. Materialni parametri zunanje kompozitne skorje so bili dolo~eni z identifikacijskim procesom,
izvr{enim z uporabo kombinacije metod optimiranja in simulacije s kon~nimi elementi. Materialni model pene z nizko gostoto
je bil uporabljen za jedro iz pene. Dobljeni podatki za material so bili ocenjeni z uporabo trito~kovnega upogibnega preizkusa
sestavljenega nosilca.
Klju~ne besede: sestavljena plo{~a, tkanina iz steklenih vlaken, jedro iz pene, trito~kovni upogibni preizkus, po{kodba,
optimizacija
1 INTRODUCTION
The principle of sandwich structures is based on a
low-density material (core) placed between two stiffer
outer skins. The main purpose of the core is to maintain
the distance between the outer skins and to transfer the
shear load while the outer skins carry the compressive
and tensile load. The outer skins are obviously thinner
than the core. This type of structural arrangement has a
much larger bending stiffness than a single solid plate
made of the same total weight from the same material as
a outer skin only. This fact together with other advan-
tages, such as corrosion resistance, product variability
and thermal or acoustic insulation, make sandwich struc-
tures the preferred alternative to conventional materials
in all types of structural applications where the weight
must be kept to a minimum value.
2 EXPERIMENTAL SET-UPS AND SPECIMENS
The used sandwich composite panel was manufac-
tured using a vacuum infusion process and its overall
thickness was 12.5 mm. The outer composite skins were
made from three layers of fibre-glass fabric with the pro-
duct name Aeroglass (material density  = 390 g/m2) and
epoxy resin designated as Epicote HGS LR 285. The
thicknesses of the outer composite skins were 1.2 mm.
The core of the sandwich panel was a closed-cell, cross-
linked polymer foam Airex C70.55. This foam core is
characterized by a low resin absorption. The resultant
sandwich panel was cured for 6 h at 50 °C.
Tensile and compressive experimental tests were
performed using a Zwick/Roell Z050 testing machine on
the separated specimens of the laminated composite skin
and foam core of the resultant sandwich structure. The
foam core was additionally subjected to a three-point
bending test too.
Three types of specimens of laminated composite
skin were used, the chosen material principal direction
(weft) form angle 0° (Type A), 45° (Type C) and 90°
(Type B) with the direction of the loading force. The
specimens of the laminated outer skin had the dimen-
sions 135 mm × 15 mm and a thickness 1.2 mm. The
dimensions of the specimens of the isotropic foam core
were 150 mm × 15 mm with a thickness 10 mm. The
loading velocity during the tensile tests was 2.0 mm/min
in the case of the composite skin and 1.0 mm/min in the
Materiali in tehnologije / Materials and technology 48 (2014) 4, 593–597 593
UDK 66.017:620.17 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(4)593(2014)
case of the foam core. Four specimens for each angle of
composite skin and for the foam core were tested.
The three-point bending tests of the foam core were
performed on specimens with the same dimensions that
were used for the tensile tests. The distance of the sup-
ports of the testing device was 100 mm and the diameters
of all three supports were 10 mm.
The resulting sandwich beam with the dimensions
330 mm × 50 mm and a total thickness 12.5 mm was
subjected to a three-point bending test in order to vali-
date the material data obtained independently for the
composite skin and the foam core. The thicknesses of the
composite skins were 1.2 mm. The distance of the sup-
ports of the testing device was 250 mm and the supports
consisted of rotationally joined cylinders with diameters
of 30 mm. The loading velocity of the sandwich beam
was 20 mm/min. Three specimens were tested in total.
2.1 Material model of the skin
A user-defined material model of the composite skin
was implemented in Abaqus software using the VUMAT
subroutine written in the Fortran code. The non-linear
function describing the stress-strain relationship starting
from the deformation 0i (i = 1, 2) was assumed in the
case of the principal material directions 1 and 2 (equat-
ions (2) and (4)). The non-linear function with a constant
asymptote was considered in the case of the shear in
plane 12 1 (equation (8)). The following equations des-
cribe the stress-strain relationship of the laminated outer
skin:
   1 11 1 12 2 13 3= ⋅ + ⋅ + ⋅C C C for 1 < 01 (1)
    
  
1 11 1
1
01
2
1
2
1 01
01 1 12 2
2
= ⋅ + ⋅ − − ⋅ ⋅
− + ⋅
( ( ( )
( ))
C
A
A
C + ⋅ ⋅ −C D13 3 1 ) ( )
for 1  01 (2)
   2 12 1 22 2 23 3= ⋅ + ⋅ + ⋅C C C for 2 < 02 (3)
    
  
2 12 1 22 2
1
02
2
2
2
2 02 02 2
2
= ⋅ + ⋅ + ⋅ − −
⋅ ⋅ −
( ( ( )
( ))
C C
A
A + ⋅ ⋅ −C D23 3 1 ) ( )
for 2  02 (4)
   3 13 1 23 2 33 3 1= ⋅ + ⋅ + ⋅ ⋅ −( ) ( )C C C D (5)
 23 23 23 1= ⋅ ⋅ −( ) ( )G D (6)
 13 13 13 1= ⋅ ⋅ −( ) ( )G D (7)





12
12
0
12
12
0
12
12
0
1
1
12 12
=
⋅
+
⋅⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟
⎛
⎝
⎜
⎜
⎞
⎠
⎟
⎟
⎛
G
G
n n⎝
⎜⎜
⎞
⎠
⎟⎟
⋅ −( )1 D (8)
The constants occurring in equations (1) to (5) are:
C
E v v
C
E v v v
C
E v
11
1 23 32
12
1 21 23 32
13
1 3
1
=
⋅ − ⋅
=
⋅ + ⋅
=
⋅
( ) ( )
(
Δ Δ
1 32 21
22
2 31 13
23
2 32 31 12
1− ⋅
=
⋅ − ⋅
=
⋅ − ⋅
v v
C
E v v
C
E v v v
) ( )
( )
Δ Δ
Δ Δ
Δ =1−
C
E v v
v v v v v v
33
3 12 21
12 21 23 32 12 23
1
2
=
⋅ − ⋅
⋅ − ⋅ − ⋅ ⋅ ⋅
( )
v 31
(9)
where E1, E2 and E3 are the Young’s moduli in the prin-
cipal directions 1, 2 and 3 and 12, 23, 31 are the
Poison’s ratios in the planes defined by the principal
directions 1, 2 and 3. The shear moduli in planes 23 and
13 are designated as G23 and G13, respectively. The
non-linear behavior in shear in plane 12 (8) is described
using the initial shear modulus G12
0 , the asymptotic
value of the shear stress 
12
0 and the shape parameter n12.
The parameters A1 and A2 in (2) and (4) describe the
straightening of the yarns of the fiber-glass fabric and
the loss of stiffness in the directions 1 and 2, respec-
tively. The values of the deformations 01 and 02 indi-
cate the the transitions between the linear and non-linear
parts of the stress-strain relationship in the given direc-
tions 1 and 2 during the loading.
The maximum stress failure criterion was used to
predict failure on the composite skin:
F
X
F
X
F
Y
F
Y
F
S
1
1
1
1
2
2
2
2
12
12
T
T
C
C
T
T
C
C L
= = =
= =
  
 
(10)
where the subscripts T and C denote the tension and
compression, X and Y are the strengths in the principal
directions 1 and 2, respectively, and SL denotes the shear
strength.
The values of the degradation variable D are depen-
dent on the kind of occurred failure.2 The principle of
material degradation is shown in Figure 1 and the degra-
dation parameters are summarized in (11). The degra-
dation parameter after failure initiation was implemented
in the case of shear failure (F12  1.0 and   12 12
F )
from.3 Due to the non-linear behavior of the composite
skin the degradation parameters are assumed in the form:
F D F D
F D F D
F
1 1
2 2
12
1 10 1 06
1 10 1 06
T C
T C
1
≥ ⇒ = ≥ ⇒ =
≥ ⇒ = ≥ ⇒ =
≥
. .
. .
and
1 and
12
1
 

< ⇒ = −
≥
⎛
⎝
⎜⎜
⎞
⎠
⎟⎟
12
1
12
1 12
12
12
F m
F
D e
F
m( )
2 ≥ ⇒ =12 10
F D .
(11)
In the case of shear failure the non-negative material
constant m12 is represented by the integer and 12
F is the
ultimate deformation when the material is fully
damaged.
Mathematical optimization was used to identify the
material parameters on data from the performed tensile
tests of composite skin. The optimization process was
T. MANDYS et al.: PROGRESSIVE FAILURE ANALYSIS OF COMPOSITE SANDWICH BEAM ...
594 Materiali in tehnologije / Materials and technology 48 (2014) 4, 593–597
handled using Optislang 3.2.0. The material parameters
that do not have a significant influence on the results
were kept constant during the optimization process. All
the material parameters are summarized in Table 1.
The parameters E3, G13 and G23 were taken from the
literature.4 The strengths of the composite skin were
determined directly from the experimental data.
Table 1: Material parameters of the composite skin
Tabela 1: Parametri materiala kompozitne skorje
Optimized values: Constant values:
E1 GPa 16.9 v12 – 0.337
E2 GPa 18.5 v23 – 0.337
G12
0 GPa 4.96 v31 – 0.28

12
0 MPa 39.66 G13 GPa 4.0
n12 – 0.9 G23 GPa 2.75
A1 – 10.0 E3 GPa 8.0
A2 – 14.0  12
F – 0.324
01 – 0.0008 C kg/m
3 1554
02 – 0.005 XT MPa 325
m12 – 5 YT MPa 347
XC MPa 65
YC MPa 67
SL MPa 35
2.2 Material model of the foam core
The used material model of the foam core was the
Low-Density Foam model from the Abaqus software
library.5 This material model is intended for highly com-
pressible elastomeric foams. The material behavior was
specified directly via unaxial stress-strain curves for
tension and compression (Figure 2). In the case of ten-
sion the unaxial stress-strain behavior was described via
a curve added in the form:
   

( ) . .
. .
= ⋅ ⋅ − ⋅ ⋅ +
+ ⋅ ⋅ − ⋅
435 10 8 76 10
6 09 10 144 10
9 3 8 2
7 4
(12)
The material is fully damaged after reaching the ten-
sile strength RmT. The compressive behavior of the foam
core was described as an ideally elastoplastic material
T. MANDYS et al.: PROGRESSIVE FAILURE ANALYSIS OF COMPOSITE SANDWICH BEAM ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 593–597 595
Figure 1: The principle of material degradation in the principal direc-
tion 1
Slika 1: Na~elo degradacije materiala v glavni smeri 1
Figure 3: The resultant force-displacement diagrams of the composite
skin types A (top), B (center) and C (bottom)
Slika 3: Diagrami sila – raztezek kompozitne skorje A (na vrhu), B (v
sredini) in C (spodaj)
Figure 2: Tensile and compressive unaxial stress-strain behavior of
foam core
Slika 2: Vedenje jedra iz pene pri enoosni natezni in tla~ni obreme-
nitvi
with the Young’s modulus E and the yield limit RmC. The
material parameters are summarized in Table 2.
Table 2: Material parameters of foam core
Tabela 2: Parametri materiala pene v jedru
E RmC RmT f  U
MPa MPa MPa kg/m3 - -
50 1.2 1.5 60 0.0 0.53
2.3 Numerical simulations and results
The simulations were performed as quasi-static expli-
cit analyses in the FEM software Abaqus 6.11 using
finite-strain theory. The numerical models of the outer
skin and the foam core were meshed using 8-node solid
elements (element type C3D8R).
Figure 3 shows the resulting force-displacement
dependencies from averaged experiments and the nume-
rical simulations for specimens of type A, B and C. The
resulting force-displacement diagrams of the tensile test
and the force-deflection diagram of three-point bending
test of the foam core is shown in Figure 4.
The numerical model of the sandwich beam was cre-
ated as a fully contact problem of four bodies – sandwich
beam and three supports. The friction between the sand-
wich beam and supports has been neglected.
The failure of the upper composite skin occurred in
compression in principal direction 2 during loading in
place under the center support, both in the case of
experiments and the numerical simulation. This situation
is shown in Figure 5. The comparison of the force-
deflection diagrams for the three-point bending tests of
the sandwich beam is shown in Figure 6.
3 CONCLUSION
The user-defined material model describing the
non-linear orthotropic elastic behaviour, considering the
progressive failure analysis, was implemented in the
FEM system Abaqus. The material parameters of the
composite outer skin of the sandwich panel were
T. MANDYS et al.: PROGRESSIVE FAILURE ANALYSIS OF COMPOSITE SANDWICH BEAM ...
596 Materiali in tehnologije / Materials and technology 48 (2014) 4, 593–597
Figure 6: The force-deflection diagram of the three-point bending test
of the resulting sandwich beam
Slika 6: Diagram sila – upogib pri trito~kovnem upogibnem preizkusu
nosilca
Figure 5: a) The resulting sandwich beam subjected to a three-point
bending test and b) the numerical model of loaded sandwich beam
Slika 5: a) Sestavljen nosilec, izpostavljen trito~kovnemu upogibu in
b) numeri~ni model obremenjenega sestavljenega nosilca
Figure 4: a) The resulting force-displacement diagrams of the tensile
test and b) the force-deflection diagram of the three-point bending test
of the foam core
Slika 4: a) Diagrami sila – raztezek pri nateznem preizkusu in b)
diagram sila – uklon pri trito~kovnem upogibnem preizkusu jedra iz
pene
identified using the mathematical optimization method.
In the case of the foam core the Low-Density Foam
material model from FEM software library was used.
The obtained material parameters of the composite skin
and foam core were verified via a three-point bending
test of the resulting sandwich beam. The future work will
focus on low-velocity impact events involving sandwich
plates.
Acknowledgement
The work has been supported by the European
Regional Development Fund (ERDF), project "NTIS –
New Technologies for Information Society", European
Centre of Excellence, CZ.1.05/1.1.00/02.0090, the stu-
dent research project of Ministry of Education of Czech
Republic No. SGS-2013-036 and the project of Grant
Agency of Czech Republic No. GA^R P101/11/0288.
4 REFERENCES
1 T. Kroupa, V. La{, R. Zem~ík, Improved nonlinear stress-strain rela-
tion for carbon-epoxy composites and identification of material para-
meters, Journal of Composite Materials, 45 (2011) 9, 1045–1057
2 V. La{, R. Zem~ík, Progressive damage of undirectional composite
panel, Journal of Composite Materials, 42 (2008) 1, 22–44
3 C. F. Yen, Ballistic Impact modeling of Composite materials, 7th
International LSDyna User’s Conference, Dearborn, Michigan, 2006
4 R. Zem~ík, V. La{, T. Kroupa, H. Pur{, Identification of material cha-
racteristics of sandwich panels, Bulletin of Applied Mechanics, 26
(2011), 26–30
5 Abaqus 6.11 Documentation, Dassault Systèmes Simulia Corp., 2011
T. MANDYS et al.: PROGRESSIVE FAILURE ANALYSIS OF COMPOSITE SANDWICH BEAM ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 593–597 597

P. SLÁMA et al.: INFLUENCE OF HEAT TREATMENT ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
INFLUENCE OF HEAT TREATMENT ON THE
MICROSTRUCTURE AND MECHANICAL PROPERTIES
OF ALUMINIUM BRONZE
VPLIV TOPLOTNE OBDELAVE NA MIKROSTRUKTURO IN
MEHANSKE LASTNOSTI ALUMINIJEVEGA BRONA
Peter Sláma, Jaromír Dlouhý, Michal Kövér
COMTES FHT a.s., Prumyslova 995, 334 41 Dobrany, Czech Republic
peter.slama@comtesfht.cz
Prejem rokopisa – received: 2013-10-03; sprejem za objavo – accepted for publication: 2013-10-24
The paper deals with the influence of heat treatment (annealing, quenching and aging) on the microstructure and mechanical
properties of pressed bars made from the CuAl10Ni5Fe4 alloy. The microstructures were observed in light and scanning
electron microscopes. The appearance and area fractions of the  and  phases and their influence on the mechanical properties
were examined. Using DSC and EBSD methods, the presence of the 2 phase was monitored, as it is a very hard and brittle
phase that impairs the material’s corrosion resistance. The results were compared with data for the high-alloyed CuAl14Fe5
aluminium bronze, in which high hardness and wear resistance are imparted by the 	 + 2 phases.
Keywords: aluminium bronze, heat treatment, microstructure, mechanical properties, DSC, EBSD, 2-phase
^lanek obravnava vpliv toplotne obdelave (`arjenje, kaljenje in staranje) na mikrostrukturo in mehanske lastnosti stiskanih palic
iz zlitine CuAl10Ni5Fe4. Mikrostruktura je bila opazovana s svetlobnim in vrsti~nim elektronskim mikroskopom. Preu~evan je
bil pojav in povr{inski dele` faze  in  ter njun vpliv na mehanske lastnosti. Z metodama DSC in EBSD je bila pregledana
prisotnost trde in krhke 2-faze, ki poslab{a korozijsko odpornost materiala. Rezultati so bili primerjani s podatki za mo~no
legiran CuAl14Fe5 aluminijev bron, pri katerem (	 + 2)-faza zni`ujeta veliko trdoto in odpornost proti obrabi.
Klju~ne besede: aluminijev bron, toplotna obdelava, mikrostruktura, mehanske lastnosti, DSC, EBSD, 2-faza
1 INTRODUCTION
CuAl10Ni5Fe4 aluminium bronze is a copper alloy
that retains its high strength even at elevated tempera-
tures and which possesses a good corrosion resistance
and a high wear resistance.1–3 It has a two-phase micro-
structure consisting of the poor-formability  phase and
the high-temperature 	 phase, which exhibits excellent
hot formability. The addition of nickel increases the
alloy’s strength without diminishing its excellent ducti-
lity, toughness and corrosion resistance. Typical applica-
tions of the alloy include valve seats, plunger tips,
marine engine shafts, valve guides, aircraft components
and pump shafts. The alloy exhibits limited cold forma-
bility (due to rapid work hardening), but excellent hot
formability in the  + 	 region. The recommended form-
ing temperatures are between 700 °C and 900 °C.1,2
Depending on the cooling rate and the subsequent
heat treatment, the 	 phase may undergo a martensitic
transformation into the unstable phase 	’, which, being
very hard and brittle, enhances the strength and reduces
the ductility of the material. In addition, other phases are
found in the microstructure, which are termed , which
consist mostly of Fe and Al or Ni,3–5 or the 2 phase
known to occur in Cu-Al binary alloys. The 2 phase in
alloys containing less than 11.8 % aluminium forms
during slow cooling or in the course of annealing at
temperatures below 565 °C. These phases increase the
strength and reduce the ductility of the alloy. The
microstructure therefore consists of the  phase and the
 +  + 2 eutectoid, as shown in the phase diagram in
Figure 1a.6
Several types of the kappa phase have been classi-
fied: the sources3–5,7,8 list four types, denoted as I, II,
III, IV. The I phase forms large dendrite-shape particles
that are rich in Fe, which consist of Fe3Al with Cu and
Ni and the cubic structure DO3. The phase II forms
smaller, globular particles with identical composition
Materiali in tehnologije / Materials and technology 48 (2014) 4, 599–604 599
UDK 669.35:621.785:620.181 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 48(4)599(2014)
Figure 1: a) Vertical section of the phase diagram of nickel aluminum
bronze with 5 % Ni and 5 % Fe, b) vertical section of phase diagram
of Cu-Al-Fe at 3 % Fe
Slika 1: a) Navpi~ni prerez faznega diagrama nikelj – aluminijev bron
s 5 % Ni in 5 % Fe, b) navpi~ni prerez faznega diagrama Cu-Al-Fe pri
3 % Fe
and structure to I. The phase (precipitates) III are
NiAl-based and have the cubic B2 structure and a plate
or lamellar shape. The IV particles are small precipitates
based on Fe3Al. The composition of the 2 phase is
Cu9Al4 and its structure is the cubic P-43m type.
Where tough requirements for hardness and wear
resistance are to be met, high-alloyed Cu-Al-Fe alumi-
nium bronzes are used with an aluminium level above
13 % and an iron content of up to 5 %. The micro-
structure of these alloys consists of two phases as well,
as shown in the phase diagram in Figure 1b9:  and 2.
The high hardness of these alloys is due to the latter, the
2 phase, which also exhibits a high hardness and brittle-
ness, but a lower corrosion resistance. Owing to their
hardness, excellent compressive strength and wear
resistance, and sliding properties, these alloys are used
for forming and drawing stainless steels and for worm
wheels, valve guides and seats.
2 EXPERIMENTAL
Hot-pressed bars with a diameter of 32 mm from an
alloy containing 10 % Al, 5 % Ni and 4 % Fe were used
for the experiment. The material is denoted as
CuAl10Ni5Fe4 (CW307G) according to EN 12163.
Specimens of bars were used for investigating the
influence of annealing between 500 °C 30 min and 850 °C
30 min, quenching from 930 °C and aging of the
quenched specimens at 30 °C 30 min to 400 °C 30 min
upon the microstructure and mechanical properties.
Attention was also paid to the influence of the slow
cooling in the furnace and the subsequent rapid cooling
after removal from the furnace. The results were com-
pared with data for the high-alloyed CuAl14Fe5
aluminium bronze (sample of a worm wheel).
The chemical compositions of both alloys are given
in Table 1.
The microstructures were observed on conventional
metallographic sections in light (LM) and scanning elec-
tron microscopes (SEM). The HV hardness of the mate-
rials was measured. The evolution of the microstructure
and the formation of the precipitates and phases were
monitored using DSC (Differential Scanning Calori-
metry) with a LINNSEIS DSC PT 1600 instrument and
by means of EBSD (Electron Backscatter Diffraction) in
a JEOL JSM 7400F scanning electron microscope
equipped with a HKL-NORDLYSS camera and Chan-
nel5 software.
3 RESULTS AND DISCUSSION
3.1 CuAl10Ni5Fe4 Alloy
3.1.1 HV Hardness
The HV30 hardness of the as-received specimens
(AR) upon annealing was measured. The results are
shown in Figure 2. The graph shows the hardness levels
after annealing of the as-received material and upon
annealing (tempering) of the quenched material. The
hardness of the as-received material is 265 HV30.
Quenching increases this value to 305, and aging, to 400
(Table 2).
Table 2: HV30 hardness
Tabela 2: Trdota HV30
Condition Annealed
cooling Furnace Air
AR 265
500 °C 262 260
600 °C 255 253
650 °C 256 254
700 °C 245 253
750 °C 215 239
800 °C 198 234
850 °C 184 230
Quenched 930 °C 30 min 305
Quenched + Aged 400 °C 30 min 400
The hardness decreases during annealing; first
slowly, while the temperature is below 700 °C, but more
rapidly when it reaches 750 °C. A greater drop in hard-
ness was seen in the specimens cooled slowly in a fur-
nace than in those cooled in air.
P. SLÁMA et al.: INFLUENCE OF HEAT TREATMENT ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
600 Materiali in tehnologije / Materials and technology 48 (2014) 4, 599–604
Figure 2: HV30 hardness of annealed specimens
Slika 2: Trdota HV30 vzorcev po `arjenju
Table 1: Chemical composition of alloys in mass fractions (w/%)
Tabela 1: Kemijska sestava zlitin v masnih dele`ih (w/%)
Alloy Cu Al Ni Fe Mn Sn Si Zn Pb
CuAl10Ni5Fe4 balance 9.9 5.4 4.5 0.77 0.02 0.13 0.02 0.01
CuAl14Fe5 balance 14.5 0.12 4.9 0.97 0.06 0.22 0.006 0.07
3.1.2 Microstructure
The micrographs of the specimens prior to etching
are shown in Figure 3. Unetched sections show grey
globular particles of the II phase (according to EDS ana-
lysis, they contain 55 % iron, 11 % aluminium and
nickel, copper, manganese and silicon), light grains of
the phase  and a eutectoid consisting of  +  or a
balance of 	. Annealing at the quenching temperature of
930 °C dissolves the  phase in the eutectoid (except the
globular II) and the 	 phase transforms to the non-equi-
librium 	’ phase,10 as seen in Figure 3b.
The microstructures upon etching are shown in
Figures 4 and 5. Etching reveals light grains of the phase
, grey to light-blue globular particles of the phase II
and the dark eutectoid  + III.7 In the as-received state,
the  grain size is very small: 1 μm to 2 μm. With
increasing annealing temperature, the size of the  grains
grows, as does the size of the eutectoid areas between
them. Slow cooling in the furnace leads to coarsening of
the  precipitates and to the formation of larger amounts
of lamellar III precipitates in the eutectoid. The phase 
also precipitates along the  grain boundaries. With
increasing  grain size the strength and hardness de-
crease. The greatest drop is seen in specimens annealed
at 850 °C with slow cooling. However, the elongation
decreases as well, which may be attributed to  preci-
pitating more frequently at the  grain boundaries.
Rapid cooling leads to a higher proportion of the
eutectoid at the expense of the  proportion. In addition,
a certain fraction of 	 phase is retained in the micro-
structure (more dark areas in the eutectoid). This is
shown in Figures 4c and 4d, where short etching with
NH4OH+H2O2 revealed an orange 	 phase, a bright 
phase and the  phase and eutectoid in light-blue colour.
P. SLÁMA et al.: INFLUENCE OF HEAT TREATMENT ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 599–604 601
Figure 4: Micrographs of etched specimens: a) as-received condition,
b) condition upon annealing at 750 °C, furnace, c) condition upon
annealing at 850 °C, furnace, d) condition upon annealing at 850 °C,
air
Slika 4: Mikrostruktura jedkanih vzorcev: a) dobavljeno stanje, b) po
`arjenju na 750 °C, ohlajanje v pe~i, c) po `arjenju na 850 °C, ohlaja-
nje v pe~i, d) po `arjenju na 850 °C, ohlajanje na zraku
Figure 5: SEM micrographs of etched specimens: a) condition upon
annealing at 800 °C, furnace, b) condition upon annealing at 800 °C,
air
Slika 5: SEM-posnetka jedkanih vzorcev: a) po `arjenju na 800 °C,
ohlajanje v pe~i, b) po `arjenju na 800 °C, ohlajanje na zraku
Figure 3: Micrographs of unetched specimens: a) as-received condi-
tion (AR), b) 17 – upon quenching
Slika 3: Mikrostruktura nejedkanih vzorcev: a) dobavljeno stanje
(AR), b) 17 – po kaljenju
(a)
(b)
In contrast to the products of slow cooling, the 
precipitates are finer. Globular precipitates of IV formed
as well. The decrease in the strength and hardness is not
as great as on slow cooling, which can be explained by a
greater proportion of eutectoid and residual 	 in the
microstructure.
The SEM micrographs also confirm that slow cooling
leads to coarser  precipitates than rapid cooling (Fig-
ures 5a and 5b).
Upon quenching and aging, fine IV precipitates form
in the martensitic 	’ phase, as shown in the light (Figure
6a) and scanning electron (Figure 6b) micrographs. The
martensitic 	’ phase and fine precipitates provide the
quenched and aged material with high hardness.
3.2 CuAl14Fe5 Alloy
3.2.1 Hardness and microstructure
The specimen of CuAl14Fe5 alloy was made from a
sample taken from a worm-wheel, in which very high
hardness was required. The alloy corresponds to the
AMPCO® 22 material by AMPCOMETAL.11 According
to the company information, it contains 	 and 2 phases
and has a nominal hardness of 332 HB30.
The specimen was used for an annealing trial at 950 °C
30 min with subsequent cooling in air. The hardness of
the as-received material was 422 HV10. During
annealing, it rose to 462 HV10. Light micrographs are
shown in Figures 7a and 7c, and a scanning electron
micrograph is shown in Figure 7b. The microstructure
consists of light-blue grains of 2 embedded in bright 	
phase matrix, of large dark II particles in the 2 phase
and of minute IV precipitates within the 	 phase. In the
as-received specimen, the microhardness of the 2 phase
is 573 HV0.05 and that of the 	 phase reaches 341
HV0.05. The presence of the 	 (	’) phase can be de-
duced from the material’s high hardness. Due to anneal-
ing, the proportion of the 2 phase increased and, as a
result of its high hardness, so did the hardness of the
P. SLÁMA et al.: INFLUENCE OF HEAT TREATMENT ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
602 Materiali in tehnologije / Materials and technology 48 (2014) 4, 599–604
Figure 7: LM and SEM micrographs of CuAl14Fe5: a) as-received
condition, unetched, HV0.05 hardness testing indentations (LM), b)
as-received condition, etched (SEM), c) annealed at 950 °C 30 min,
unetched (LM)
Slika 7: Svetlobni in SEM-posnetek CuAl14Fe5: a) dobavljeno stanje,
nejedkano, odtiski trdote HV0,05 (SM), b) dobavljeno stanje, jedkano
(SEM), c) `arjeno pri 950 °C 30 min, nejedkano (SM)
Figure 6: Microstructures of quenched and aged specimen, etched: a)
quenched and aged at 400 °C (LM), b) quenched and aged at 400 °C
(SEM)
Slika 6: Mikrostruktura hitro ohlajenega in staranega vzorca, jedkano:
a) hitro ohlajeno in starano pri 400 °C (SM), b) hitro ohlajeno in
starano pri 400 °C (SEM)
entire specimen. The SEM observation reveals needles of
non-equilibrium 	’ phase in the 	 phase (Figure 7b).
3.3 DSC Analysis
Using DSC, the precipitation of the  phases and the
	   + 2 eutectoid transformation were monitored.12,13
DSC curves were measured at heating and cooling rates
of 10 K/min. Figure 8 shows the DSC curves for run 2
(after one cycle heating–cooling). The focus was the
changes taking place at temperatures around 500 °C, in
the vicinity of the eutectoid transformation 	  +2 at
565 °C. Small peaks were recorded for the
CuAl10Ni5Fe4 alloy. Temperature peaks between
930 °C and 830 °C on the cooling curve correspond to
the formation of  phases. (+	  +	+) The peak at
approx. 500 °C corresponding to the eutectoid transfor-
mation is very indistinct.
According to7 at 930 °C the globular II phase (	 
	+II transformation) precipitates from the 	 phase and at
800 °C the +III lamellar eutectoid forms (  +III).
At 850 °C fine IV precipitates form in the  phase.
In the CuAl14Fe5 alloy, the peaks due to the forma-
tion of the  and 2 phases are much more pronounced.
At temperatures below 800 °C, two peaks can be ob-
served, which could indicate the formation of the  and
2 phases in the 	 phase. The 500 °C peak may corres-
pond to the 	   + 2 eutectoid transformation.
The transformation (formation of phases) at tempera-
tures below 800 °C can be described according to Figure
1b as follows: 	  	+  	++2  ++2.
The peaks at temperatures of 300–400 °C correspond
to the martensitic transformations of the 	 phase (forma-
tion of 	1, 	’).10,13,14
3.4 Identification of phases using EBSD
As the phases of the different structures produce
different diffraction patterns (EBSP), the EBSD analysis
can be used for identifying the various phases.15 The
EBSD analysis was aimed at identifying the  and 2
phases in both aluminium bronzes. The phase maps and
band contrast images are shown in Figure 9.
With the CuAl10Ni5Fe4 alloy, a higher band contrast
was obtained (better quality of EBSP) in phases identi-
fied as  (large grey-coloured areas) and II – Fe3Al
(small, bright spots). Due to the low band contrast, the
phases in the eutectoid were not identified, most pro-
bably owing to the lamellar structure of  + III.
In the CuAl14Fe5 alloy, the phases identified in-
cluded 2 (large grey areas) and II (small bright spots).
No regions with the 	 phase were identified, probably
due to the presence of needles of non-equilibrium mar-
tensitic 	’ phase within the 	 phase.
P. SLÁMA et al.: INFLUENCE OF HEAT TREATMENT ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
Materiali in tehnologije / Materials and technology 48 (2014) 4, 599–604 603
Figure 9: Band contrast and phase maps: a) alloy CuAl10Ni5Fe4,
after annealing at 800 °C, air; grey phase – phase , light phase –
phase II, b) alloy CuAl14Fe5, grey phase – phase 2, light phase –
phase II
Slika 9: Trakast kontrast in razporeditev faz: a) zlitina CuAl10Ni5Fe4
po `arjenju na 800 °C, ohlajeno na zraku; siva faza – faza , svetla
faza – faza II, b) zlitina CuAl14Fe5, siva faza – faza 2, svetla faza –
faza II
Figure 8: DSC curves for heating and cooling: a) CuAl10Ni5Fe4
alloy, b) CuAl14Fe5 alloy
Slika 8: DCS-krivulje pri ogrevanju in ohlajanju: a) zlitina CuAl10
Ni5Fe4, b) zlitina CuAl14Fe5
4 CONCLUSIONS
The outcomes of the investigation of the effect of
heat treatment on the microstructure and properties of
pressed bars from CuAl10Ni5Fe4 alloy can be sum-
marized as follows:
• Microstructure of the bars upon hot pressing is fine,
consisting of  phase grains and the  +  eutectoid.
Owing to rapid cooling, the 	 phase may form as
well.
• Annealing causes  grains and the eutectoid regions
with III-type precipitates to coarsen.
• Maximum hardness was obtained by quenching and
aging at 400 °C: 400 HV30. This is due to the fine
dispersion of particles  in the martensitic 	’ phase.
• The 2 phase was not detected in the CuAl10Ni5Fe4
alloy by DSC or EBSD analysis.
CuAl10Ni5Fe4 Alloy
• The hardness of the CuAl14Fe5 is higher. It is deter-
mined by the proportion of the very hard 2 phase
(with a microhardness of more than 550 HV0.05) in
the 	 phase. The as-received specimen had a hardness
of 422 HV10. Upon annealing, with subsequent air
cooling, the fraction of 2 increased. The hardness
increased as well: to a level of 462 HV10.
• The microstructure consists of 	 phases with a
greater fraction of 2, particles of II within the 2
phase and minute IV dispersoids in the 	 phase.
• The presence of  and 2 phases was confirmed by
DSC and EBSD analyses.
Acknowledgement
The results presented in this paper were developed
under the project "West-Bohemian Centre of Materials
and Metallurgy" CZ.1.05/2.1.00/03.0077 co-funded from
European Regional Development Fund.
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ves of Foundry Engineering, 13 (2013) 3, 72–79
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P. SLÁMA et al.: INFLUENCE OF HEAT TREATMENT ON THE MICROSTRUCTURE AND MECHANICAL PROPERTIES ...
604 Materiali in tehnologije / Materials and technology 48 (2014) 4, 599–604
[OLSKI CENTER [KOFJA LOKA SE JE PREDSTAVIL
TUDI PREDSEDNIKU RS BORUTU PAHORJU
Po uspe{ni vzpostavitvi Medpodjetni{kega
izobra`evalnega centra (MIC) in energetski sanaciji
je [olski center [kofja Loka 18. 4. 2014 obiskal
predsednik RS Borut Pahor. Predsednik se je sre~al z
direktorjem [olskega centra ([C) g. Martinom Piv-
kom, z vodjo MIC-a g. Alojzom Kokaljem in z ravna-
telji {olskega centra. »Kot predsednik ob razli~nih
te`avah, s katerimi se sre~ujem, potrebujem eno krepko
dozo optimizma. To je moje dana{nje sre~anje z mladi-
mi,« je obisk komentiral Borut Pahor.
Na [olskem centru [kofja Loka so predsedniku
Pahorju predstavili usposabljanja za tehni{ke poklice,
predvsem iz strojni{tva in lesarstva: nagrajene projekte iz
lesarstva (modularno pisarno, kopalni{ko opremo, klub-
sko mizico in leseno kolo), delo z robotom IRB 120 in
avtomatizacijo procesov. Vodja projekta energetske
sanacije Primo` Praper s podjetja EUTRIP je predsed-
niku Pahorju predstavil sistem digitalnega energetskega
monitoringa in spremljanje doseganja energetske u~inko-
vitosti. Predsednik se je sprehodil tudi med dijaki, ki so
izvajali postopke CAD-na~rtovanja in 3D-modeliranja
(3D-tiskalnik in 3D-skener), v laboratoriju modernih
CNC-tehnologij pa so mu predstavili inovativni pripo-
mo~ek za slepe, ki je nastal v {olskem centru. Predsednik
se je ustavil tudi ob predstavitvi sne`nega topa in podjet-
ni{kega projekta e-skiro, zapeljal pa se je tudi z elektri-
~nim vozilom Smart, ki je nastal v delavnici {olskega
centra.
»Prav bi bilo, da bi med na{imi ljudmi bolj odme-
valo, kako inteligentno, ambiciozno, inovativno mlado
generacijo imamo. Tukaj so vsi bodo~i in`enirji, tudi
tisti, ki bodo {li v proizvodnjo, imajo tehni{ki poklic in
jih ne skrbi za prihodnost zaposlitve. Dobro bi bilo, da se
vidi te `arke upanja. To so mladi samozavestni ljudje, ki
delajo stvari, ki se zdijo nemogo~e. Zdaj odhajam spet z
`arom do `ivljenja,« je ob slovesu povedal Borut Pahor,
ki je tudi poudaril, da danes ni ve~ mogo~e izobraziti
~loveka za eno podro~je, na katerem bo delal vse `iv-
ljenje, ampak da je potrebna iznajdljivost, inovativnost,
prilagodljivost. »Najprej izobra`ujemo osebnosti. Mi {e
vedno mislimo, da je glavni proces {olskega procesa
podajanje znanja. Ni tako. Informacijo lahko danes
dobi{ kjer koli. Problem je, kako to informacijo razumeti
in kako jo vpeti v na~rte za prihodnje. To jim lahko dajo
profesorji. Ne le s tehni~nimi znanji, ampak da nastanejo
to kreativne osebnosti, da mislijo s svojo glavo, da pro-
fesorji spodbujajo mlade k razmi{ljanju, da u~enci lahko
re~ejo: šJaz sem bolj{i od tebe’, in da so tudi u~itelji
zadovoljni, ko vidijo, da so u~enci bolj{i od njih.
Strokovne {ole lahko veliko ponudijo.«
Direktor Martin Pivk, ki je gostil predsednika
dr`ave, je povedal: »Za obisk predsednika nismo pri-
pravili ni~ posebnega. Pokazali smo samo tisto, kar se
vsak dan dogaja v {olskem centru. Tako da lahko re~emo,
da ti prikazi niso odsev enodnevnega obiska, ampak od-
sev sodelovanja dijakov ter pomo~i mentorjev, u~iteljev
pri njihovem vsakdanjem delu.« Direktor je {e izrazil
Materiali in tehnologije / Materials and technology 48 (2014) 4, 605–606 605
UDK 330.322.3:37(497.12) ISSN 1580-2949
Dogodki/Events MTAEC9, 48(4)605(2014)
Slika 2: Predsednik Borut Pahor in direktor [C Martin Pivk med pred-
stavitvijo elektri~nega vozila Smart
Slika 1: Dijaki predsedniku Borutu Pahorju predstavljajo delo z robo-
tom
veselje, »da imamo tako uspe{ne in prisr~ne dijake, ki
imajo znanje. To je dokaz, da ima strokovno in poklicno
izobra`evanje v [kofji Loki {e velike mo`nosti.«
Vodja MIC-a Alojzij Kokalj je ob tem dogodku
povedal: »Veseli me, da so vedno bolj prepoznavni
projekti, ki v sebi skrivajo veliko realnega znanja, ki so
povezani s podjetji in tehnologijami, ki se dogajajo na
trgu ali v industriji v tem trenutku. Pravi u`itek je
spremljati mlade na tej ustvarjalni poti, ko opazuje{,
kako raste njihova motivacija ob konkretnih rezultatih,
kako se mladi razvijajo ob problemih, s katerimi se
sre~ujejo. Ob tem se ~lovek vpra{a, ali jim lahko sploh {e
kaj ve~ damo. To je odli~na popotnica za njihovo `iv-
ljenjsko in karierno pot. Veseli me, da je tudi predsednik
dr`ave ob dana{njem obisku prepoznal potrebo po
predstavitvah in potrditvah tovrstnih dose`kov mladih.«
[olski center [kofja Loka nadgrajuje svoje izobra`e-
valno-razvojno poslanstvo in odmevne projekte. V letu
2013 je v sodelovanju z In{titutom Metron `e predelal in
v uporabo predal elektri~ni avto, v septembru 2013 pa je
kon~al investiciji gradnje Medpodjetni{kega izobra`e-
valnega centra (MIC) in energetske sanacije {olskih
objektov. V okviru projekta vzpostavitve MIC-a [kofja
Loka, ki se je za~el marca 2009, je bil na povr{ini 2200
m2 zgrajen izredno napredno opremljen objekt, ki
omogo~a izvajanje kakovostnega poklicnega in strokov-
nega izobra`evanja, usposabljanja ter povezovanja z
gospodarstvom. Vzpostavitev MIC-a je bila sofinan-
cirana iz Evropskega sklada za regionalni razvoj, projekt
energetske sanacije pa iz Kohezijskega sklada. Skupna
vrednost obeh projektov je bila pribli`no 5,5 milijona
evrov.
Avtor fotografij: Lado Bukovec, [olski center [kofja Loka
606 Materiali in tehnologije / Materials and technology 48 (2014) 4, 605–606