VSEBINA – CONTENTS
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Material failure of an AISI 316L stainless steel hip prosthesis
Napake materiala v kol~ni protezi iz nerjavnega jekla AISI 316L
M. Godec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85
A comparison of the corrosion behaviour of austenitic stainless steels in artificial seawater
Primerjava korozijskih lastnosti avstenitnih nerjavnih jekel v simulirani morski vodi
A. Kocijan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 91
Influence of the foaming precursor’s composition and density on the foaming efficiency, microstructure development and
mechanical properties of aluminium foams
Vpliv sestave in gostote prekurzorjev za penjenje na u~inkovitost penjenja ter razvoj mikrostrukture in mehanskih lastnosti aluminijskih
pen
V. Kevorkijan, S. D. [kapin, I. Paulin, B. [u{tar{i~, M. Jenko, M. La`eta. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 95
Multi-objective optimization of the cutting forces in turning operations using the Grey-based Taguchi method
Multi namenska optimizacija stru`enja z uporabo Taguchi metode na Grey podlagi
Y. Kazancoglu, U. Esme, M. Bayramo

glu, O. Guven, S. Ozgun . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 105
Thermodynamic conditions for the nucleation of boron compounds during the cooling of steel
Termodinami~ni pogoji za nukleacijo borovih spojin pri ohlajanju jekla
Z. Adolf, J. Ba`an, L. Socha . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111
A micro-damage investigation on a low-alloy steel tested using a 7.62-mm AP projectile
Raziskava mikropo{kodb malolegiranega jekla po preizkusu s kroglo AP 7,62 mm
T. Demir, M. Übeyli . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 115
Activation of polymer polyethylene terephthalate (PET) by exposure to CO2 and O2 plasma
Aktivacija polimera polietilentereftalata (PET) s CO2- ali O2-plazmo
A. Vesel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 121
The impact on rigid PVC pipes: a study of the correlation between the length of the crazed zone and the area of the impacted
region
Udar togih PVC-cevi: {tudija korelacije med dol`ino razpokane zone in povr{ino zone udara
C. B. Fokam, M. Chergui, K. Mansouri, M. El Ghorba, M. Mazouzi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 125
A comparative analysis of theoretical models and experimental research for spray drying
Primerjalna analiza teoreti~nih modelov in eksperimentalna raziskava razpr{ilnega su{enja
D. Tolmac, S. Prvulovic, D. Dimitrijevic, J. Tolmac . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 131
Creep resistance of microstructure of welds of creep resistant steels
Odpornost proti lezenju pri mikrostrukturi zvarov jekel, odpornih proti lezenju
F. Vodopivec, M. Jenko, R. Celin, B. @u`ek, D. A. Skobir . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 139
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Effect of the martensite volume fraction on the machining of a dual-phase steel using a milling operation
Vpliv volumenskega dele`a martenzita na obdelavo dvofaznega dualnega jekla z rezkanjem
O. Topçu, M. Übeyli, A. Acir . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 145
Degradation of a Ni-Cr-Fe alloy in a pressurised-water nuclear power plant
Degradacija zlitin Ni-Cr-Fe v tla~novodnih jedrskih elektrarnah
R. Celin, F. Tehovnik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 151
Investigation into the mechanical properties of micro-alloyed as-cast steel
Raziskave mehanskih lastnosti mikrolegiranih jekel
B. Chokkalingam, S. S. M. Nazirudeen, S. S. Ramakrishnan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 159
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 45(2)83–166(2011)
MATER.
TEHNOL.
LETNIK
VOLUME 45
[TEV.
NO. 2
STR.
P. 83–166
LJUBLJANA
SLOVENIJA
MAR.–APR.
2011
The effect of electromagnetic stirring on the crystallization of concast billets
Kristalizacija kontinuirno ulitih gredic v elektromagnetnem polju
K. Stransky, F. Kavicka, B. Sekanina, J. Stetina,V. Gontarev, J. Dobrovska. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 163
19. KONFERENCA O MATERIALIH IN TEHNOLOGIJAH, 21. – 23. november 2011, Portoro`, Slovenija
19th CONFERENCE ON MATERIALS AND TECHNOLOGY, November 21–23, 2011, Portoro`, Slovenia . . . . . . . . . . . . . . . . . . . . . 167
M. GODEC: MATERIAL FAILURE OF AN AISI 316L STAINLESS STEEL HIP PROSTHESIS
MATERIAL FAILURE OF AN AISI 316L STAINLESS
STEEL HIP PROSTHESIS
NAPAKE MATERIALA V KOL^NI PROTEZI IZ NERJAVNEGA
JEKLA AISI 316L
Matja` Godec
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
matjaz.godec@imt.si
Prejem rokopisa – received: 2011-03-15; sprejem za objavo – accepted for publication: 2011-03-25
One of the major successes of modern surgery is the total replacement of joints by timplantation of prosthetic components by
which the pain is relieved and deformity is corrected. The aseptic loosening of a prosthetic-joint component is the most common
reason for joint-revision surgery. Furthermore, it is thought that wear particles are one of the major contributors to the
development and perpetuation of aseptic loosening. It was ascertained that besides wear particles originated from polyethylene
acetabular cup, the main reason of loosening hip prosthesis was due to the material failure. Based on stem material examination
by different techniques it was found out that improper microstructure caused additional abrasion by hard particles and aseptic
loosening took place.
Keywords: hip prosthesis, aseptic loosening, wear, stainless steel, polyethylene, non-metallic inclusions
Velik uspeh moderne kirurgije je med drugimi tudi popolna nadomestitev sklepov s proteti~nimi nadomestki, s katerimi se
omilijo bole~ine in korigira deformacija. Pogosto pride do asepti~nega omajanja in potrebe po ponovni operaciji. Vzrok za to je
ravno v obrabi in nastanku drobnih obrabnih delcev, ki ve~inoma vodijo do komplikacij in omajanja proteze. V raziskavi smo
pokazali, da je poleg delcev, ki so izvirali iz polietilenske ~a{ice bil glavni razlog omajanja ravno v napaki materiala. Na osnovi
analize stebla materiala z razli~nimi tehnikami smo pokazali, da je neustrezna mikrostruktura povzro~ila dodatno abrazijo,
zaradi trdih delcev in tako povzro~ila omajanje.
Klju~ne besede: kol~na, proteza, asepti~no omajanje, obraba, nerjavno jeklo, polietilen, nekovinski vklju~ki
1 INTRODUCTION
A total hip replacement is a surgical procedure
whereby the diseased cartilage and bone of the hip joint
is surgically replaced with artificial materials. An
increasingly ageing population means that absolute
numbers of people with a predilection for osteoarthritis
is set to rise 1. Hip joint replacement is usually done in
people age 60 and older. It is expected that total hip
replacement will increase due to population ageing.
Two types of arthroplasty are carried out nowadays
with reference to mode of fixation; cemented und
uncemented. Furthermore, two types of bearing surfaces
have been utilised in total hip replacement. Hard-on-soft
bearings have included couplings where the acetabular
liner has been polyethylene (PE) and the femoral heads
have been metallic, usually cobalt-chrome alloy, or cera-
mic. In order to assure lower abrasive wear and reduce
the rate of aseptic loosening ceramic-on-ceramic or
metal-on-metal bearing have been applied recently.
There are three main materials used for implant stem:
cobalt-chrome, stainless steel and titanium alloy.
Although cobalt-chrome alloys are extensively used for
manufacture of implants they have among good corro-
sion resistance and easily workability also some less
desirable properties. One of them is high modulus of
elasticity and not completely known long-term effects of
metallosis arising from cobalt-chrome alloy 2,3. Stainless
steel is easy to work, has high strength and low price,
however lower fatigue strength and tendency to
corrosion cannot put it on the top 4–23. One of the best is
titanium alloy due to superior biocompatibility, relatively
low elastic modulus and good corrosion performances
4,24–26. However, it is softer than other metals and cannot
be used for femoral heads due to wear 1.
Younger people who have a hip replaced may put
extra stress on the artificial hip. That extra stress can
cause it to wear out. The main cause of failure in these
patients remains loosening due to osteolysis and the
focus in future is going to be on extending the durability
and survivorship of these components in a younger
patient age group 1. The aseptic loosening of a prosthe-
tic-joint component is the most common reason for
joint-revision surgery among all group of patient.
However, there have been a few reported cases where a
fracture of the stem neck has occurred, most likely due to
the development of an inappropriate microstructure 27,28.
While abnormal wear at the bearing surface might give
rise to excessive particle generation, the bone loss in
relation to inflammation caused by these particles might
in turn result in the loosening of the device and sub-
sequent abnormal mechanical loading. In any case, wear
particles are now considered to be one of the major
contributors to the development and perpetuation of
aseptic loosening 29. Metal implants usually exhibit a low
level of wear (40–100 times lower than ultra-high-
Materiali in tehnologije / Materials and technology 45 (2011) 2, 85–90 85
UDK 669.14.018.8:620.17./.19 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 45(2)85(2011)
molecular-weight-polyethylene UHMWPE), a good
surface finish and excellent mechanical properties. The
drawbacks of metals are the metallic, electrochemical
corrosion risks (low bio-compatibility) accelerated by
the ions present in the body fluids and the presence of
oxygen 2,4,5.
In the present study an AISI 316L stainless steel total
hip prosthesis with unusual microstructure regarding the
presence of non-metallic inclusions was investigated in
order to find eventual connections with aseptic loosening
and material microstructure.
2 EXPERIMENTAL DETAILS
A 63 years old male patient underwent total hip
arthroplasty revision surgery due to aseptic loosening of
both the femoral and acetabular components. The
prosthesis was 106 months in use before exchange due to
severe pain caused by loosening of stem. The femoral
component was AISI 316L stainless steel stem (0.03
wt% C, 16–18 wt% Cr, 10–14 wt% Ni, 2–3 wt% Mo)
and UHMWPE acetabular cup cemented with poly-
methylmethacrylate (PMMA). The hip prosthesis was
immediately cleaned, using only distilled water, after
surgical removal. The specimens for the surface exami-
nation were cut from the hip stem and the polyethylene
cup and cleaned in ethanol in an ultrasonic bath. The
metal specimen for the microstructural examination was
metallographically prepared by a standard procedure
using 3 and 1 μm diamond polishing paste, and then
etched by glyceregia to reveal the grain boundaries. The
specimens for electron backscatter diffraction (EBSD)
analyses were, after polishing, given an additional polish
with colloidal silica oxide for 3 min and cleaned in
ultrasonic bath. Furthermore, ion etching was performed
in PECS 682 Gatan using 2,5 kV beam energy in rotation
mode at 50° gracing angle. The specimens were ana-
lyzed using a light microscope Microphot FXA (Nikon)
and with a FE-SEM JEOL JSM 6500F field-emission
scanning electron microscope with attached EDX (an
INCA X-SIGHT LN2 type detector, INCA ENERGY
450 software and an HKL Nordlys II EBSD camera
using Channel5 software). For the EDX analysis, a 15
kV accelerating voltage and a probe current of 0.8 nA
were used. The EBSD was performed at a 20 kV
accelerating voltage and 2.7 nA probe current. The
polyethylene specimen was covered with a 4 nm AuPd
conductive layer for a subsequent SEM examination.
3 RESULTS AND DISCUSSION
After stainless steel (AISI 316L) total hip prosthesis
had been surgical replaced, it was microscopically and
spectroscopically examined. Polyethylene acetabular cup
as well as stainless steel stem were studied together with
the cement binder in order to find out the cause of its
loosening (Figure 1). Bone cement (PMMA) is an
excellent type of fixation for the hip prostheses,
especially for elderly, relatively inactive patients who are
usually also suffering from osteoporosis. During patient
movement the wear between the polyethylene acetabular
cup and the metallic head caused formation of very fine,
irregularly shaped and elongated particles not longer
than 1 μm. Furthermore, tribological activity of the
acetabular cup and the metallic head caused tearing of
larger particles from the inner acetabular cup wall. It is
supposed that in overload situations adhesion caused
conditions that part of the acetabular cup can be pulled
out. As a result, the polished surface of the inner wall of
the polyethylene acetabular cup with 300 μm deep
cavities randomly distributed was formed (Figure 2) and
consecutively looseness the hip joint.
Accurate stainless stem surface examination showed
that original rough surface, caused by sand-blasting as a
last step in manufacturing process of the stem, in some
areas transformed into smooth surface with some rough
islands residues. During this process the metallic
particles and ions accumulate in the body tissues and
fluids. Debonding between the cement and implant
might reduce the life-time of the implant fixation. Small
M. GODEC: MATERIAL FAILURE OF AN AISI 316L STAINLESS STEEL HIP PROSTHESIS
86 Materiali in tehnologije / Materials and technology 45 (2011) 2, 85–90
Figure 2: Surface modification of the polyethylene acetabular cup’s
inner wall with deep cavities, SE image
Slika 2:. Sprememba povr{ine polietilenske ~a{ice z globokimi
jamicami, SE slika
Figure 1: Hip prosthesis together with acetabular cup and residues of
cement after being surgery removed
Slika 1: Kol~na proteza z acetabularno ~a{ico in ostanki cementa po
operaciji
M. GODEC: MATERIAL FAILURE OF AN AISI 316L STAINLESS STEEL HIP PROSTHESIS
Materiali in tehnologije / Materials and technology 45 (2011) 2, 85–90 87
Figure 6: EDX spot analysis, (a) SE image of stem specimen with
marked spots of EDX analysis, (b), (c) and (d) EDX spot analysis of
non-metallic inclusions
Slika 6: EDX to~kovna analiza stebla, (a) SE slika stebla z ozna~e-
nimi mesti EDX analize, (b), (c) in (d) EDX to~kovne analize
nekovinskih vklju~kov
Figure 4: (a) SE image of alumina particle due to residue of surface
sand blasting with marked spot of the EDX analysis; (b) EDX spot
analysis of an alumina particle
Slika 4: (a) SE slika aluminijevega oksida, ki je ostal na povr{ini po
peskanju, z ozna~enim mestom EDX analize, (b) EDX to~kovna
analiza aluminijevega oksida
Figure 5: Microstructure of stem specimen showing plenty of Al2O3
non-metallic inclusions, BE image
Slika 5: Mikrostruktura vzorca stebla, ki prikazuje veliko Al2O3
nekovinskih vklju~kov, BE slika
Figure 3: (a) Original sand blasted surface, SE image; (b) surface
modification due to wear process, SE image
Slika 3: (a) Originalna povr{ina po peskanju, SE slika; (b) sprememba
povr{ine zaradi obrabe, SE slika
cement fractures are thought to play a significant role in
the initiation of the cement failure 1. The prevailing
stress between the stem and the cement is mostly caused
by shear, whereas compressive-stress occurs in aceta-
bulum, especially among active patients. The wear parti-
cles from the cement, the metal and the polyethylene
(PE) are claimed to play a major role in the aseptic
loosening. Osteolysis without any signs of infection is
mostly related to release of PMMA debris in the tissues
by micro-motion which initiates the cement failure. As a
consequence a foreign-body reaction occurs and initiates
osteolysis and finally leads to loosening. Once metallic
particles are formed by wear, the new particles will be
generated by friction and apart from causing a mild
foreign-body reaction, such a metals debris can be
involved in the process of aseptic loosening mostly due
to cytoxine release from macrophages.
Analysed stainless steel stem indicated some of the
surface areas not affected by wear while some parts of
stem underwent significant changes due to wear process.
Figure 3 (a) shows original sand blasted surface and
Figure 3(b) demonstrated the surface which was
M. GODEC: MATERIAL FAILURE OF AN AISI 316L STAINLESS STEEL HIP PROSTHESIS
88 Materiali in tehnologije / Materials and technology 45 (2011) 2, 85–90
Figure 10: EBSD mapping (a) band contrast image (b) inverse pole
figure in X direction, (c) inverse pole figure in Y direction, (c) inverse
pole figure in Z direction
Slika 10: EBSD ploskovna analiza (a) slika kvalitete uklona (b)
inverzna polova figura v X smeri, (c) ) inverzna polova figura v Y
smeri, (c) inverzna polova figura v Z smeri
Figure 8: BE image of hip stem microstructure close to the surface
with the slip lines due to surface deformation by sand-blasting
Slika 8: BE slika mikrostrukture stebla blizu povr{ine z drsnimi
~rtami, zaradi deformacije povzro~ene s peskanjem
Figure 9: SE image of stem microstructure consist of austenite grains.
Grain boundaries are free of any carbides and other phases. Etched by
glyceregia
Slika 9: SE slika avstenitne mikrostrukture stebla. Na mejah ni
karbidov in ostalih faz. Jedkano z zlatotopko
Figure 7: EDX mapping, (a) BE image of stem specimen with
corresponding EDX images of (b) O K1, (c) Si K1 (d) Al K1, Ca
K1 and Fe K1
Slika 7: EDX ploskovna analiza, (a) BE slika vzorca stebla s
pripadajo~imi EDX slikami (b) O K1, (c) Si K1 (d) Al K1, Ca K1
in Fe K1
exposed to wear process. Rougher surface had plenty of
small mechanically produced pits visible at a higher
magnification. There were also plenty of Al2O3 particles
embedded in the material most likely due to sand
blasting procedure in order to make the surface rougher
and easier for bone tissue to grow into. Al2O3 particles
were detected by EDX in the stem surface (Figure 4).
Stainless steel stem microstructure across whole
inner and outer part of implant had plenty of non-
metallic inclusions (Figure 5). The inclusions were
analysed by EDX. The inclusions were combinations of
Al2O3, SiO2 and CaO particles and were consequences of
problems during steel making. The presence of
non-metallic oxide inclusions is a major cause of
incompatibility between the attainable and desirable
level of cleanliness in many grades of commercial steel.
Generally, inclusions degrade the mechanical properties
of the steel and thereby reduce the ductility of the cast
metal and increase the risk for mechanical and/or
corrosion failure of the final product. The number of
these inclusions was too large for the AISI 316L steel
grade particularly in an orthopaedic application. Oxide
inclusions originate from two sources: (a) residual
products resulting from intentionally added alloying
elements to deoxidize the molten steel after oxygen
treatment (endogenous or micro inclusions); (b) products
resulting from reactions between the melt and atmo-
sphere, slag, or refractory (exogenous or macro
inclusions) 30. Alumina inclusions occur as deoxidation
products in the aluminum-based desoxidation of steel.
Most grades of steel are treated with calcium using Ca-Si
alloy is certain amount of silicon is permitted in steel.
During calcium treatment, the alumina and silica
inclusions are converted to molten calcium aluminates
and silicate which are globular in shape because of the
surface tension effect. The change in inclusion compo-
sition and shape is known as the inclusion morphology
control. However, such high amount of that type of
inclusions is an evidence of wrong desoxidation process.
The size of inclusions was from less than 1 μm and up to
5 μm. Figures 6 and 7 show EDX spot analyses and
mapping of such inclusions, respectively. There were
some pure Al2O3 as well as pure SiO2, but most of them
were combination of both with additional CaO content.
During conventional metallographic sample preparation
due to water presence a part of calcium inclusions might
be lost. If calcium inclusions are of high interest of study
the ion etching (cross-section polisher) or some similar
techniques shall be used.
Stem microstructure consist of pure austenite phase.
No other phases as delta ferrite, sigma phase and chi
phase were observed and detected in the microstructure
even though the specimens were etched or colloidal
silica polished and observed by EBSD. Only in the very
surface region in cross-section mode it was possible to
observe slip lines caused by sand-blasting deformation
(Figure 8). If deformation is enough high the defor-
mation martensite will take place. Such a microstructure
is usually not desirable.
In some cases the formation of stress-induced
martensite might lead to micro-crack formation, which,
especially in combination with intercrystalline corrosion,
leads to material fracture in overload situations 31. Due to
the low carbon content in the examined steel, the
possibilities of intercrystalline corrosion are negligible in
contrast to the stainless steels with a higher carbon
content, where Cr23C6 carbides along the grain boun-
daries are formed. Figure 9 shows grain boundary
between two of austenite grains. The grain boundary was
free of any carbide. EBSD was performed in order to
find out any possible presence of other unwanted phases
in stainless steel grades. However, the technique was
primarily developed for texture analyses, but nowadays it
is used for phase analyses, misorientation measurements
as well as strain and stress detection 32–38. Figure 10
shows results of EBSD measurements. Band contrast
(BC) image sometimes referred as pattern quality image
and can tell us a lot of microstructure details mostly
impossible revealed by chemical etching because of
EBSD being a very surface sensitive techniques. Band
contrast image of analysed sample (Figure 10 (a)) very
precisely revealed all of those non-metallic inclusions
based on aluminium, silicon and calcium oxides39.
Unfortunately, it was almost impossible to obtain
sufficiently good Kikuchi pattern for indexing it.
Therefore, in EBSD map the only iron FCC crystal
lattice was considered. All three inverse pole figure (IPF)
images in X, Y and Z direction were performed.
Analysed steel had no texture and had equiaxed austenite
grains. Because of forging of stem part and further
annealing it is expected to have such microstructure.
4 CONCLUSIONS
Shortly after total hip surgical replacement septic
loosening might occur due to several reasons. If this is
not about to happen the patient has good chances to
undergo recovery successfully. However, after several
years usually due to the process of adhesive and abrasive
wear as well as to wear polishing, polyethylene and
metallic debris occurs in between polyethylene aceta-
bular cup and metallic ball as well as in between
stainless steel stem and cement bond in the case of
cemented prostheses. Based on the investigation
performed on an AISI 316L stainless steel hip prosthesis
it was found that in the present case wear polishing and
adhesive wear played an important role. Acetabular cup
areas exposed to adhesive wear had rough regions where
larger pieces were pulled out while wear polishing of the
polyethylene acetabular cup caused a very smooth,
glossy-like appearance of the surface. Wear polishing of
the femoral stem was observed to occur unevenly over
the stem surface. Certain areas were higher exposed to
wear and are seemed to be more curved part. The wear
M. GODEC: MATERIAL FAILURE OF AN AISI 316L STAINLESS STEEL HIP PROSTHESIS
Materiali in tehnologije / Materials and technology 45 (2011) 2, 85–90 89
polishing of the surface caused a loss of material from
the surface of the implant, the formation of metallic
particles and the possibility of metal-ion formation. The
AISI 316L stainless steel microstructure is austenitic
with typical twin grain boundaries but with no other
usually unwanted phases. Also no carbides at grain
boundaries were found. But an unusually large number
of aluminium-silicon-calcium based non-metallic inclu-
sions were detected. Such high amount of that type of
inclusions is an evidence of wrong desoxidation process
during steel production. However, the number of these
inclusions is too large for the AISI 316L steel grade
especially in an orthopaedic application. Furthermore,
Al2O3 particles due to sand-blasting were found in the
surface. The long-term wear polishing strengthened by
hard Al2O3 particles in the surface and as such might
digs out hard non-metallic inclusions as well and thus
complementing the wear-polishing mechanism with
abrasive wear. Usually the loosening of the components
is triggered by the micromotion of cemented components
which causes the abrasion of the implant’s surface,
producing metal-wear debris and a proliferative soft-
tissue reaction. Based on case study investigation it is no
doubt that such a bad quality stainless steel by huge
amount of hard non-metallic inclusions accelerating the
loosening process much faster due to potentially toxic
and human body irritating debris formed during wear.
Acknowledgments
The author would like to express great acknow-
ledgment to Prof. Dr. V. Antoli~ and Dr. D. Dolinar for
providing surgically replaced total hip prosthesis for
further investigation.
5 REFERENCES
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M. GODEC: MATERIAL FAILURE OF AN AISI 316L STAINLESS STEEL HIP PROSTHESIS
90 Materiali in tehnologije / Materials and technology 45 (2011) 2, 85–90
A. KOCIJAN: A COMPARISON OF THE CORROSION BEHAVIOUR OF AUSTENITIC STAINLESS STEELS ...
A COMPARISON OF THE CORROSION BEHAVIOUR OF
AUSTENITIC STAINLESS STEELS IN ARTIFICIAL
SEAWATER
PRIMERJAVA KOROZIJSKIH LASTNOSTI AVSTENITNIH
NERJAVNIH JEKEL V SIMULIRANI MORSKI VODI
Aleksandra Kocijan
Institute of Metals and Technology, Lepi pot 11, SI-1000 Ljubljana, Slovenia
aleksandra.kocijan@imt.si
Prejem rokopisa – received: 2011-03-15; sprejem za objavo – accepted for publication: 2011-03-25
The evolution of the passive films formed on AISI 317LNM and AISI 316L stainless steels in artificial seawater was studied
using the electrochemical techniques of cyclic voltammetry and potentiodynamic measurements. The extent of the passive range
increased significantly for the AISI 317LNM stainless steel compared to the AISI 316L in the investigated solution. The cyclic
voltammograms of the AISI 317 LMN and AISI 316L were recorded in artificial seawater. The current-density peaks formed
during the cyclic voltammetry were ascribed to the corresponding electrochemical processes taking place on the surface of the
investigated materials, and the influence of the potential scan rate was also studied. The current density of the peaks increased
linearly with the scan rate. The potentials of the anodic and cathodic peaks moved slightly, with the increasing scan rate,
towards more positive and negative values, respectively.
Keywords: stainless steel, sea water, chloride, cyclic voltammetry, potentiodynamic measurements
Z elektrokemijskimi tehnikami cikli~ne voltametrije in potenciodinamskih meritev smo raziskovali tvorbo pasivne plasti na
povr{ini nerjavnih jekel AISI 317LNM in AISI 316L v raztopini simulirane morske vode. [irina pasivnega obmo~ja je pri
nerjavnem jeklu AISI 317 LMN bistveno ve~ja kot pri AISI 316L v preiskovani raztopini. S cikli~no voltametrijo smo opisali
nastale vrhove s procesi, ki potekajo na povr{ini preiskovanih vzorcev. Raziskovali smo tudi vpliv spreminjanja potenciala s
~asom na obliko cikli~nih voltamogramov. Ugotovili smo, da z nara{~ajo~o hitrostjo spreminjanja potenciala s ~asom prihaja do
rahlega nara{~anja in premika vrhov.
Klju~ne besede: nerjavno jeklo, morska voda, kloridi, cikli~na voltametrija, potenciodinamske meritve
1 INTRODUCTION
The austenitic stainless steel AISI 316L is one of the
most commonly used materials. Its mechanical proper-
ties, such as ductility and wear resistance, make it
attractive for particular applications. The corrosion
resistance of stainless steel is relatively good. However,
it is challenged by the hostile environment in marine
applications, as it is susceptible to localised corrosion in
environments containing chloride 1.
The austenitic stainless steel AISI 317LNM is a
nitrogen-alloyed austenitic stainless steel with a high
molybdenum addition. It exhibits an austenitic micro-
structure free of deleterious carbide precipitations at the
grain boundaries. The grade contains some residual
ferrite, up to 2 % after solution annealing at temperatures
1100–1150 °C and water quenching. Its low carbon
content avoids the intergranular corrosion, even on
welded pieces without an ulterior water quenching. The
high molybdenum content gives this steel a higher
resistance to corrosion in chloride-containing environ-
ments than standard grades, i.e., AISI 316L. Nitrogen
additions and a low silicon content have a stabilizing
effect on the austenitic structure, reducing the precipi-
tation of inter-metallic phases during welding. The
nitrogen addition also increases the yield strength com-
pared to AISI 317L. Its main properties are a high
ductility, an easy weldability and a high corrosion
resistance. The main applications are chemical and
petrochemical applications, pollution control equipment
and chemical tankers 2.
Several techniques have been used to study the
corrosion behaviour of stainless steels in different media
1,3–27. The behaviour of AISI 316L stainless steel in
chloride-containing solutions has been studied by
various authors 28–31. Platt et al. 28 compared the corrosion
behaviour of 2205 duplex stainless steel and AISI 316L
in a 0.9 % chloride solution. Electrochemical testing
indicated that the 2205 duplex stainless steel had a wider
passivation range than the AISI 316L. Blanco et al. 29
investigated the corrosion behaviour of two traditional
austenitic stainless steels in reinforced concrete struc-
tures exposed to highly chloride-contaminated atmo-
spheres. The effects of ion nitriding on the corrosion
performance of a 316L stainless steel were evaluated in a
0.9 % sodium chloride solution by Gil et al. 30. It was
shown that an ion nitriding treatment in a 25 % N2–75 %
H2 atmosphere performed at a temperature of 410 °C
improves the surface hardness of the AISI 316L stainless
steel. However, under the experimental conditions
carried out in this research, the nitrided steel is as prone
to localized corrosion as the untreated one. It is consi-
Materiali in tehnologije / Materials and technology 45 (2011) 2, 91–94 91
UDK 669.14.018.8.: 669.193 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 45(2)91(2011)
dered that this behaviour is mainly due to the presence of
CrN, which precipitates during processing, contributing
to the depletion of chromium from the adjacent matrix
and leading to a galvanic corrosion mechanism.
Hryniewicz et al. 31 compared the corrosion behaviour of
austenitic stainless steels after two electropolishing
processes carried out in the absence and presence of a
magnetic field. It has been shown that the proposed
magnetic field process shifts the corrosion potentials in
the direction of greater corrosion resistance.
The influence of artificial seawater on the corrosion
properties of AISI 317LMN stainless steel has not been
investigated extensively. Among the more recent work,
only one paper concerning the corrosion resistance of
AISI 317LMN stainless steel for an application in
pre-stressed concrete structures was published 2.
In the present work AISI 317LMN stainless steel and
AISI 316L stainless steel were studied in artificial
seawater. The study was conducted using the electroche-
mical techniques of cyclic voltammetry and potentio-
dynamic measurements.
2 EXPERIMENTAL
AISI 317 LMN stainless steel and AISI 316L
stainless steel were investigated. Their compositions
were confirmed by analytical chemical methods, as
shown in Table 1. The concentrations of Cr, Ni, Mn and
Mo were determined by using inductively coupled
plasma atomic emission spectroscopy (ICP-AES,
ICP-AES Perkin Elmer Optima 3100 RL instrument), the
concentration of Si was determined gravimetrically
(Mettler H35AR balance). A spectrophotometric method
using the bismuth phosphomolybdate complex was
applied for the determination of phosphorus. A blue
complex formed was extracted with methyl isobutyl
ketone. The absorbance was measured at 625 nm (Opton
PM 6 spectrophotometer). The concentration of carbon
was determined by the oxidation of the sample in an
induction furnace by heating in an oxygen atmosphere to
form CO2, which was then measured with an infrared
detector (Eltra CS-800 instrument). The measuring
principle is based on the infrared-radiation-absorbing
properties of the gases. In the case of the nitrogen
determination the sample was melted at high tempera-
tures up to 3000 °C in an electrically heated graphite
crucible in the furnace. The concentration of nitrogen
was determined with a thermal conductivity detector
(Eltra ON-900 instrument).
The experiments were carried out in artificial
seawater with a composition of 3.5 % NaCl (Merck,
Darmstadt, Germany).
The test specimens were cut into discs of 15 mm
diameter. The specimens were ground with SiC emery
paper down to 1000 grit prior to the electrochemical
studies, and then rinsed with distilled water. The
specimens were then embedded in a Teflon PAR holder
and employed as a working electrode. The reference
electrode was a saturated calomel electrode (SCE, 0.242
V vs. SHE) and the counter electrode was a high-purity
graphite rod.
The cyclic voltammetry and potentiodynamic
measurements were recorded using an EG&G PAR
PC-controlled potentiostat/galvanostat Model 273 with
M252 Softcorr III computer programs. In the case of
potentiodynamic measurements the specimens were
immersed in the solution 1 h prior to the measurement in
order to stabilize the surface at the open-circuit potential.
The potentiodynamic curves were recorded, starting at
250 mV more negative than the open-circuit potential.
The potential was then increased, using a scan rate of 1
mV/s, until the transpassive region was reached.
3 RESULTS AND DISSCUSION
The potentiodynamic behaviour of AISI 317LMN
and AISI 316L stainless steels in artificial seawater is
shown in Figure 1. The differences in the alloys’ com-
positions affected the polarisation and the passivation
behaviours of the tested materials. After 1 h of
stabilization at the open-circuit potential, the corrosion
potential (Ecorr) for the AISI 317LMN in the artificial
seawater was approximately –0.29 V. Following the Tafel
region, the alloy exhibited a broad range of passivation.
The breakdown potential (Eb) for the AISI 317LMN in
artificial seawater was approximately 1.2 V. In the case
of the AISI 316L, the Ecorr in the artificial seawater was
essentially the same as in the case of AISI 317LMN. The
range of passivation was significantly narrowed
compared to the AISI 317LMN specimen and the Eb was
0.20 V. The corrosion-current densities in the passive
range were similar for both the tested specimens.
The cyclic voltammograms of the AISI 317LMN and
AISI 316L recorded in the artificial seawater enabled us
to ascribe the current-density peaks to the corresponding
electrochemical processes taking place on the surface of
the investigated materials, and the influence of the
potential scan rate () was also studied (Figures 2 and
3). The cyclic voltammograms were recorded for the
A. KOCIJAN: A COMPARISON OF THE CORROSION BEHAVIOUR OF AUSTENITIC STAINLESS STEELS ...
92 Materiali in tehnologije / Materials and technology 45 (2011) 2, 91–94
Table 1: The composition of AISI 317LMN and AISI 316L stainless steels in mass fractions, w/%
Tabela 1: Kemijska sestava nerjavnih jekel AISI 317LMN in AISI 316L, masni dele`, w/%
w(Cr) w(Ni) w(Mn) w(Si) w(P) w(S) w(C) w(Mo)
AISI 316L 17.00 10.00 1.40 0.38 0.041 < 0.005 0.021 2.10
AISI 317 LMN 17.33 12.82 1.64 0.39 0.027 < 0.0005 0.014 4.06
AISI 316L and AISI 317LMN at different scan rates, in
the potential ranges from –1.0 V to 0.4 V and 1.2 V,
respectively. In the presence of the artificial seawater
four peaks are observed in the cyclic voltammogram for
the AISI 317 LMN (Figure 2). The first anodic peak A1
at a potential of –0.6 V can be ascribed to the electro-
formation of Fe(II) oxide upon the Cr(III)-containing
passivating layer, existing on the electrode at such
negative potentials according to the reaction: Fe + H2O
 FeO + 2H+ + 2e– 32. The second anodic peak A2 at the
potential of –0.1 V is ascribed to the oxidation of the
Fe(II) species to the Fe(III) species according to the
reaction: 2Fe2+ + 3H2O  Fe2O3 + 6H+ + 2e– 32. This is
followed by the narrow region with a constant current
density, up to 0.5 V, where the current density starts to
increase and the transpassive region is reached. In that
region the transpassive oxidation of the Cr(III) species to
the Cr(VI) species occurs, the Ni(II) species formed
during the passivation process might have been oxidised
to Ni(IV) oxide (NiO2) in this potential range, too 32. In
the reduction cycle in the potential range of peak C2 at
0.2 V the Cr(VI) is reduced to Cr(III) and the iron
oxide-hydroxide layer is largely reduced in the potential
range of the peak C1 at a potential of –0.8 V 20.
In the case of the AISI 316L in artificial seawater
(Figure 3), the anodic peak at a potential of –0.1 V is
observed in the cyclic voltammograms and can be
ascribed to the electro-formation of Fe(II) oxide. After
the passive region the range of the transpassive oxidation
is reached at a potential of 0.35 V. In the case of slower
scan rates, i.e., 5–10 mV/s, a hysteresis is formed in the
reverse cycle, indicating the pronounced dissolution of
the oxide layer formed at higher potentials. In the
reduction cycle both peaks are less evident on that scale;
however, in the potential range of peak C2 at 0.2 V the
Cr(VI) is reduced to Cr(III) and the second peak in the
cathodic cycle at –0.7 V corresponds to the reduction of
the iron oxide-hydroxide layer.
The influence of the scan rate on the oxidation-
reduction processes for both investigated materials in
cyclic voltammograms is also presented in Figures 2
and 3. The current density of the anodic and cathodic
peaks increases linearly with the scan rate. The poten-
tials of the anodic and cathodic peaks move slightly, with
the increasing scan rate, towards more positive and
negative values, respectively.
4 CONCLUSIONS
The corrosion behaviours of the AISI 317LMN and
AISI 316L stainless steels were studied in artificial
seawater. The study was conducted using the electro-
chemical techniques of cyclic voltammetry and potentio-
dynamic measurements.
A. KOCIJAN: A COMPARISON OF THE CORROSION BEHAVIOUR OF AUSTENITIC STAINLESS STEELS ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 91–94 93
Figure 3: Cyclic voltammograms recorded for AISI 316L stainless
steel with increasing scan rate in 3.5 % NaCl
Slika 3: Cikli~ni voltamogrami za nerjavno jeklo AISI 316L v
3,5-odstotni raztopini NaCl z nara{~ajo~o hitrostjo spreminjanja
potenciala s ~asom
Figure 1: Polarisation curves recorded for AISI 317LMN and AISI
316L stainless steel in 3.5 % NaCl
Slika 1: Polarizacijske krivulje za nerjavni jekli AISI 317LMN in
AISI 316L v 3,5-odstotni raztopini NaCl
Figure 2: Cyclic voltammograms recorded for AISI 317LMN
stainless steel with increasing scan rate in 3.5 % NaCl
Slika 2: Cikli~ni voltamogrami za nerjavno jeklo AISI 317LMN v
3,5-odstotni raztopini NaCl z nara{~ajo~o hitrostjo spreminjanja
potenciala s ~asom
The potentiodynamic measurements showed the
superior corrosion stability of AISI 317 LMN stainless
steel in comparison to AISI 316L stainless steel. The
range of passivation was significantly broader for the
AISI 317LMN compared to the AISI 316L specimen.
The corrosion-current densities in the passive range were
similar for both the tested specimens.
The cyclic voltammograms revealed the similar
electrochemical behaviour of both the tested materials in
the presence of the artificial seawater, although in a
different potential range. In the case of the AISI
317LMN stainless steel, four peaks were observed
corresponding to the electro-formation of the Fe(II) and
Fe(III) species, and subsequently to the reduction of
Cr(VI) to Cr(III) oxide and the reduction of the iron
oxide-hydroxide layer. In the case of the AISI 316L
stainless steel the anodic peak was ascribed to the
electro-formation of Fe(II) oxide. In the cathodic cycle
the reduction of Cr(VI) and the reduction of the iron
oxide-hydroxide layer occurs, similar to the results for
the AISI 317LMN stainless steel.
The results of the present study show that the
electrochemical characteristics of AISI 317LMN and
AISI 316L stainless steel are similar. However, the
corrosion resistance of AISI 317LMN stainless steel
under the investigated conditions is superior compared to
AISI 316L, which indicates the possibilities of its
application in a marine environment.
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A. KOCIJAN: A COMPARISON OF THE CORROSION BEHAVIOUR OF AUSTENITIC STAINLESS STEELS ...
94 Materiali in tehnologije / Materials and technology 45 (2011) 2, 91–94
V. KEVORKIJAN et al.: INFLUENCE OF THE FOAMING PRECURSOR’S COMPOSITION AND DENSITY ...
INFLUENCE OF THE FOAMING PRECURSOR’S
COMPOSITION AND DENSITY ON THE FOAMING
EFFICIENCY, MICROSTRUCTURE DEVELOPMENT AND
MECHANICAL PROPERTIES OF ALUMINIUM FOAMS
VPLIV SESTAVE IN GOSTOTE PREKURZORJEV ZA PENJENJE
NA U^INKOVITOST PENJENJA TER RAZVOJ MIKROSTRUKTURE
IN MEHANSKIH LASTNOSTI ALUMINIJSKIH PEN
Varu`an Kevorkijan1, Sre~o Davor [kapin2, Irena Paulin3,4, Borivoj [u{tar{i~3,
Monika Jenko3, Marjana La`eta5
1Zasebni raziskovalec, Betnavska cesta 6, 2000 Maribor, Slovenija
2Institut "Jo`ef Stefan", Odsek za raziskave sodobnih materialov, Jamova 39, 1000 Ljubljana, Slovenija
3In{titut za kovinske materiale in tehnologije, Lepi pot 11, 1000 Ljubljana, Slovenija
4TALUM, d. d., Tovarni{ka cesta 10, 2325 Kidri~evo, Slovenija
5Impol, d. o. o., Partizanska 38, 2310 Slovenska Bistrica, Slovenija
varuzan.kevorkijan@impol.si
Prejem rokopisa – received: 2010-11-08; sprejem za objavo – accepted for publication: 2011-02-11
In this work, the influence of the composition, density and porosity of foaming precursors on the foaming efficiency,
microstructure development and mechanical properties of aluminium foams are presented and discussed. The foams were
prepared, starting from precursors made either by powder metallurgy (PM) or by the melt route. Following the PM route,
precursors were made by mixing Al powder and 3–10 % of volume fractions of dolomite or calcium carbonate particles of
particle size from 20 μm to 120 μm and cold isostatically pressing the mixture at 700 MPa. In the case of the melting route,
precursors were made by introducing dolomite or calcium carbonate particles directly into the molten aluminium at 700 °C.
After melt stirring, the precursors were prepared by casting the semi-solid slurry into a cylindrical, water-cooled mould.
Finally, aluminium foams were made in all cases by inserting precursors into a cylindrical stainless-steel mould and heating the
arrangement at 750 °C for 10 min. After that, the mould was removed from the furnace and the foaming process was stopped by
cooling in air to room temperature.
The microstructure of the obtained foams was investigated by optical and scanning electron microscopy (SEM-EDS), while
XRD was applied for a detailed identification of phases.
The quality of the precursors was evaluated by determining their mechanical properties (uniaxial room-temperature compression
stress-strain curve, compressive strength and energy absorption after a 30 % strain) and the foaming efficiency (the relative
density of the foam obtained). The concentration of the foaming agent and the density of precursors were found to have a
detrimental influence on the foaming efficiency as well as on the foam’s microstructure and mechanical properties. The foaming
of precursors with open porosity were inefficient.
Key words: aluminium foams, foaming agents, calcium carbonate, dolomite, characterization
V delu poro~amo o vplivu sestave ter gostote oz. poroznosti prekurzorja za penjenje na u~inkovitost penjenja ter razvoj
mikrostrukture in mehanskih lastnosti aluminijevih pen. Pene smo izdelovali z uporabo prekurzorjev na osnovi Al s homogeno
porazdeljenimi delci dolomita ali kalcijevega karbonata. Prekurzorje smo pripravljali po postopku pra{ne metalurgije (PM) in z
litjem taline, kar je cenej{e, zagotavlja pa manj homogeno porazdelitev sredstva za penjenje. S PM-postopkom smo prekurzorje
za penjenje izdelovali iz zmesi Al prahu in 3–10 % dolomita ali kalcijevega karbonata razli~ne povpre~ne velikosti (od 20 μm do
120 μm), ki smo jo izostatsko stisnili pri 700 MPa. Z litjem smo prekurzorje pripravljali tako, da smo delce penila uvajali v
Al-talino, segreto do najve~ 700 °C, preme{ali in nastalo suspenzijo ulili v cilindri~en, vodno hlajen jekleni model. Pene smo iz
obeh vrst prekurzorjev izdelovali tako, da smo prekurzor vstavili v zaprt jeklen model za penjenje, segrevali pri 750 °C 10 min.
ter nato ohlajali na zraku do sobne temperature.
Mikrostrukturo nastalih pen smo preu~evali z opti~no in elektronsko (SEM/EDS) mikroskopijo, in sicer tako, da smo ugotavljali
morfologijo in povpre~no velikost por, sestavo vsebovanih faz pa z metodo rentgenske (XRD) difrakcije. Gostoto prekurzorjev
in pen smo dolo~ali na osnovi mase in izra~unane prostornine strojno obdelanih vzorcev, u~inkovitost penjenja (relativno
gostoto dobljene pene) pa na osnovi primerjave dejansko dose`ene gostote pene in gostote aluminija. Primerjalno smo gostoto
pen dolo~ali tudi z Arhimedovo metodo.
Kakovost izdelanih pen smo ocenjevali na osnovi njihovih mehanskih lastnosti (krivulje napetost – deformacija pri sobni
temperaturi, tla~ne trdnosti in sposobnosti absorpcije energije pri 30-odstotni deformaciji). Ugotovili smo, da na u~inkovitost
penjenja ter razvoj mikrostrukture in mehanskih lastnosti vplivata predvsem kemijska sestava in gostota prekurzorja za penjenje,
pri ~emer je bilo prekurzorje z odprto poroznostjo nemogo~e peniti.
Klju~ne besede: aluminijske pene, sredstvo za penjenje, kalcijev karbonat, dolomit, karakterizacija
Materiali in tehnologije / Materials and technology 45 (2011) 2, 95–103 95
UDK 549.74:621.762:620.18 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 45(2)95(2011)
1 INTRODUCTION
Closed-cell-foamed aluminium is a macro-composite
material consisting of an aluminium alloy matrix, usually
discontinuously reinforced with various ceramic parti-
culates and closed pores filled with gas distributed
throughout the matrix. This unique structure possesses
an unusual combination of properties, such as a low
density, a high weight-specific stiffness, extraordinary
energy absorption and a remarkable vibration attenu-
ation, which are important for a number of engineering
applications. Aluminium foams are also non-flammable,
ecologically harmless and easily recyclable.
Methods to produce closed-cell aluminium foams
have been known for a number of years and can be
generally separated into two fundamental groups: (i)
foaming liquid metal (casting procedures) and (ii)
foaming metallic precursors (powder metallurgy-PM
procedures)1.
Titanium hydride (TiH2) was mainly applied as a
blowing agent for both the casting and powder
metallurgical procedures of foaming for aluminium and
aluminium alloys. The role of the blowing agent in both
– the casting and powder metallurgical procedures of
foaming of aluminium – is to release gas, since the metal
is transferred in the liquid or the semi-liquid viscous
state. However, as discussed in detail by Gergely et al.2,
several factors are involved in the selection of a blowing
agent for the foaming of aluminium alloys. These
include: (i) the kinetics and thermodynamic characte-
ristics of decomposition reactions and reactions between
the blowing agent particles and the molten or semi-
molten aluminium alloy; (ii) the wetting of foaming-
agent particles with molten metal; (iii) the influence of
decomposition products and foaming gas on the
stabilisation of foams; and (iv) the availability, cost and
user-friendly handling of the powdered agent concerned.
As a foaming agent for the production of aluminium
alloy foams by the casting and powder metallurgical
routes, TiH2 has several limitations that significantly
influence the mass application of Al foams. The main
limitations of the foaming technologies using TiH2 as a
foaming agent are the following:
• TiH2 is very expensive.
• The decomposition temperature of TiH2 is very low –
starting at about 400 °C for the untreated hydride.
• TiH2 particles act only as a blowing agent and are not
involved in the foam stability.
• The density of TiH2 (approx. 3.9 g/cm3) is signi-
ficantly higher than the density of molten aluminium
(approx. 2.7 g/cm3). Thus, during foaming the
settling of TiH2 particles occurs by gravity, resulting
in a non-uniform microstructure of the foam.
The cost-effective and highly promising alternatives
to the TiH2 blowing agent are CaCO3, as a marble
powder or synthetic calcium carbonate, and dolomite
powder. In contrast to TiH2, calcium carbonate and
dolomite are of low cost, significantly lower than the
cost of aluminium, and with a density (2.71 g/cm3 to
2.84 g/cm3) almost identical to the density of molten
aluminium. Moreover, the CaCO3 and dolomite decom-
position temperatures are above the melting point of
aluminium, usually in the temperature interval between
660 °C and 930 °C. Therefore, CaCO3 and dolomite are
particularly suitable for the melt-route, settling-free
production of foamed aluminium-based materials.
The "Foamcarp" process of the indirect foaming of
aluminium using CaCO3 particles previously introduced
into a solidified precursor and, more recently, an addi-
tional indirect foaming procedure of molten aluminium
with a CaCO3-containing preform isostatically pressed
from a mixture of CaCO3 and AA6061 powders were
reported2,3.
In some early patents4 and in the recent work of
Nakamura et al.5, the use of CaCO3 was found to be
potentially suitable as a foaming agent for direct
(melt-route) foam manufacturing. Moreover, Bryat et al.6
reported that CaCO3 acts in contact with the molten
aluminium as a foaming agent and, at the same time,
through the decomposition products has a significant
effect on the foam stabilisation, enabling the formation
of "self-stabilized aluminium foams" (i.e., foams solidi-
fied from a foamable suspension created through the
controlled decomposition of carbonate powders with
molten aluminium). A similar cell face stabilising
mechanism operating in carbonate-foamed melts was
also reported by Gergely et al.2.
The additional advantage of CaCO3 in comparison
with a TiH2 blowing agent of the same average particle
size is in achieving foams with higher porosity levels and
finer cell sizes5. Gergely et al.2 showed that by replacing
TiH2 by CaCO3 as a foaming agent, foams can be
produced having appreciably finer cells and more uni-
form cell structures at a significantly reduced raw-mate-
rial cost.
On the other hand, Alcoa patented7 a method of
making an aluminum foam product by adding reactive
gas producing particles to a molten aluminum alloy at a
temperature that is above the decomposition temperature
of the reactive-gas producing particles. The reactive-gas
producing particles are selected from the group con-
sisting of magnesium carbonate, calcium carbonate,
dolomite and mixtures thereof.
Although the applicability of CaCO3 and dolomite
powders as a foaming agent in the direct and indirect
foaming of aluminium alloys and composites has already
been investigated, there is a need for an additional study
of the influence of the CaCO3 and dolomite powder
morphology (particularly the average particle size and
particle size distribution) and the role of CaCO3 and
dolomite decomposition products on the achievement of
a better foam stability and a higher foam quality.
Hence, in this paper, the performance of synthetic
calcium carbonate and natural dolomite powders as
V. KEVORKIJAN et al.: INFLUENCE OF THE FOAMING PRECURSOR’S COMPOSITION AND DENSITY ...
96 Materiali in tehnologije / Materials and technology 45 (2011) 2, 95–103
cost-effective foaming agents was investigated by both
the powder metallurgical and the melt route of foam
preparation. The influence of CaCO3 and dolomite
particle morphology and volume fraction on the foaming
behaviour and the development of the foam micro-
structure were also monitored. As a result, a clearer
interrelation between the foam morphology and struc-
ture, on the one hand, and its mechanical properties, on
the other, was established.
2 EXPERIMENTAL
All the foams made in this work were prepared by
indirect foaming methods starting from a solid foamable
precursor consisting of a metallic matrix containing
uniformly dispersed blowing-agent particles. Foamable
precursors were made either by: (i) the powder metallur-
gical route, or (ii) the melt route using the same blowing
agent, i.e., CaCO3 powders (types C-A, C-B and C-C)
with various average particle sizes ((38, 72 and 120) μm,
respectively) and dolomite powders (type D-A, D-B and
D-C) with various average particle sizes ((44, 76 and 97)
μm, respectively).
Following the powder metallurgical (PM) route,
foamable precursors with CaCO3 particles were made by
mixing Al powder with an average particle size of 63 μm
(purity (mass fractions, w: 99.7 %, oxygen content: 0.25 %)
and (3, 5, 7 and 12) % of the blowing agent, followed by
cold compaction in a lubricated 20-mm diameter die to a
pressure of 600 MPa to 900 MPa.
In the case of the melt route, foamable precursors
with the same concentration of CaCO3 blowing agent
((3, 5, 7 and 12) %) and the same geometry were
prepared by an induction-heated, batch-type, stir-casting
method by which the aluminium powder (the same as
used for the PM route) was induction melted, followed
by the addition of CaCO3 particles, stirring and casting.
Once the molten aluminium had reached 750 °C, the
power was switched off and the melt stirring was
initiated until the temperature of the melt decreased to
700 °C. After that, the blowing agent/aluminium powder
mixture (1 : 2 mass ratio) was introduced and the melt
was stirred (at approximately 1200 r/min) for an additi-
onal 30–90 s. Finally, the foamable precursors were
prepared by casting the semi-solid slurry into a room-
temperature mould with a diameter of 20 mm.
In the case of the powder metallurgy (P/M) route, the
fabrication of foamable precursors with dolomite parti-
cles was conducted by mixing Al powder with an
average particle size of 63 μm (purity: 99.7 %, oxygen
content: 0.25 %), 5 % of SiC particles with an average
particle size of 10 μm and (3, 5, 7 and 12) % of blowing
agent, followed by cold compaction, in a lubricated
20-mm diameter die to a pressure of 600 MPa to 900
MPa.
Following the melt route, foamable precursors with
the same concentration of dolomite blowing agent ((3, 5,
7 and 12) %) and the same geometry were prepared by
induction-heated, batch-type, stir-casting in which
aluminium powder (the same as used for the P/M route)
was induction melted, followed by the addition of
dolomite particles, stirring and casting. Once the molten
aluminium was heated to 700 °C, the power was
switched off and melt stirring was initiated until the
temperature of the melt decreased to 685 °C. After that,
the blowing agent/aluminium powder mixture (1 : 2 mass
ratio) was introduced and the melt was stirred (at appro-
ximately 1200 r/min) for an additional 30–90 s. Finally,
the foamable precursors were prepared by casting the
semi-solid slurry into a room-temperature mould with
20-mm diameter.
The solidified precursors were machined and some of
the samples were additionally cold isostatically pressed.
The density of the foamable precursors as well as of the
foams obtained was calculated from the mass and
geometry of the samples and, in addition, measured by
Archimedes’ method. The distribution of the blowing-
agent particles inside the Al matrix was examined by an
assessment of the optical and scanning electron
micrographs of as-polished bars.
All the precursors were foamed in a conventional
batch furnace with air atmosphere circulation under the
same experimental conditions (temperature, time,
cooling method). Before foaming, the individual
precursors were inserted into a cylindrical (40-mm
diameter, 70-mm long) stainless-steel mould coated with
a boron nitride suspension. The mould dimensions and
the precursor size (20-mm diameter and 60-mm long)
were selected to allow the expansion of the precursor to a
foam with a theoretical density close to 0.6 g/cm3. The
arrangement was placed inside a pre-heated batch
furnace at 750 °C for 10 min. After that period of time,
the mould was removed from the furnace and the
foaming process was stopped by rapid cooling with
pressurised air to room temperature. The thermal history
of the foam sample was recorded, using a thermocouple
located directly in the precursor material.
The porosity of the foam was calculated using
equation: 1 – (foam density/ aluminium density). Macro
and microstructural examinations were performed on
sections obtained by precision wire-cutting across the
samples and on samples mounted in epoxy resin using
optical and scanning electron microscopy (SEM-EDS).
The average size of the pores in the foams was
estimated by an assessment of the optical and scanning
electron micrographs of as-polished foam bars using the
point-counting method and image-analysis and pro-
cessing software.
Regarding the mechanical properties of the foams,
uniaxial room-temperature compressive tests were
carried out on a Zwick 1474 testing machine at a con-
stant 5-mm/min crosshead displacement. Testing was
performed on standard prismatic foam specimens of 50
mm × 12 mm × 17 mm and each point of the stress-strain
V. KEVORKIJAN et al.: INFLUENCE OF THE FOAMING PRECURSOR’S COMPOSITION AND DENSITY ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 95–103 97
curve was determined as an average of four individual
measurements. Compression was stopped whenever
either 80 % strain or 95 kN force (equivalent to 61.9
MPa) was reached. As a result of testing, the uniaxial
compression stress-strain curve, compressive strength
and energy absorption after a 30 % strain were
determined and correlated with the density, the average
pore size and microstructure of the foam samples.
3 RESULTS AND DISCUSSION
3.1 Characterisation of the foamable precursors
The measured and calculated densities of the
foamable precursors obtained by the PM route, Table
1a,b, confirmed that under isostatic pressing with an
applied pressure of 700 MPa the precursors prepared by
PM possessed a closed porosity and densities above 98 %
of theoretical, whereas as-machined precursors obtained
by the melt route (Table 2a,b) had a significant fraction
of open porosity and thus were not suitable for foaming
to the desired foam densities (usually about 0.5 g/cm3 to
0.7 g/cm3). However, after additional isostatic pressing,
the porosity in these precursors was successfully reduced
to below 2.0 % (volume fraction /%) (Table 3 a,b) and
foam samples with densities between 0.62 g/cm3 and
0.80 g/cm3 were obtained.
It is important to note that very high precursor
densities (>99 % of theoretical) were achieved only in
the precursors prepared by the powder metallurgical
route with 3–7 % of CaCO3 and dolomite particles of
Type-C-A or Type D-A (Table 1 a,b). With a higher
particle content, and by using coarser CaCO3 or dolomite
powders of Type-C-B or Type-C-C as well as Type D-B
and Type D-C, this could not be achieved and resulted in
a lower foaming efficiency, as evident in Tables 4 a,b–6
a,b.
The foaming efficiency of the precursors was
evaluated from the relative density of the foam obtained,
, calculated by dividing the apparent density of the
foam, F, by the density of aluminium, Al. Thus, the
foaming efficiency is expressed as:
 = 1 –  = 1 – (F/A) (1)
which actually corresponds to the volume fraction of
pores in the foam samples. The lower the foam density,
the higher is the foaming efficiency.
In all cases the experimental results clearly indicate
that the porosity measured in foamable precursors and
the apparent density achieved in aluminium foam
samples are inversely proportional. Generally, foamable
precursors with a lower porosity resulted in foam
samples with a higher apparent density and a lower
foaming efficiency.
Under the same foaming conditions (temperature,
time), the average pore size of the foam samples was
influenced by the density of foaming precursors and the
initial size of the foaming particles. As a rule, in foams
made from precursors with high density (99 % of
theoretical), the average pore size remained below 1.0
mm. On the other hand, in foams made from precursors
with a lower density (below 99 % of theoretical), the
pores grew to a 20 % to 50 % higher average size pore.
Regarding the initial size of the foaming particles,
which also influences the density of the precursor and,
hence, the density of the foam samples, the increase of
the average particle size of CaCO3 or dolomite foaming
agent was observed to have a detrimental influence on
the average size of the pores. Coarser CaCO3 and
dolomite powders led to the formation of larger bubbles
in the foam structure.
Table 1 a: Porosity of CaCO3 particles containing foamable pre-
cursors obtained by the PM route
Tabela 1 a: Poroznost prekurzorjev s CaCO3-penilom, izdelanih s po-
stopkom pra{ne metalurgije
Chemical composition of
precursors (w/%)
Porosity
(/%)
CaCO3 Al powder Calculated Measured
Type C-A
3 97 0.8 ± 0.08 0.7 ± 0.04
5 95 0.8 ± 0.08 0.8 ± 0.04
7 93 0.9 ± 0.09 1.0 ± 0.05
10 90 1.1 ± 0.11 1.2 ± 0.06
Type C-B
3 97 1.1 ± 0.10 1.0 ± 0.05
5 95 1.1 ± 0.11 1.1 ± 0.06
7 93 1.3 ± 0.13 1.3 ± 0.07
10 90 1.7 ± 0.17 1.8 ± 0.09
Type C-C
3 97 1.3 ± 0.11 1.4 ± 0.07
5 95 1.5 ± 0.15 1.4 ± 0.07
7 93 1.6 ± 0.16 1.7 ± 0.09
10 90 1.8 ± 0.18 2.0 ± 0.10
Table 1 b: Porosity of dolomite particles containing foamable pre-
cursors obtained by the PM route.
Tabela 1 b: Poroznost prekurzorjev z dolomitom kot penilom, izde-
lanih po postopku pra{ne metalurgije
Chemical composition of performs
w/%
Porosity
/%
Dolomite SiC Al powder Calculated Measured
Type D-A
3 5 92 0.7 ± 0.07 0.7 ± 0.04
5 5 90 0.8 ± 0.08 0.8 ± 0.04
7 5 88 0.9 ± 0.09 1.0 ± 0.05
10 5 85 1.2 ± 0.12 1.3 ± 0.07
Type D-B
3 5 92 1.0 ± 0.10 1.0 ± 0.05
5 5 90 1.0 ± 0.10 1.1 ± 0.06
7 5 88 1.2 ± 0.12 1.3 ± 0.07
10 5 85 1.6 ± 0.16 1.8 ± 0.09
Type D-C
3 5 92 1.1 ± 0.11 1.1 ± 0.06
5 5 90 1.2 ± 0.12 1.3 ± 0.06
7 5 88 1.3 ± 0.13 1.4 ± 0.07
10 5 85 1.8 ± 0.18 2.0 ± 0.10
V. KEVORKIJAN et al.: INFLUENCE OF THE FOAMING PRECURSOR’S COMPOSITION AND DENSITY ...
98 Materiali in tehnologije / Materials and technology 45 (2011) 2, 95–103
Table 2a: Porosity of as-machined CaCO3 containing foamable pre-
cursors prepared by the melt route
Tabela 2a: Poroznost strojno obdelanih prekurzorjev s CaCO3 peni-
lom, izdelanih po livarskem postopku
Chemical composition of
precursors (w/%)
Porosity
(/%)
CaCO3 Al powder Calculated Measured
Type C-A
3 97 4.2 ± 0.42 4.1 ± 0.21
5 95 4.6 ± 0.46 4.4 ± 0.22
7 93 4.9 ± 0.49 4.8 ± 0.24
10 90 5.3 ± 0.53 5.2 ± 0.26
Type C-B
3 97 4.1 ± 0.41 4.2 ±0.21
5 95 4.7 ± 0.47 4.7 ± 0.24
7 93 5.0 ± 0.50 5.1 ± 0.26
10 90 5.4 ± 0.54 5.5 ± 0.28
Type C-C
3 97 3.9 ± 0.39 4.1 ± 0.21
5 95 4.2 ± 0.42 4.2 ± 0.21
7 93 4.3 ± 0.43 4.4 ± 0.22
10 90 4.8 ± 0.48 4.9 ± 0.25
Table 2b: Porosity of as-machined dolomite particles containing
foamable performs prepared by the melt route.
Tabela 2b: Poroznost strojno obdelanih prekurzorjev z dolomitom kot
penilom, izdelanih po livarskem postopku.
Chemical composition of performs
w/%
Porosity
/%
Dolomite SiC Al powder Calculated Measured
Type D-A
3 5 92 4.7 ± 0.47 4.9 ± 0.25
5 5 90 4.9 ± 0.49 5.1 ± 0.26
7 5 88 5.4 ± 0.54 5.8 ± 0.29
10 5 85 6.1 ± 0.61 6.6 ± 0.33
Type D-B
3 5 92 4.5 ± 0.45 4.7 ± 0.24
5 5 90 4.7 ± 0.47 5.0 ± 0.25
7 5 88 5.0 ± 0.50 5.3 ± 0.27
10 5 85 5.7 ± 0.57 6.2 ± 0.31
Type D-C
3 5 92 4.4 ± 0.44 4.5 ± 0.23
5 5 90 4.5 ± 0.45 4.7 ± 0.24
7 5 88 4.7 ± 0.47 5.0 ± 0.25
10 5 85 4.9 ± 0.49 5.3 ± 0.27
Table 3a: Porosity of CaCO3 containing foamable precursors obtained
by the melt route improved by additional isostatic pressing
Tabela 3a: Poroznost prekurzorjev s CaCO3-penilom, izdelanih po
livarskem postopku, izbolj{ana z dodatnim hladnim izostatskim
stiskanjem
Chemical composition of
precursors (w/%)
Porosity
(/%)
CaCO3 Al powder Calculated Measured
Type C-A
3 97 1.3 ± 0.13 1.3 ± 0.06
5 95 1.6 ± 0.16 1.7 ± 0.09
7 93 1.9 ± 0.19 2.0 ± 0.10
10 90 2.0 ± 0.20 2.2 ± 0.11
Type C-B
3 97 1.1 ± 0.11 1.2 ± 0.06
5 95 1.2 ± 0.12 1.2 ± 0.06
7 93 1.4 ± 0.14 1.5 ± 0.08
10 90 1.7 ± 0.17 1.9 ± 0.09
Type C-C
3 97 1.0 ± 0.10 1.0 ± 0.10
5 95 1.2 ± 0.12 1.1 ± 0.06
7 93 1.3 ± 0.13 1.4 ± 0.07
10 90 1.6 ± 0.16 1.8 ± 0.09
Table 3b: Porosity of dolomite particles containing foamable
performs obtained by the melt route improved by additional isostatic
pressing
Tabela 3b: Poroznost prekurzorjev z dolomitom kot penilom,
izdelanih po livarskem postopku, izbolj{ana z dodatnim hladnim
izostatskim stiskanjem
Chemical composition of performs
w/%
Porosity
/%
Dolomite SiC Al powder Calculated Measured
Type D-A
3 5 92 0.9 ± 0.09 0.9 ± 0.05
5 5 90 0.9 ± 0.09 0.9 ± 0.05
7 5 88 1.0 ± 0.10 1.1 ± 0.06
10 5 85 1.1 ± 0.11 1.2 ± 0.06
Type D-B
3 5 92 0.9 ± 0.09 0.9 ± 0.05
5 5 90 0.9 ± 0.09 0.9 ± 0.05
7 5 88 0.9 ± 0.09 1.0 ± 0.05
10 5 85 1.0 ± 0.10 1.1 ± 0.06
Type D-C
3 5 92 0.8 ± 0.08 0.8 ± 0.04
5 5 90 0.8 ± 0.08 0.8 ± 0.04
7 5 88 0.8 ± 0.08 0.9 ± 0.05
10 5 85 0.9 ± 0.09 1.0 ± 0.05
V. KEVORKIJAN et al.: INFLUENCE OF THE FOAMING PRECURSOR’S COMPOSITION AND DENSITY ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 95–103 99
Table 4a: Density, foaming efficiency and average pore size of
aluminium foams prepared by the PM route using CaCO3 foaming
agent
Tabela 4a: Gostota, u~inkovitost penjenja in povpre~na velikost por v
aluminijskih penah, izdelanih po postopku pra{ne metalurgije z
uporabo CaCO3 kot sredstva za penjenje
Initial composition of
foamable precursors
(w/%)
Selected properties of foamed
samples
CaCO3 Al powder Density(g/cm3)
Foaming
efficiency
(%)
Average
pore size
(mm)
Type D-A
3 97 0.42 ± 0.02 84.4 0.8 ± 0.08
5 95 0.47 ± 0.03 82.6 0.9 ± 0.09
7 93 0.51 ± 0.03 81.1 1.0 ± 0.10
10 90 0.55 ± 0.03 79.6 1.2 ± 0.12
Type D-B
3 97 0.46 ± 0.03 83.0 0.5 ± 0.05
5 95 0.49 ± 0.03 81.9 0.7 ± 0.07
7 93 0.53 ± 0.03 80.4 0.8 ± 0.08
10 90 0.59 ± 0.03 78.1 0.9 ± 0.09
Type D-C
3 97 0.53 ± 0.03 80.4 0.5 ± 0.05
5 95 0.55 ± 0.03 79.6 0.5 ± 0.05
7 93 0.59 ± 0.03 78.1 0.6 ± 0.06
10 90 0.61 ± 0.03 77.4 0.8 ± 0.08
Table 4b: Density, foaming efficiency and the average pore size of
aluminium foams prepared by the PM route using dolomite as
foaming agent
Tabela 4b: Gostota, u~inkovitost penjenja in povpre~na velikost por v
aluminijskih penah, izdelanih po postopku pra{ne metalurgije z
uporabo dolomita kot sredstva za penjenje
Initial composition of
foamable performs
w/%
Selected properties of foamed
samples
Dolo-
mite SiC
Al
powder
Density
(g/cm3)
Foaming
efficiency
(%)
Average
pore size
(mm)
Type D-A
3 5 92 0.56 ± 0.03 79,3 0.9 ± 0.09
5 5 90 0.59 ± 0.03 78,1 0.9 ± 0.09
7 5 88 0.63 ± 0.03 76,7 1.1 ± 0.11
10 5 85 0.69 ± 0.03 74,4 1.3 ± 0.13
Type D-B
3 5 92 0.51 ± 0.03 81,1 0.6 ± 0.06
5 5 90 0.53 ± 0.03 80,4 0.7 ± 0.07
7 5 88 0.57 ± 0.03 78,9 0.9 ± 0.09
10 5 85 0.59 ± 0.03 78,1 0.9 ± 0.09
Type D-C
3 5 92 0.50 ± 0.03 81,5 0.6 ± 0.06
5 5 90 0.52 ± 0.03 80,7 0.6 ± 0.06
7 5 88 0.55 ± 0.03 79,6 0.8 ± 0.08
10 5 85 0.56 ± 0.03 79,3 0.9 ± 0.09
Table 5a: Density, foaming efficiency and average pore size of
aluminium foams prepared from as-machined foamable preforms
fabricated by the melt route using CaCO3 foaming agent
Tabela 5a: Gostota, u~inkovitost penjenja in povpre~na velikost por v
aluminijskih penah, izdelanih iz strojno obdelanih prekurzorjev s
CaCO3 kot sredstvom za penjenje, dobljenih po livarskem postopku
Initial composition of
foamable precursors
(w/%)
Selected properties of foamed
samples
CaCO3 Al powder Density(g/cm3)
Foaming
efficiency
(%)
Average
pore size
(mm)
Type C-A
3 97 0.89 ± 0.05 67.0 1.1 ± 0.11
5 95 0.92 ± 0.05 65.0 1.3 ± 0.13
7 93 0.97 ± 0.05 64.1 1.4 ± 0.14
10 90 0.99 ± 0.05 63.3 1.4 ± 0.14
Type C-B
3 97 0.84 ± 0.03 68.9 0.8 ± 0.08
5 95 0.89 ± 0.03 67.0 0.9 ± 0.09
7 93 0.93 ± 0.04 65.6 0.9 ± 0.09
10 90 0.95 ± 0.04 64.8 1.1 ± 0.11
Type C-C
3 97 0.79 ± 0.04 70.7 0.7 ± 0.07
5 95 0.81 ± 0.04 70.0 0.8 ± 0.08
7 93 0.85 ± 0.04 68.5 1.2 ± 0.12
10 90 0.88 ± 0.04 67.4 1.5 ± 0.15
Table 5b: Density, foaming efficiency and average pore size of
aluminium foams prepared from as-machined foamable performs
fabricated by the melt route using dolomite as foaming agent
Tabela 5b: Gostota, u~inkovitost penjenja in povpre~na velikost por v
aluminijskih penah, izdelanih iz strojno obdelanih prekurzorjev z
dolomitom kot sredstvom za penjenje, dobljenih po livarskem
postopku
Initial composition of
foamable performs
w/%
Selected properties of foamed
samples
Dolo-
mite SiC
Al
powder
Density
(g/cm3)
Foaming
efficiency
(%)
Average
pore size
(mm)
Type D-A
3 5 92 0.71 ± 0.04 73,7 1.2 ± 0.12
5 5 90 0.72 ± 0.04 73,3 1.4 ± 0.14
7 5 88 0.75 ± 0.04 72,2 1.4 ± 0.14
10 5 85 0.81 ± 0.04 70,0 1.5 ± 0.15
Type D-B
3 5 92 0.61 ± 0.03 77,4 0.8 ± 0.08
5 5 90 0.63 ± 0.03 76,6 0.9 ± 0.09
7 5 88 0.66 ± 0.03 75,5 1.0 ± 0.10
10 5 85 0.70 ± 0.04 74,1 1.1 ± 0.11
Type D-C
3 5 92 0.65 ± 0.03 75,9 0.8 ± 0.08
5 5 90 0.66 ± 0.03 75,5 0.9 ± 0.09
7 5 88 0.68 ± 0.03 74,8 1.1 ± 0.11
10 5 85 0.72 ± 0.04 73,3 1.5 ± 0.15
V. KEVORKIJAN et al.: INFLUENCE OF THE FOAMING PRECURSOR’S COMPOSITION AND DENSITY ...
100 Materiali in tehnologije / Materials and technology 45 (2011) 2, 95–103
Table 6a: Density, foaming efficiency and average pore size of
aluminium foams prepared by the melt route from as-machined and
additionally isostatically pressed foamable precursors with CaCO3
foaming agent
Tabela 6a: Gostota, u~inkovitost penjenja in povpre~na velikost por v
aluminijskih penah, izdelanih iz strojno obdelanih in hladno izostatsko
stisnjenih prekurzorjev s CaCO3 kot sredstvom za penjenje, dobljenih
po livarskem postopku
Initial composition of
foamable precursors
(w/%)
Selected properties of foamed
samples
CaCO3 Al powder Density(g/cm3)
Foaming
efficiency
(%)
Average
pore size
(mm)
Type C-A
3 97 0.69 ± 0.03 74.4 1.1 ± 0.11
5 95 0.72 ± 0.04 73.3 1.2 ± 0.12
7 93 0.76 ± 0.04 71.9 1.3 ± 0.13
10 90 0.80 ± 0.04 70.4 1.6 ± 0.16
Type C-B
3 97 0.64 ± 0.03 76.3 0.8 ± 0.08
5 95 0.69 ± 0.03 74.4 0.9 ± 0.09
7 93 0.72 ± 0.04 73.3 1.3 ± 0.13
10 90 0.74 ± 0.04 72.6 1.4 ± 0.14
Type C-C
3 97 0.62 ± 0.03 77.0 0.9 ±0.09
5 95 0.67 ± 0.03 75.2 1.1 ± 0.11
7 93 0.71 ± 0.04 73.7 1.3 ± 0.13
10 90 0.73 ± 0.04 73.0 1.6 ± 0.16
Table 6b: Density, foaming efficiency and average pore size of
aluminium foams prepared by the melt route from as-machined and
additionally isostatically pressed foamable performs using dolomite as
foaming agent
Tabela 6b: Gostota, u~inkovitost penjenja in povpre~na velikost por v
aluminijskih penah, izdelanih iz strojno obdelanih in hladno izostatsko
stisnjenih prekurzorjev z dolomitom kot sredstvom za penjenje,
dobljenih po livarskem postopku
Initial composition of
foamable performs
w/%
Selected properties of foamed
samples
Dolo-
mite SiC
Al
powder
Density
(g/cm3)
Foaming
efficiency
(%)
Average
pore size
(mm)
Type D-A
3 5 92 0.63 ± 0.03 76,7 1.1 ± 0.11
5 5 90 0.67 ± 0.03 75,2 1.2 ± 0.12
7 5 88 0.71 ± 0.03 73,7 1.3 ± 0.13
10 5 85 0.78 ± 0.03 71,1 1.6 ± 0.16
Type D-B
3 5 92 0.57 ± 0.03 78,9 0.7 ± 0.07
5 5 90 0.59 ± 0.03 78,1 0.9 ± 0.09
7 5 88 0.62 ± 0.03 77,0 1.1 ± 0.11
10 5 85 0.63 ± 0.03 76,7 1.4 ± 0.14
Type D-C
3 5 92 0.58 ± 0.03 78,5 0.7 ± 0.07
5 5 90 0.60 ± 0.03 77,8 0.8 ± 0.08
7 5 88 0.64 ± 0.03 76,3 1.0 ± 0.10
10 5 85 0.67 ± 0.03 75,2 1.1 ± 0.11
3.2 Microstructural investigation of aluminium foam
samples
A similar cellular structural development with
spherical, closed pores was obtained by both the powder
metallurgy (Figure 1) and the melt processing route
(Figure 2). However, as is evident in Figure 1, the
samples obtained by the powder metallurgical route had
a more uniform microstructure consisting of well-sepa-
rated individual cells. On the other hand, the micro-
structure of the samples obtained by the melt processing
route, Figure 2, revealed the presence of some indi-
vidual, non-uniformities created by flow (the movement
of bubbles with respect to each other), drainage (flow of
liquid metal through the intersection of three foam
V. KEVORKIJAN et al.: INFLUENCE OF THE FOAMING PRECURSOR’S COMPOSITION AND DENSITY ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 95–103 101
Figure 2: Cross-section of an aluminium foam obtained by the melt
route with a characteristic channel network and foam drainage
Slika 2: Posnetek pre~nega prereza vzorca aluminijske pene, izdela-
nega po livarskem postopku, na katerem so poleg posameznih por
opazne tudi nehomogenosti, povzro~ene z zlitjem por in odtekanjem
taline skozi zna~ilne kanale v mikrostrukturi
Figure 1: Cross-section of an aluminium foam obtained by the
powder metallurgical route with well-separated individual cells and
relatively uniform microstructure.
Slika 1: Posnetek pre~nega prereza vzorca aluminijske pene, izdela-
nega po postopku pra{ne metalurgije, z izra`enimi posameznimi
porami in relativno enakomerno mikrostrukturo
films), coalescence (sudden instability in foam film) and
coarsening (slow diffusion of gas from smaller bubbles
to larger ones).
The absence of considerable pore coarsening and
drainage suggests that there is a cell-face stabilising
mechanism operating in the carbonate-foamed melts2,
slowing down the cell-face rupturing process and hence,
inhibiting cell coarsening. The mechanism is likely to be
a result of the foaming gas (CO2)/melt or semi-solid
slurry reaction during the foaming procedure, as was
discussed in detail by Gergely et al.2.
Concerning the average pore size and the uniformity
in cell size distribution, foams made by the powder
metallurgical route have finer pores and a more regular
morphology than samples made by the melt route, par-
ticularly those from as-machined precursors. However,
an additional cold isostatic pressing of the as-machined
precursor obtained by the melt route was found to help in
achieving more uniform foams with a smaller average
pore size, similar to those obtained by the powder
metallurgical route. The improvement is most probably
caused by better compacting of the individual CaCO3 or
dolomite particles and the aluminium matrix, resulting in
a higher density of the foamable precursor.
3.3 Mechanical properties
Figure 3 shows an example of the stress-strain
response of samples foamed from preforms prepared by
the PM route in which the compressive strength of the
foams was correlated with their density.
Because of the closed cell structure, the compressive
foam behaviour in all cases showed a typical stress-strain
diagram with a division into three parts: a linear increase
in stress mainly caused by elastic deformation, followed
by a plateau caused by homogeneous plastic deformation
and a final steep increase due to the collapse of the cells.
The compressive strength was taken as the initial peak
stress. Foams made by the PM route possessed the
highest compressive strength, while samples foamed
from as-machined precursors had significantly lower
values. For the interval of foam densities analysed in this
work (from 0.42g/cm3 to 0.55 g/cm3), it was found that
in more dense foam samples the position of the plateau
shifted toward higher stress values.
The energy absorbed per unit volume (E-energy
absorption capacity), which is one of the most important
characteristics of aluminium foams, was determined
from the area under the stress-strain plots as follows8:
E
l
= ∫   
0
( )d (2)
Where  is the compressive stress, l is the limit of the
strain concerned and  is the compressive strain. The
calculated values of energy-absorption capacity for
samples are plotted in Figure 2 and correlated with the
foam density. The typical response was found to be a
quasi-Gaussian function with a maximum energy-
absorption capacity in a very narrow density range.
The maximum energy-absorption capacity for various
foams is summarized in Figure 3. For foams made by
the PM route, the maximum energy-absorption capacity
of 6.14 MJ/m3 was achieved in foams with a density of
0.53 g/cm3. On the other hand, in samples foamed from
as-machined precursors fabricated by melt route, a
maximum energy-absorption capacity of only 5.41
MJ/m3 was found. The maximum appeared at a foam
density of 0.58 g/cm3. Finally, in the melt route fabri-
cated precursors, additionally isostatically pressed before
foaming, an intermediate maximum energy-absorption
capacity of 5.82 MJ/m3 was found in samples with a
density of 0.58 g/cm3.
V. KEVORKIJAN et al.: INFLUENCE OF THE FOAMING PRECURSOR’S COMPOSITION AND DENSITY ...
102 Materiali in tehnologije / Materials and technology 45 (2011) 2, 95–103
Figure 4: Example of optimization of the aluminium foam density
range for the maximum energy absorption capacity: A)-foams
obtained by powder metallurgy, B)-foams obtained from as-machined
precursors fabricated by the melt route, and C)-foams obtained from
as-machined and cold isostatically pressed precursors fabricated by
the melt route. The foaming agent was CaCO3.
Slika 4: Primer optimiziranja gostote vzorcev aluminijskih pen za
doseganje najve~je sposobnosti absorpcije energije: A) pene, izdelane
po postopku pra{ne metalurgije; B) pene iz strojno obdelanih pre-
kuzorjev, narejenih po livarskem postopku; C) pene iz strojno
obdelanih in hladno izostatsko stisnjenih prekurzorjev, narejenih po
livarskem postopku. Kot sredstvo za penjenje je bil uporabljen CaCO3.
Figure 3: The stress-strain response of various aluminium foam
samples from preforms obtained by the PM route using the CaCO3
foaming agent.
Slika 3: Krivulja napetost – deformacija za vzorce aluminijskih pen
na osnovi predoblik s CaCO3 penilom, izdelanih po postopku pra{ne
metalurgije
The foaming process does not materially affect the
properties of the cell-wall material. However, it leads to
a unique spatial distribution of the aluminium which
results in significantly different properties of the foamed
component in comparison with the bulk part. It is
obvious that the properties of the aluminium foam
significantly depend on its porosity, so that a desired
property (or combination of properties) can be tailored
by selecting the foam density.
The mechanical properties of the foams obtained by
applying dolomite powder as a foaming agent are fully
comparable with the corresponding properties of foams
fabricated using TiH2.
4 CONCLUSION
The following conclusions can be drawn from this
work.
• TiH2 powder as foaming agent was successfully
replaced by commercial CaCO3 or dolomite powders
of a different average particle size.
• Foaming precursors with different proportions ( =
3–10 %) of CaCO3 or dolomite powder particles as a
foaming agent were routinely prepared either by the
powder metallurgical or melt route.
• Precursors obtained by powder metallurgy had
superior homogeneity and densities 98% of theore-
tical. Moreover, in precursors obtained by the PM
route containing  = 3–7 % of CaCO3 or dolomite
particles of an average particle size of below 50 μm,
densities 99 % of theoretical were achieved.
• With greater addition of CaCO3 or dolomite particles
and by using CaCO3 or dolomite powders with a
higher average particle size (above 70 μm), densities
99 % of theoretical could not be achieved.
• The foaming efficiency of experimentally prepared
precursors was evaluated based on the relative
density of foams obtained (the apparent density of the
foam divided by the density of aluminium). The
experimental findings showed that the apparent
density of the foam samples is inversely proportional
to the density of the foaming precursor. Thus,
foamable precursors with a higher density resulted in
foam samples with a lower apparent density and a
higher foaming efficiency. On the other hand, the
foaming efficiency and the average pore size of the
foamed samples are generally reciprocally dependent.
Thus, a higher foaming efficiency results in a foam
microstructure with finer pores.
• The mechanical properties (compression strength and
energy absorption capacity) of the foamed samples
are also strongly influenced by the foaming efficien-
cy. For the range of foam densities analysed, the
compression strength, considered as the initial peak
stress, was found to be superior (approx. 13 MPa) in
samples with increased density (0.55 g/cm3) and
hence, lower foaming efficiency (79.6 %). In contrast
to this, the maximum energy-absorption capacity was
achieved in foams with the highest foaming efficien-
cy.
• From the experimental findings is obvious that the
properties of an aluminium foam significantly depend
on its porosity and the desired property (or combina-
tion of properties) can be tailored by the foam
density.
• The experimental findings confirm that the micro-
structure, compression strength and energy-absorp-
tion capacity of aluminium foams prepared with
CaCO3 or dolomite powder as foaming agent are
quite comparable with their counterparts foamed by
TiH2.
Acknowledgement
This work was supported by funding from the Public
Agency for Research and Development of the Republic
of Slovenia, as well as the Impol Aluminium Company
and Bistral, d. o. o. from Slovenska Bistrica under
contract No. 2410-0206-09.
5 REFERENCES
1 J. Banhart, Adv. Eng. Mater., 8 (2006) 9, 781
2 V. Gergely, D. C. Curra, T. W. Clyne, Composites Science and
Technology, 63 (2003) 2301
3 L. E. G. Cambronero, J. M. Ruiz-Roman, F. A. Corpas, J. M. Ruiz
Prieto, Journal of Materials Processing Technology, 209 (2009) 1803
4 US Patent No. 2.751.289, June 19, 1956
5 T. Nakamura, S. V. Gbyloskurenco, K. Savanto, A. V. Byakova, R.
Ishikawa, Mater. Trans., 43 (2002) 5, 1191–1196
6 J. D. Bryant, M. D. Clowley, M. D. Wilhelmy, J. A. Kallivayalil, W.
Wang, In Metfoam 2007, DEStech Publications, Inc., Lancaster, PA,
2008, pp. 27
7 US Patent No. 7.452.402, November 18, 2008
8 S. Asavavisithchai, R. Tantisiriphaiboon, Chiang Mai J. Sci., 36
(2009) 3, 302
V. KEVORKIJAN et al.: INFLUENCE OF THE FOAMING PRECURSOR’S COMPOSITION AND DENSITY ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 95–103 103

Y. KAZANCOGLU et al.: MULTI-OBJECTIVE OPTIMIZATION OF THE CUTTING FORCES ...
MULTI-OBJECTIVE OPTIMIZATION OF THE CUTTING
FORCES IN TURNING OPERATIONS USING THE
GREY-BASED TAGUCHI METHOD
MULTI NAMENSKA OPTIMIZACIJA STRU@ENJA Z UPORABO
TAGUCHI METODE NA GREY PODLAGI
Yigit Kazancoglu1, Ugur Esme2, Melih Bayramoglu3, Onur Guven4,
Sueda Ozgun5
1Izmir University of Economics, Department of Business Administration, 35330, Balcova-Izmir/Turkey
2Mersin University, Tarsus Technical Education Faculty, Department of Mechanical Education, 33480, Tarsus-Mersin/Turkey
3Cukurova University, Engineering and Architecture Faculty, Department of Mechanical Engineering, 01030, Balcali-Adana/Turkey
4Mersin University, Engineering Faculty, Department of Mechanical Engineering, 33400, Mersin/Turkey
5Mersin University Vocational School of Gülnar, 33400, Gülnar-Mersin/Turkey
uguresme@gmail.com
Prejem rokopisa – received: 2010-06-18; sprejem za objavo – accepted for publication: 2011-02-02
This study investigated the multi-response optimization of the turning process for an optimal parametric combination to yield
the minimum cutting forces and surface roughness with the maximum material-removal rate (MRR) using a combination of a
Grey relational analysis (GRA) and the Taguchi method. Nine experimental runs based on an orthogonal array of the Taguchi
method were performed to derive objective functions to be optimized within the experimental domain. The objective functions
were selected in relation to the parameters of the cutting process: cutting force, surface roughness and MRR. The Taguchi
approach was followed by the Grey relational analysis to solve the multi-response optimization problem. The significance of the
factors on the overall quality characteristics of the cutting process was also evaluated quantitatively using the
analysis-of-variance method (ANOVA). Optimal results were verified through additional experiments. This shows that a proper
selection of the cutting parameters produces a high material-removal rate with a better surface roughness and a lower cutting
force.
Keywords: turning, cutting, Grey relation analysis, Taguchi method, optimization
Raziskani so odgovori optimizacije procesa stru`enja z optimalno kombinacijo parametrov s ciljem dose~i minimalne sile
rezanja in hrapavost povr{ine pri maksimalni odstranitvi materiala (MRR) z uporabo kombinacije Grey odvisnostne analize
(GRA) in Taguchi metode. Devet preizkusov na podlagi ortogonalne ureditve po Taguchi metodi je bilo izvr{eno za razvoj
objektivnih funkcij in njihovo optimizacijo v podro~ju preizkusov. Objektivne funkcije so bile izbrane glede na proces rezanja:
sila rezanja, hrapavost povr{ine in MRR. Taguchi pribli`ek z Grey analizo odvisnosti je bil uporabljen za re{itev problema
optimizacije z ve~ odgovori. Pomen dejavnikov na kakovostne zna~ilnosti procesa rezanja je bil kvantitativno ocenjen z uporabo
metode analize variance (ANOVA). Optimalni rezultati so bili verificirani z dopolnilnimi preizkusi. Rezultati ka`ejo, da prava
izbira parametrov rezanja zagotovi visoko hitrost odstranjevanja materiala pri bolj{i kakovosti povr{ine in manj{i sili rezanja.
Klju~ne besede: stru`enje, rezkanje, Grey analiza, Taguchi metoda, optimizacija
1 INTRODUCTION
Turning is a very important machining process in
which a single-point cutting tool removes material from
the surface of a rotating cylindrical workpiece. The
cutting tool is fed linearly in a direction parallel to the
axis of rotation 1.
As indicated in Figure 1, the turning is carried out on
a lathe that provides the power to turn the workpiece at a
given rotational speed and to feed the cutting tool at a
specified rate and depth of cut. Therefore, three cutting
parameters, i.e., cutting speed (V), feed rate (F), and
depth of cut (d), should be properly selected for a better
surface finish with a lower cutting force.
In a turning operation, it is an important task to select
the cutting parameters to achieve a high cutting
performance. Usually, the desired cutting parameters are
determined based on experience or by using a handbook
1. However, this does not ensure that the selected cutting
parameters have optimal or near optimal cutting perfor-
mance for a particular machine and environment. To
select the cutting parameters properly, several mathema-
tical models 1–6 based on statistical regression techniques
or neural computing have been constructed to establish
Materiali in tehnologije / Materials and technology 45 (2011) 2, 105–110 105
UDK 621.9.025.5:669.14 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 45(2)105(2011)
Figure 1: Schematic representation of the turning process 1
Slika 1: Shema procesa stru`enja 1
the relationship between the cutting performance and the
cutting parameters. Then, an objective function with
constraints is formulated to solve the optimal cutting
parameters using optimization techniques. Therefore,
considerable knowledge and experience are required to
use this modern approach 1. Furthermore, a large number
of cutting experiments has to be performed and analyzed
in order to build the mathematical models. Thus, the
required model building is very costly in terms of time
and materials 1.
Basically, the Taguchi method is a powerful tool for
the design of high-quality systems. It provides a simple,
efficient and systematic approach to optimize the designs
for performance, quality, and cost 1–6. The methodology
is valuable when the design parameters are qualitative
and discrete. Taguchi parameter design can optimize the
performance characteristics through the settings of the
design parameters and reduce the sensitivity of the
system performance to sources of variation. In recent
years, the rapid growth of interest in the Taguchi method
has led to numerous applications of the method in a
world-wide range of industries and countries 1,7,8.
Therefore, this study applied a Taguchi L9 orthogonal
array to plan the experiments on the turning process. The
three controlling factors, including the cutting speed (V),
the depth of cut (d) and feed rate (f), were selected. The
Grey relational analysis is then applied to examine how
the cutting factors influence the cutting force (F), the
surface roughness (Ra) and the material removal rate
(MRR). An optimal parameter combination was then
obtained. Through analyzing the Grey relational grade
matrix, the most influential factors for individual quality
targets of the turning process can be identified. Addition-
ally, an analysis of variance (ANOVA) was also utilized
to examine the most significant factors for the F, Ra and
MRR in the turning process.
2 GREY RELATIONAL ANALYSIS (GRA)
2.1 Data Preprocessing
In a Grey relational analysis, experimental data, i.e.,
measured features of the quality characteristics, are first
normalized, ranging from zero to one. This process is
known as Grey relational generation. Next, based on
normalized experimental data, the Grey relational
coefficient is calculated to represent the correlation
between the desired and the actual experimental data.
Then overall Grey relational grade is determined by
averaging the Grey relational coefficient corresponding
to selected responses 9. The overall performance cha-
racteristic of the multiple response process depends on
the calculated Grey relational grade. This approach con-
verts a multiple-response process-optimization problem
into a single-response optimization situation with the
objective function of the overall Grey relational grade.
The optimal parametric combination is then evaluated,
which would result in the highest Grey relational grade.
The optimal factor setting for maximizing the overall
Grey relational grade can be performed using the
Taguchi method 9,10.
In Grey relational generation, the normalized F and
Ra corresponding to the smaller-the-better (SB) criterion
which can be expressed as:
x k
y k y k
y k y ki
i i
i i
( )
max ( ) ( )
max ( ) min ( )
=
−
−
(1)
MRR should follow the larger-the-better (LB) crite-
rion, which can be expressed as:
x k
y k y k
y k y kj
j j
j j
( )
( ) min ( )
max ( ) min ( )
=
−
−
(2)
where xi(k) and xj(k) are the value after the Grey
relational generation for the SB and LB criteria,
respectively. Min yi(k) is the smallest value of yi(k) and
for the kth response, and max yi(k) is the largest value of
yi(k) for the kth response 9. An ideal sequence is x0(k)
(k=1, 2,…,m) for the responses. The definition of the
Grey relational grade in the course of the Grey relational
analysis is to reveal the degree of relation between the 9
sequences [x0(k) and xi(k), k = 1,2,…,m and i =
1,2,…,9]. The Grey relational coefficient i(k) can be
calculated as:


i i
k
k
( )
( )
min max
max
=
−
+
Δ Δ
Δ Δ0
(3)
where Δ 0 0i ix k x k= −( ) ( ) is the difference of the
absolute value x0(k) and xi(k);  is the distinguishing
coefficient 0 
  
 1; Δ min
min min( )k j i k= ∀ ∈ ∀
x k x kj0 ( ) ( )− is the smallest value of 0i; and
Δ max
max max( ) ( ) ( )k j i k x k x kj= ∀ ∈ ∀ −0 is the largest
value of 0i. After averaging the Grey relational coeffi-
cients, the Grey relational grade i can be computed as:
 i i
k
n
n
k=
=
∑1
1
( ) (4)
where n is the number of process responses. The higher
value of the Grey relational grade corresponds to an in-
tense relational degree between the reference sequence
x0(k) and the given sequence xi(k). The reference
sequence x0(k) represents the best process sequence;
therefore, a higher Grey relational grade means that the
corresponding parameter combination is closer to the
optimal 9. The mean response for the Grey relational
grade with its grand mean and the main effect plot of the
Grey relational grade are very important because the
optimal process condition can be evaluated from this
plot 9.
106 Materiali in tehnologije / Materials and technology 45 (2011) 2, 105–110
Y. KAZANCOGLU et al.: MULTI-OBJECTIVE OPTIMIZATION OF THE CUTTING FORCES ...
3 EXPERIMENTAL PROCEDURE AND TEST
RESULTS
3.1 Experimental Details
The cutting experiments were carried out on an
experimental lathe setup using a HSS cutting tool for the
machining of the AISI 1050 steel bar, which is 30 mm in
diameter and 80 mm in length. The mechanical
properties and percent composition of the workpiece
material is listed in Table 1.
A Phynix TR-100 model surface-roughness tester was
used to measure the surface roughness of the machined
samples. The cut-off length () was chosen as 0.3 for
each roughness measurement. An average of six
measurements of the surface roughness was taken to use
in the multi-criteria optimization. Also, the MRR
(mm3/min) was calculated using Eq. (5);
MRR = 1000 Vfd (5)
where f/(mm/r) denotes the feed rate, d/mm describes
the cutting depth and V/(m/min) represents the cutting
speed of the turning operation.
3.2 Process Parameters and Test Results
In full factorial design, the number of experimental
runs exponentially increases as the number of factors, as
well as their level increases. This results in a huge
experimentation cost and considerable time periods 9. So,
in order to compromise these two adverse factors and to
search for the optimal process condition through a
limited number of experimental runs Taguchi’s L9
orthogonal array consisting of 9 sets of data was selected
to optimize the multiple performance characteristics of
the turning process. Experiments were conducted with
the process parameters given in Table 2, to obtain the
machined surface on the AISI 1050 medium-carbon
steel. The feasible space for the cutting parameters was
defined by varying the cutting speed in the range
110–600 m/min, the feed rate in the range 0.2–0.6
mm/min, and the depth of cut in the range 0.5–1.5 mm.
The initial cutting parameters were selected as:
cutting speed of 110 m/min; feed rate of 0.20 mm/min;
and depth of cut of 0.5 mm. In the cutting parameter
design, three levels of the cutting parameters were
selected, as shown in Table 2. In order to prevent a
sudden increase of the cutting forces due to the dullness
of the cutting edge, the HSS tool was changed after three
repetitions of each experiment.
Table 3 shows the selected design matrix based on
the Taguchi L9 orthogonal array consisting of 9 sets of
coded conditions and the experimental results for the
responses of F, Ra and MRR. All these data were utilized
for the analysis and evaluation of the optimal parameter
combination required to achieve the desired quality
within the experimental domain.
4 PARAMETRIC OPTIMIZATION OF THE
CUTTING PROCESS
4.1 Evaluation of the Optimal Process Condition
First, by using Eqs. (1) and (2), the experimental data
were normalized to obtain the Grey relational generation
9. The normalized data and 0i(k) for each of the
responses are listed in Table 4 and Table 5, respectively.
For MRR the larger-the-better (LB) and for F and Ra the
smaller-the-better (SB) criteria were selected.
Y. KAZANCOGLU et al.: MULTI-OBJECTIVE OPTIMIZATION OF THE CUTTING FORCES ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 105–110 107
Table 1: Chemical and mechanical properties of AISI 1050 medium carbon steel
Tabela 1: Kemi~na sestava in mehanske lastnosti jekla AISI 1050 s srednjim ogljikom
Chemical composition
w/%
C P S Mn Cr Fe Ni Cu
0.49 0.02 0.02 0.78 0.08 97.99 0.10 0.26
Mechanical properties
Yield strength (MPa) Tensile strength (MPa) Elongation (%) Vickers Hardness (HV)
365 636 24 261
Table 2: Cutting parameters and their limits
Tabela 2: Parametri rezanja in njihove meje
Cutting Parameters Notation Unit
Levels of factors
1 2 3
Cutting speed V m/min 110* 300 600
Feed rate f mm/min 0.2* 0.4 0.6
Depth of cut d mm 0.5* 1.0 1.5
*Initial cutting parameter
Table 3: Orthogonal array L9 of the experimental runs and results
Tabela 3: Ortogonalna razporeditev L9 eksperimentov in rezultati
Run no
Parameter level Experimental results
V f d MRR/(mm3/min) F/N Ra/μm
1 1 1 1 0.11 123 0.87
2 1 2 2 0.44 179 2.33
3 1 3 3 0.99 364 6.62
4 2 1 2 0.60 166 1.98
5 2 2 3 1.80 295 3.82
6 2 3 1 0.90 255 3.96
7 3 1 3 1.80 340 0.92
8 3 2 1 1.20 218 1.22
9 3 3 2 3.60 268 5.60
Table 4: Grey relational generation of each performance characte-
ristics
Tabela 4: Grey relacijska generacija vsake karakteristike performance
Run no
MRR F Ra
Larger-the-
better
Smaller-the-
better
Smaller-the-
better
Ideal sequence 1.000 1.000 1.000
1 0.000 1.000 1.000
2 0.095 0.768 0.746
3 0.252 0.000 0.000
4 0.140 0.822 0.807
5 0.484 0.286 0.487
6 0.226 0.452 0.463
7 0.484 0.100 0.991
8 0.312 0.606 0.939
9 1.000 0.398 0.177
Table 5: Evaluation of 0i(k) for each of the responses
Tabela 5: Ocena 0i(k) za vsak odgovor
Run no MRR F Ra
Ideal sequence 1.000 1.000 1.000
1 1.000 0.000 0.000
2 0.905 0.232 0.254
3 0.748 1.000 1.000
4 0.860 0.178 0.193
5 0.516 0.714 0.513
6 0.774 0.548 0.537
7 0.516 0.900 0.009
8 0.688 0.394 0.061
9 0.000 0.602 0.823
Table 6 shows the calculated Grey relational coeffi-
cients (with the weights of MRR = 0.33, F = 0.33 and
Ra = 0.33) of each performance characteristic using Eq.
(3).
Table 6: Grey relational coefficient of each performance characte-
ristics (MRR = 0.33, F = 0.33 and Ra = 0.33)
Tabela 6: Grey odvisnostni koeficient za vsako zna~ilnost performance
(MRR = 0.33, F = 0.33 and Ra = 0.33)
Run no MRR F Ra
Ideal sequence 1.000 1.000 1.000
1 0.248 1.000 1.000
2 0.267 0.587 0.565
3 0.306 0.248 0.248
4 0.277 0.649 0.631
5 0.390 0.316 0.391
6 0.299 0.376 0.380
7 0.390 0.268 0.974
8 0.324 0.456 0.844
9 1.000 0.354 0.286
The Grey relational coefficients, given in Table 7, for
each response have been accumulated by using Eq. (4) to
evaluate the Grey relational grade, which is the overall
representative of all the features of the cutting-process
quality. Thus, the multi-criteria optimization problem has
been transformed into a single equivalent objective
function optimization problem using a combination of
the Taguchi approach and Grey relational analyses. The
higher is the value of the Grey relational grade, the
corresponding factor combination is said to be close to
the optimal 9.
Table 7: Grey relational grade
Tabela 7: Grey stopnja odvisnosti
Run no Grey relationalgrade Rank
1 0.7119 1
2 0.4683 6
3 0.2648 9
4 0.5139 5
5 0.3623 7
6 0.3483 8
7 0.5388 3
8 0.5360 4
9 0.5414 2
The signal-to-noise (S/N) ratio is a measure of the
magnitude of a data set relative to the standard deviation.
If the S/N is large, the magnitude of the signal is large
relative the noise, as measured with the standard
deviation 11. Table 8 shows the S/N ratio based on the
larger-the-better criterion for the overall Grey relational
grade calculated using Eq. (6).
S N
n y ii
n
/ log= −
⎡
⎣⎢
⎤
⎦⎥=
∑10 1 12
1
(6)
where n is the number of measurements, and yi is the
measured characteristic value.
Table 8: S/N ratio for overall Grey relational grade
Tabela 8: S/N razmerje za splo{no Grey stopnjo
Run no S/N
1 –2.59
2 –6.59
3 –11.54
4 –5.78
5 –8.82
6 –9.16
7 –5.37
8 –5.42
9 –5.33
Table 9: Response Table for the mean Grey relational grade
Tabela 9: Tabela odgovorov za povpre~no Grey stopnjo odvisnosti
Factors Grey relational grade
Level 1 Level 2 Level 3 max-min
V 0.49 0.41 0.54 0.13
f 0.60 0.46 0.38 0.22
d 0.54 0.51 0.39 0.15
Total mean Grey relational grade = 0.48
A graphical representation of the S/N ratio for the
overall Grey relational grade is shown in Figure 2. The
dashed line is the value of the total mean of the S/N ratio.
Y. KAZANCOGLU et al.: MULTI-OBJECTIVE OPTIMIZATION OF THE CUTTING FORCES ...
108 Materiali in tehnologije / Materials and technology 45 (2011) 2, 105–110
As indicated in Figure 2, the optimal condition for
the turning of the AISI 1050 medium-carbon steel
becomes V3 f1 d1. Table 9 shows the mean Grey relational
grade ratio for each level of the process parameters.
4.2 Analysis of Variance (ANOVA)
The purpose of the analysis of variance (ANOVA) is
to investigate which turning parameters significantly
affect the performance characteristics 8–10. This is accom-
plished by separating the total variability of the grey
relational grades, which is measured by the sum of the
squared deviations from the total mean of the grey
relational grade, into contributions from each of the
turning parameters and the error 9. Thus;
SS SS SST F e= + (7)
where
SS j m
j
p
T = −
=
∑ ( )  2
1
(8)
SST – Total sum of the squared deviations about the mean
j – Mean response for the jth experiment
m – Grand mean of the response
p – Number of experiments in the orthogonal array
SSF – Sum of the squared deviations due to each factor
SSe – Sum of the squared deviations due to error
In addition, the F test was used to determine which
turning parameters have a significant effect on the per-
formance characteristic. Usually, the change of the
turning parameter has a significant effect on the
performance characteristics when the F value is large
8–10. The ANOVA for the overall Grey relational grade is
shown in Table 10.
Table 10: ANOVA results of turning process parameters
Tabela 10: ANOVA rezultati parametrov procesa stru`enja
Parameter Degree ofFreedom
Sum of
Square
Mean
Square F
Contribu-
tion (%)
V 2 0.026 0.013 1.21 17.81
f 2 0.071 0.035 3.27 48.63
d 2 0.039 0.019 1.79 26.71
Error 2 0.010 0.011 6.85
Total 8 0.158 100
According to this analysis, the most effective
parameters with respect to the material-removal rate, the
cutting force and the surface roughness are the feed rate,
the depth of cut and the cutting speed. The percentage
contribution indicates the relative power of a factor to
reduce the variation. For a factor with a high percentage
contribution, there is a great influence on the perfor-
mance. The percent contributions of the cutting para-
meters on the material-removal rate, the cutting force
and the surface roughness are shown in Table 10 and
Figure 3. The feed rate was found to be the major factor
affecting the material-removal rate, the cutting force and
the surface roughness (48.63 %), whereas the depth of
cut (26.71 %) and the cutting speed (17.81 %) were
found to be the second- and third-ranking factors
respectively.
4.3 Confirmation Test
After evaluating the optimal parameter settings, the
next step is to predict and verify the enhancement of the
quality characteristics using the optimal parametric
combination 9. The estimated Grey relational grade 
using the optimal level of the design parameters can be
calculated as:
 ( )   = + −
=
∑m j m
i
o
1
(9)
where m is the total mean Grey relational grade, i is
the mean Grey relational grade at the optimal level, and
Materiali in tehnologije / Materials and technology 45 (2011) 2, 105–110 109
Y. KAZANCOGLU et al.: MULTI-OBJECTIVE OPTIMIZATION OF THE CUTTING FORCES ...
Figure 2: (a) Mean plot, (b) S/N plot for the Grey relational grade
Slika 2: (a) Povpre~na odvisnost, (b) S/N odvisnost za Grey stopnjo
odvisnosti
Figure 3: Contribution percentage of the cutting parameters
Slika 3: Prispevni procent parametrov rezanja
o is the number of the main design parameters that affect
the quality characteristics 9. Table 11 indicates the
comparison of the predicted tensile strength and
elongation with that of the actual by using the optimal
turning conditions. Good agreement between the actual
and the predicted results has been observed (the impro-
vement in the overall Grey relational grade was found to
be as 0.20).
In the Taguchi method, the only performance feature
is the overall Grey relational grade and the aim should be
to search for a parameter setting that can achieve the
highest overall Grey relational grade 9. The Grey rela-
tional grade is a representative of all the individual
performance characteristics. In the present study, the
objective functions were selected in relation to the
parameters of the material-removal rate, the cutting force
and the surface roughness. The importance weights of
the material-removal rate, the cutting force and the
surface roughness were equally adjusted to be 0.33.
The results show that using the optimal parameter
setting (V3 f1 d1) causes a lower cutting force and surface
roughness with a higher material removal rate and hence
a better surface finish.
5 CONCLUSIONS
In this study, the Grey-based Taguchi method was
applied for the multiple performance characteristics of
turning operations. A grey relational analysis of the
material-removal rate, the cutting force and the surface
roughness obtained from the Taguchi method reduced
from the multiple performance characteristics to a single
performance characteristic which is called the grey
relational grade. Therefore, the optimization of the
complicated multiple performance characteristics of the
processes can be greatly simplified using the Grey-based
Taguchi method. It is also shown that the performance
characteristics of the turning operations, such as the
material removal rate, the cutting force and the surface
roughness are greatly enhanced by using this method.
6 REFERENCES
1 W. H. Yang, Y. S. Tarng, Design Optimization of cutting parameters
for turning operations based on the Taguchi method, Journal of
Materials Processing Technology, 84 (1988), 122–129
2 P. L. B. Oxley, Modelling machining processes with a view to their
optimization and the adaptive control of metal cutting machine tools,
Robot. Comput.-Integrated Manuf., 4 (1988), 103–119
3 G. Chryssolouris, M. Guillot, A comparison of statistical and AI
approaches to the selection of process parameters in intelligent
machining, ASME J. Eng. Ind., 112 (1990), 122–131
4 Y. Yao, X. D. Fang, Modelling of multivariate time series for tool
wear estimation in finish turning, Int. J. Mach. Tools Manuf., 32
(1992) 4, 495–508
5 C. Zhou, R A. Wysk, An integrated system for selecting optimum
cutting speeds and tool replacement times, Int. J. Mach. Tools
Manuf. 32 (1992) 5, 695–707
6 M. S. Chua, M. Rahman, Y. S. Wong, H. T. Loh, Determination of
optimal cutting conditions using design of experiments and
optimization techniques, Int. J. Mach. Tools Manuf., 33 (1993) 2,
297–305
7 A. Bendell, J. Disney, W. A. Pridmore, Taguchi methods: Appli-
cations in world industry, IFS Publications, UK, 1989
8 D. C. Montgomery, Design and analysis of experiments, Wiley,
Singapore, 1991
9 S. Datta, A. Bandyopadhyay, P. K. Pal, Grey-based taguchi method
for optimization of bead geometry in submerged arc bead-on-plate
welding. Int J Adv Manuf Technol., 39 (2008) 11, 1136–1143
10 U. Esme, M. Bayramoglu, Y. Kazancoglu, S. Özgun, Optimization of
weld bead geometry in TIG welding process using Grey Relation
Analysis and Taguchi Method, Mater. Tehnol., 43 (2009), 143–149
11 D. S. Holmes, A. E. Mergen, Signal to noise ratio – What is the right
size, www.qualitymag.com/.../Manuscript%20Holmes%20&%20
Mergen.pdf, USA, 1996, 1–6
Y. KAZANCOGLU et al.: MULTI-OBJECTIVE OPTIMIZATION OF THE CUTTING FORCES ...
110 Materiali in tehnologije / Materials and technology 45 (2011) 2, 105–110
Table 11: Results of confirmation test
Tabela 11: Rezultati potrditvenega preizkusa
Initial factor settings
Optimal process condition
Prediction Experiment
Factor levels V1 f1 d1 V3 f1 d1 V3 f1 d1
MRR/ (mm3/min) 0.11 0.30
F/N 123 115
Ra/ìm 0.87 0.65
S/N ratio of overall Grey relational grade –2.59 –2.21 -–2.18
Overall Grey relational grade 0.71 0.72 0.77
Improvement in Grey relational grade = 0.06
Z. ADOLF et al.: THERMODYNAMIC CONDITIONS FOR THE NUCLEATION OF BORON COMPOUNDS ...
THERMODYNAMIC CONDITIONS FOR THE
NUCLEATION OF BORON COMPOUNDS DURING THE
COOLING OF STEEL
TERMODINAMI^NI POGOJI ZA NUKLEACIJO BOROVIH SPOJIN
PRI OHLAJANJU JEKLA
Zdenìk Adolf, Jiøí Ba`an, Ladislav Socha
V[B-Technical University of Ostrava, Department of Metallurgy, 17. listopadu 15/2172, 708 33 Ostrava-Poruba, Czech Republic
zdenek.adolf@vsb.cz
Prejem rokopisa – received: 2010-11-08; sprejem za objavo – accepted for publication: 2011-01-31
The higher the value of the product of boron and oxygen concentrations, or of boron and nitrogen concentrations, than the value
corresponding to the balance for the given temperature, is the thermo-dynamic criterion for the nucleation of a new phase (B2O3
or BN). It follows from the calculations that the theoretical temperature of the beginning of B2O3 nucleation is higher than the
temperature of the beginning of BN nitride nucleation. During the solidification and cooling down of steel the boron oxide will
be formed preferentially before the boron nitride.
Keywords: boron steel, nucleation, non-metallic inclusions
Ve~ji produkt vsebnosti bora in kisika ter bora in du{ika kot ravnote`na vrednost pri izbrani temperaturi je termodinami~ni
pogoj za nukleacijo nove faze (B2O3 oz. BN). Izra~uni so pokazali, da je teoreti~no temperatura za~etka nukleacije oksida B2O3
vi{ja kot pri nitridu BN. Zato bo pri ohlajanju jekla borov oksid nastal pred nitridom.
Klju~ne besede: bor, jeklo, nukleacija, nekovinski vklju~ki
1 INTRODUCTION
This paper presents a thermodynamic analysis of the
probability of the formation of boron oxide and nitride in
boron- and nitrogen-microalloyed stainless steels. The
steels are designated for forgings for the production of
valves of nuclear power plants’ primary circuits and their
chemical composition is given in Table 1.
Table 1: Chemical composition of the steel (in mass fractions w/%)
Tabela 1: Kemi~na sestava jekel (v masnih dele`ih w/%)
Element. (w/%)
C 0.04 S 0.001 V 0.05 Nb 0.017
Mn 1.57 Cr 17.5 W 0.02 B 0.004
Si 0.6 Ni 10.5 Al 0.050 N 0.0126
P 0.020 Mo 0.07 Ti 0.40
The objective of this work is to determine at which
temperatures the B2O3 or BN is formed during the
cooling and solidification of the steel. The formation of a
new phase (inclusion) is related to the content of boron,
nitrogen and oxygen in the steel.
2 THERMODYNAMIC BALANCE
The temperature dependencies of the Gibbs energy
for the formation of B2O3 and BN were derived from
table data1, 2 with the use of the equations:
[B] + 3/2 [O] =
1/2 B2O3(l) (1)
[B] + [N] = BN(s) (2)
ΔG1
0 = –411 990 + 143,585 T (3)
ΔG 2
0 = –227 737 + 97,95 T (4)
Due to the fact that the melting temperature of B2O3
is 450 °C, this oxide is at the temperatures of steel soli-
dification in the liquid state2.
It is possible to derive from equations (3) and (4) the
following relations for the temperature dependencies of
the equilibrium constants:
lg K
T1
21517
7 50= − , (5)
lg K
T2
11894
5116= − , (6)
The following is valid for the equilibrium constants
of the reactions (1) and (2):
( )K
a
a a
1 =
⋅
B O
1 /2
equilibrium
2 3
[B] [O]
3/ 2
(7)
( )K
a
a a
2 =
⋅
BN
equilibrium[B] [N]
(8)
Assuming that pure boron oxide and boron nitride are
formed, it is possible to consider their activities to be
equal to one. It is similarly possible to assume unequivo-
cally that the solutions of boron, oxygen and nitrogen in
steel are diluted and that the activities of these elements
Materiali in tehnologije / Materials and technology 45 (2011) 2, 111–113 111
UDK 669.14.018.8:621.785 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 45(2)111(2011)
are equal to a mass percentage. It is then possible to
adjust the equations (7) and (8) to these forms:
( )[B] [O]3/ 2⋅ =
equilibrium
1
1K
(9)
( )[B] [N]⋅ =
equilibrium
1
2K
(10)
Due to the fact that the equilibrium constants K1 and
K2 are a function of temperature only in accordance with
equations (5) and (6), the equilibrium products of the
concentrations of boron and oxygen, as well as boron
and nitrogen, depend only on the temperature (see
equations (11) and (12)) and can be calculated from the
temperature dependencies.
( )lg ,[B] [O]3/ 2⋅ = − +
equilibrium
21517
7 50
T
(11)
( )lg ,[B] [N]⋅ = − +
equilibrium
11894
5116
T
(12)
It is subsequently possible to affirm logically that the
formation of boron trioxide or boron oxide at the
temperature T is conditioned by a higher value of the
real product of boron and oxygen, or a boron and
nitrogen concentration that would correspond to the
equilibrium.
( ) ( )[B] [O] [B] [O]3/ 2 3/ 2⋅ ≥ ⋅
real equilibrium
(13)
( ) ( )[B] [N] [B] [N]⋅ ≥ ⋅
real equilibrium
(14)
It follows from equations (11) and (12) that with
decreasing temperature the value of equilibrium products
(13) and (14) also decreases, and therefore the proba-
bility of the formation of the inclusions B2O3 and BN
increases, since the real products (13) and (14) remain
constant.
3 DISCUSSION OF THE RESULTS
The derived relationships were applied to the steel
microalloyed with boron and nitrogen of required chemi-
cal composition – see Table 1.
The theoretical dependencies of the temperatures of
the beginning of formation of the B2O3 or BN on the
content of oxygen and nitrogen are given in Figures 1
and 2.
These dependencies were calculated from equations
(5) and (6), adjusted for mass fractions 0.001 %, 0.006 %,
0.03 % and 0.05 % of boron.
It follows from Figure 2 that, for example, the
theoretical temperature of the beginning of nucleation of
BN nitride is for the nitrogen content of 100 μg/g (ppm)
in the interval 903 °C, 1001 °C, 1104 °C and 1040 °C, or
for 200 μg/g of nitrogen in the interval 939 °C, 10043
°C, 1154 °C and 1193 °C. The theoretical temperature
for the start of nucleation of the oxide B2O3 is for the
achieved oxygen contents (10 μg/g) higher, i.e., 1161 °C,
1240 °C, 1318 °C and 1345 °C, and for 20 μg/g of
oxygen it is 1206 °C, 1290 °C, 1373 °C and 1402 °C. It
follows, therefore, that boron oxide will be formed pre-
ferentially before boron nitride during the cooling of the
steel.
For the steels with the above-mentioned boron
contents the equilibrium temperatures and equilibrium
Z. ADOLF et al.: THERMODYNAMIC CONDITIONS FOR THE NUCLEATION OF BORON COMPOUNDS ...
112 Materiali in tehnologije / Materials and technology 45 (2011) 2, 111–113
Figure 3: Equilibrium temperatures and equilibrium contents of
nitrogen corresponding to the oxygen content in steel of 10–100 μg/g
Slika 3: Ravnote`ne temperature in ravnote`ne vsebnosti du{ika pri
vsebnosti 10–100 μg/g kisika v jeklu
Figure 1: Dependence of the beginning temperature of formation of
B2O3 and the content of oxygen in steel
Slika 1: Za~etna temperatura tvorbe B2O3 v odvisnosti od vsebnosti
kisika v jeklu
Figure 2: Dependence of the starting temperature on the formation of
BN and the content of nitrogen in steel
Slika 2: Za~etna temperatura tvorbe B2O3 v odvisnosti od vsebnost
du{ika v jeklu
nitrogen contents corresponding to 10–100 μg/g of
oxygen in steel are shown in Figure 3. It is evident from
the figure that B2O3 oxides can be formed only at higher
contents of oxygen than would correspond to an
equilibrium (right to the curve). Similarly, boron nitrides
can be formed only at higher nitrogen contents than
would correspond to an equilibrium (above the curve),
since inequalities (13) and (14) are fulfilled.
4 CONCLUSIONS
The thermodynamic balance of probability for the
formation of oxide B2O3 and nitride BN in boron-
microalloyed steels was calculated. The balance proved
that the oxide is more stable than the boron nitride and,
therefore, during the cooling of steels it is formed
preferentially.
The work was prepared at the conclusion of the
projects FR-TI1/477 and FR-TI1/222 under the financial
support of the Ministry of Industry and Trade (MPO
^R).
5 REFERENCES
1 J. Fruehan, et al. The Making Shaping and Treating of Steel. Pitts-
burgh 1998, 767 pp., ISBN 0-930767-02-0
2 J. Leitner. Database of thermodynamic data for admixtures in iron
based melts. V[CHT Praha, 2002, 23 pp.
Z. ADOLF et al.: THERMODYNAMIC CONDITIONS FOR THE NUCLEATION OF BORON COMPOUNDS ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 111–113 113

T. DEMÝR, M. ÜBEYLÝ: A MICRO-DAMAGE INVESTIGATION ON A LOW-ALLOY STEEL ...
A MICRO-DAMAGE INVESTIGATION ON A
LOW-ALLOY STEEL TESTED USING A 7.62-mm AP
PROJECTILE
RAZISKAVA MIKROPO[KODB MALOLEGIRANEGA JEKLA PO
PREIZKUSU S KROGLO AP 7,62 mm
Teyfik Demir1, Mustafa Übeyli2
1TOBB University of Economics and Technology, Engineering Faculty, Mechanical Engineering, 06560 Ankara - TÜRKÝYE
2Osmaniye Korkut Ata University, Engineering Faculty, Mechanical Engineering, 80000 Osmaniye - TÜRKÝYE
mubeyli@gmail.com
Prejem rokopisa – received: 2011-01-21; sprejem za objavo – accepted for publication: 2011-02-23
Low-alloy steels have been widely used as an armor material in defense systems due to their superior mechanical and
weldability properties. In this study, microdamage formation in an AISI 4140 steel, which was hit by 7.62-mm armor-piercing
(AP) projectile, was investigated. Analyses of the microstructure and microhardness after ballistic testing were performed to
correlate the microstructure-property relationship. Two different types of adiabatic shear band (ASB) were observed in the
tested samples. The experimental results indicated that the type and hardness of the bands were strongly related to the hardness
and ballistic performance of the steel specimens.
Keywords: high-strength steel, adiabatic shear band, armor.
Malolegirana jekla se uporabljajo za obrambne sisteme predvsem zaradi njihovih bolj{ih lastnosti in varivosti. V tem delu je
raziskan nastanek po{kodbe v jeklu AISI 4140 po udarcu prebojne krogle 7,62 mm v oklep. Mikrostruktura in mikrotrdota sta
po balisti~nem preizkusu korelirana z odnosom mikrostruktura-lastnosti. V preizkusnih vzorcih sta opa`eni dve vrsti adiabatskih
stri`nih pasov (ASB). Rezultati so pokazali, da sta vrsta in trdota pasov zelo povezani s trdoto in balisti~nimi lastnostmi vzorcev
jekla.
Klju~ne besede: visokotrdna jekla, adiabatski stri`ni pasovi, oklep
1 INTRODUCTION
Microstructural changes, i.e., phase transformations,
distortion of grains, elongation of grains, grain refine-
ment, are observed in metallic materials that are exposed
to high strain rates and very large strains. With these
changes, bands that are different from the original matrix
in nature or alignment may form. These bands are
termed adiabatic shear bands and they occur in a number
of processes, such as dynamic impact, metal forming,
ballistic testing, machining and high-strain-rate defor-
mation. The adiabatic shear band phenomenon was first
discovered by Zener and Hollomon 1. However, studies
on the adiabatic shear band formation accelerated after
the year 1970. The thermomechanical instability caused
by the presence of local inhomogeneity, inducing local
deformation and heating, results in the formation of
shear bands 2. Adiabatic shearing is a special case where
the heat generated in the localized bands cannot be
transferred easily to the surrounding material due to the
strain rate and thermal properties of the material 2.
However, in a real situation some of the heat loss to the
surrounding material always occurs. The reason for
using the adiabatic term is to indicate that a large part of
the heat is retained in the band 2. This leads to a local
temperature increase which may cause a microstructure
change in the material 3. Adiabatic shearing can lead to a
catastrophic failure in materials under high strain rate
deformation 4,5. There are two general classifications of
the ASBs: deformed and transformed 2. In the deformed
ASBs, there is no change in the microstructure of the
materials, but the grains or structure are highly distorted.
On the other hand, in the transformed bands, a
crystallographic phase change takes place.
Derep 6 investigated the microstructure transforma-
tion induced by the adiabatic shearing in an armor steel
tested using a hollow charge. It was observed that a
fine-grained, equiaxed structure of delta ferrite con-
sisting of very narrow martensite laths formed in the
center of ASBs. In another study, Wittman et al. 7
investigated the adiabatic shear-band formation in AISI
4340 steel. Also, an investigation of the adiabatic shear
band formation in different steels (AISI 1018, AISI 4340
and HY-100) was performed after applying torsion loads
at high strain rates 8. It was mentioned that the shear
bands were produced by the dynamic torsional
deformation in all cases and they were of the deformed
type for the steel AISI 1018, whereas those were
transformed for the steel AISI 4340. In addition, Bassim
9 studied ASB formation in the AISI 4340 steel during
high strain rate deformation using a torsional Split
Hopkinson Bar System. It was also reported that ASBs
initiated at local defects and inhomogeneities. In
addition, the specimen geometry and dimensions were
found to be contributing factors to the development of
ASBs. Furthermore, Duan et al. 10 examined the forma-
Materiali in tehnologije / Materials and technology 45 (2011) 2, 115–120 115
UDK 669.15-194:620.18 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 45(2)115(2011)
tion of adiabatic shear bands in two different grades of
steel (medium carbon steel and 30CrMnMo) subjected to
the impact of tungsten projectiles. They reported that
three kinds of ASBs formed in the 30CrMnMo, while the
medium carbon steel did not easily produce ASBs.
Moreover, Hu et al. 11 investigated the ASB formation in
various steels (4130, AISI 1045, modified rolled
homogeneous armor steel and AerMet 100) tested by 44
grain-fragment simulating projectiles. It was shown that
transformed bands were found in the steels, 4130,
modified rolled homogeneous armor steel and AerMet
100, while deformed bands were found in the AISI 1045.
Recently, the effect of heat treatment on the ASBs in
AISI 4340 steel tested with the Split Hopkinson Pressure
Bar was studied 12. It was concluded that heat treatment
did not cause any significant effect on the microstructure
and hardness of ASBs. In a more recent work, Lins et al.
13 made a microstructural characterization of the ASBs in
an interstitial-free steel deformed at high strain rates
(>2.8 · 104 s–1) in a Split Hopkinson Pressure Bar. They
found that the deformation twins occurred throughout
the microstructure. Recently, a microstructural study on
the AISI 4340 and DIN 100Cr6 steels, tested with a
7.62-mm AP projectile, was carried out 14. It was stated
that the impact of the AP projectile resulted in the
formation of ASBs in all the samples having different
hardness levels 14. In addition to the studies conducted on
steels, shear band formation in copper 15–17, aluminum
18,19 and refractory metals 20–23 was also examined.
The ballistic behavior of the high-strength, low-alloy
steel AISI 4140 with respect to its hardness and areal
density was determined against a 7.62-mm armor-
piercing (AP) projectile in our previous work 24. In this
study, a micro-damage analysis in this type of steel after
the ballistic testing with the 7.62-mm AP projectile was
made.
2 EXPERIMENTAL PROCEDURE
The AISI 4140 steel is a high-strength, low-alloy
steel that can be strengthened significantly by applying a
martensite formation heat treatment. This steel is used
mainly in the automotive industry and for structural
applications. The AISI 4140 steel used in this work had
0.40 % C, 1.0 % Cr, 0.2 % Mo, 0.7 5% Mn and 0.2 % Si
(mass fractions, w/%). Experimental studies were based
mainly on two lines: a microstructural examination and a
microhardness measurement of the ASBs formed in AISI
4140 steel that was subjected to the impact of a 7.62-mm
AP projectile 24. This type of projectile is widely used in
NATO armies. Its specification can be found in ref. 24.
The thicknesses of the steel samples utilized in this work
were 7.2 mm, 9 mm, 10.8 mm, 12.7 and 14.4 mm. Fur-
thermore, three hardness levels (38 HRC, 50 HRC and
60 HRC) were considered for each thickness of the steel
sample for the comparison with the previous study 14. In
coding of the steel specimens, the letters A, B and C
were used to indicate the hardness levels of 38 HRC, 50
HRC and 60 HRC, respectively. In addition, the
numbers, 1, 2, 3, 4 and 5 following these letters were
utilized to refer to the steel thickness of 7.2 mm, 9 mm,
10.8 mm, 12.7 and 14.4 mm, respectively.
The microstructural changes in the specimens were
firstly determined using an optical microscope to see the
type and distribution of the ASBs in the material. For
this purpose, the AISI 4140 steel specimens were etched
using 5 % nital solution to see both the phase(s) and the
adiabatic shear bands clearly. Next, the microstructures
of the samples were investigated at the impact surfaces
through the thickness of the steel and parallel to the
projectile motion (Figure 1). Secondly, an analysis using
a scanning electron microscope (SEM) was carried out to
resolve the microstructure of the white ASBs more
clearly. Moreover, a representative chemical content of
the precipitate was determined by energy-dispersive
X-ray (EDX) analysis. Finally, the microhardness
measurements were also made to determine the hardness
of the different types of adiabatic shear bands formed
upon impact.
3 EXPERIMENTAL RESULTS AND DISCUSSION
Typical microstructures of the sample before testing
are represented in Figure 2. Due to the applied heat
treatment, all the samples had a tempered martensite
phase containing ferrite with very small iron-carbide
islands. This type of phase provides the best combination
of strength and toughness for a low-alloy steel. For this
reason, this type of microstructure is frequently chosen
to be used in high-strength applications.
The kinetic energy of the projectile was able to
generate ASBs in the AISI 4140 steel. Its velocity was
measured to be (782 ± 5.4) m/s 24. The most important
factor in ballistic studies is the velocity of the impact.
Damage types change significantly with the impact
velocity. The velocities between 700 m/s and 3000 m/s
are termed intermediate impact velocities, which cause
T. DEMÝR, M. ÜBEYLÝ: A MICRO-DAMAGE INVESTIGATION ON A LOW-ALLOY STEEL ...
116 Materiali in tehnologije / Materials and technology 45 (2011) 2, 115–120
Figure 1: Location of samples for ASB investigation
Slika 1: Mesto odvzema vzorcev za preiskave ASB
quasi-static and dynamic damage to the target material 25.
Therefore, the projectile generates compressive shock
waves on impact. These waves induce phase transfor-
mations and chemical changes, which can harden metals
to a great extent 25.
After the ballistic testing of the steel samples, some
microstructural changes were recorded. These changes
were mainly the formation of ASBs. A general view of a
typical ASB in the investigated samples is seen in Figure
3. One can see that the ASBs were formed in all samples
under the effect of AP projectile and appear as long
white strips. However, the density of the ASBs varied
significantly with respect to the hardness and thickness
of the samples. The microstructure of the deformed
ASBs in the specimen group A was very similar to that
of the matrix material. In addition, the extent of the ASB
formation was fairly low in comparison to the other
groups B and C. Furthermore, only deformed bands were
observed in the samples of group A and distortion in the
bands was rather low. On the other hand, in the specimen
groups B and C, both deformed and transformed bands
were found. Moreover, a higher distortion in the
deformed bands induced by shear in these specimen
groups was seen significantly in comparison to the
samples of group A. In our previous study 14, the
transformed band formation was also found for the steel
samples (AISI 4340 and DIN 100Cr6) with the hardness
of 49 HRC and 59 HRC even though the most of the
bands were of the deformed type in all the samples.
The thickness of the bands seen in the investigated
samples is given in Table 1. It changed between 19 μm
and 56 μm, depending on the types of band and the
sample. This indicates that the local deformation and
heating in the samples is taking place. Figure 4 repre-
sents the deformed band formed in the sample A2. It is
clear that the microstructure of the band was the same
with the base tempered martensitic structure. However,
the color of the martensite laths in the deformed band
was remarkably different from that of the matrix micro-
structure. This was probably due to the shear action
causing a different orientation of the martensite laths in
the band with respect to the base microstructure. Figure
5 illustrates the band formation in the sample B5. It is
apparent that both deformed and transformed bands
appeared in this sample. The transformed bands formed
as white lines and their microstructures were invisible
with an optical microscopy analysis as opposed to
deformed bands. A similar situation was also found in
the sample group C (Figure 6). The thinning and
T. DEMÝR, M. ÜBEYLÝ: A MICRO-DAMAGE INVESTIGATION ON A LOW-ALLOY STEEL ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 115–120 117
Figure 3: General view of ASBs in samples with various hardness and thickness
Slika 3: Splo{en videz ASB v vzorcih z razli~no trdoto in debelino
Figure 2: Microstructures of the sample AISI 4140 with a hardness of
a) 40 HRC, b) 60 HRC before testing
Slika 2: Mikrostruktura vzorca AISI 4140 s trdotama a) 40 HRC in b)
60 HRC pred preizkusom
shortening of the martensite laths in a transformed band
of C5 can be clearly seen in Figure 6b. Furthermore,
white areas of this band were resolved by SEM analysis.
Figure 7 depicts the detailed microstructure features for
this transformed band in the sample C5 steel after the
ballistic impact. In white regions of the band, fine
iron-carbide precipitates, consisting of a small amount of
chromium, were observed (Figure 7). The size of preci-
pitates changed between 100 nm and 5 μm. In
addition, fine grains were also observed in the band.
Figure 8 shows the elemental analysis for the fine
precipitate found in this band by EDX analysis. The
formation of very fine grains and precipitates and very
fine ferrite grains can be explained by the dynamic
recrystallization. The impact velocity and kinetic energy
of the projectile were sufficient to generate an immediate
local heating through very narrow bands in the
investigated steel AISI 4140. A ballistic study on the
30CrMnMo steel 10 also showed the formation of
different types of shear bands. The fine-grain and
T. DEMÝR, M. ÜBEYLÝ: A MICRO-DAMAGE INVESTIGATION ON A LOW-ALLOY STEEL ...
118 Materiali in tehnologije / Materials and technology 45 (2011) 2, 115–120
Figure 6: View of the band formation in the sample C5 after the
impact of the 7.62-mm AP projectile. a) 100-times, b) 500-times.
Slika 6: Formiranje pasov v vzorcu C5 po trku AP-krogle 7,62 mm:
pove~ave: a) 100-kratna, b) 500-kratna
Figure 4: Microstructure of the sample A2 after ballistic testing.
Deformed band formation with a lighter color compared to the base
microstructure (tempered martensite) was observed.
Slika 4: Mikrostruktura vzorca A2 po balisti~nem preizkusu. Opa`eno
je formiranje deformacijskih pasov s svetlej{o barvo v primerjavi z
osnovno mikrostrukturo (popu{~eni martenzit)
Table 1: Shear band widths of the samples
Tabela 1: [irina pasov v vzorcih
Sample ID Average BandWidth (μm) Band Type
A1 27
Deformed
A2 28
A3 36
A4 25
A5 42
B1 56
Deformed
B2 49
B3 34
B4 51
B5 49
B1 21
Transformed
B2 24
B3 25
B4 32
B5 36
C1 35
Deformed
C2 41
C3 25
C4 54
C5 40
C1 19
Transformed
C2 32
C3 25
C4 25
C5 35
Figure 5: View of deformed and transformed bands in the sample B5
after ballistic testing. Transformed band was observed as a white line.
Slika 5: Videz deformiranih in premenjenih pasov v vzorcu B5 po
balisti~nem preizkusu. Premenjeni pas je viden kot bela ~rta
carbide-precipitate formations in the shear bands of the
30CrMnMo steel, observed after the impact of the
tungsten-based projectile, were very similar to those
found in this study.
The hardness of the bands varied with respect to the
sample hardness and band type. Table 2 gives the micro-
hardness values of the bands for the investigated
specimens. It is clear that the hardness of the ASBs was
greater than that of the matrix for all the samples. The
difference between the hardness levels of the matrix
material and the ASB was smaller for the specimens of
group A since the microstructural coherency and smaller
shear between the band and the matrix material took
place. On the other hand, for the specimen groups of B
and C the variation in the hardness was much higher due
to the fact that the shear process occurred to a larger
extent. The hardness of the deformed band (344 HV) was
found to be slightly higher than that of the base material
(320 HV) for the specimens of group A. However, it
reached 519 HV and 695 HV for the sample groups B
and C, respectively. On the other hand, the hardness of
the transformed bands in the sample groups B and C was
recorded as 569 HV and 722 HV, respectively. The
increment in the hardness values was around 100 HV
and 150 HV for the deformed and transformed bands in
the sample group B, respectively. However, it was 30 HV
and 50 HV for the deformed and transformed bands in
the sample group C, respectively. The much higher
hardness of transformed bands was directly caused by
the fine structure of the grains and precipitates. A
previous study 24 indicated that when the hardness of the
AISI 4140 steel increased, the resistance to the advance-
ment of the projectile decreased. Although the specimens
belonging to the group A were perforated by a ductile
hole formation mechanism under the impact of the
7.62-mm AP projectiles, most of the specimens in the
groups of B and C were broken into several pieces in a
brittle manner rather than perforated 24. In other words,
the projectile perforated the samples in group A easily
due their lower hardness. Furthermore, the higher
resistance to projectile movement in the sample groups B
T. DEMÝR, M. ÜBEYLÝ: A MICRO-DAMAGE INVESTIGATION ON A LOW-ALLOY STEEL ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 115–120 119
Figure 9: Micro-view of the sample C5 after the ballistic testing. The
very intense ASB formation and the cracks are seen in the impact
zone.
Slika 9: Mikrovidez vzorca C5 po balisti~nem preizkusu. V zoni
udarca je nastalo veliko ASB-pasov in razpok.
Table 2: Shear bands hardness
Tabela 2: Trdota pasov
Sample
Group Types of Band
Hardness of
Band (HV)
Hardness of
Main Metal (HV)
A Deformed 344 320
B
Deformed 519
417
Transformed 569
C
Deformed 695
667
Transformed 722
Figure 8: Elemental analysis of the fine precipitate found in the
transformed band of C5 by EDX analysis.
Slika 8: Elementna EDX-analiza majhnega izlo~ka v prememnjenen
pasu vzorca (c/%)
Figure 7: SEM examination of a transformed band formed in the
sample C5. Formation of fine precipitate and grain occurred. a)
4000X, b) 14000X.
Slika 7: SEM-posnetki premenjenega pasa v vzorcu C5. Nastala so
majhna zrna in izlo~ki. Pove~ave: a) 4000-kratna, b) 14000-kratna
and C led to higher local heating in the steel, which
accelerated the formation of the ASBs. Both the micro-
structure and microhardness results for the AISI 4140
steel impacted by the AP projectile were found to be
similar with those of AISI 4340 and DIN 100Cr6 14 due
to the fact that their alloy contents are similar. All the
three steels have a strong carbide-former element of
chromium, which promotes the formation of carbide
precipitates in the transformed ASBs during the impact.
In contrast to the other two steels, the AISI 4340 steel
has another alloying element of nickel, which is not a
carbide forming element and so it does not affect the
precipitation sequence in the formation of the trans-
formed bands.
Borvik et al. 26 worked on the adiabatic shear forma-
tion in the Weldox 460E Steel impacted by blunt
projectiles. They found that the target thickness heavily
influenced the shear localization process. Only the
deformed bands ranging from 100 μm to 300 μm in
width were observed in target plates with a thickness of
less than 20 mm, while the transformed bands having
much less width (10–100 μm) were found for the thicker
target plates 26. Furthermore, they measured much higher
hardness values for the transformed bands (420 HV)
compared to the original hardness of the steel (180-190
HV) 26. In a more recent study 27, it was recorded that the
impact of the 7.62-mm AP projectile to the dual-phase
steel targets caused the formation of deformed bands
with higher hardness values. An increase of 70 to 100
HV in the hardness of the samples was measured after
testing due to the formation of these bands 27. In the
current study, it was also observed that an increase in the
thickness of the samples caused more ASB formation.
This is directly related to the increase in the resistance to
projectile motion with increasing thickness. The higher
portion of the kinetic energy of the projectile is trans-
ferred to the steel sample resulting in more ASB for-
mation.
ASBs could also lead to crack formation. In other
words, they can act as crack-initiation sites in the
deformed materials under high strain rates. Typical crack
sites occurred on the impact zone are represented for the
sample C5 in Figure 9. Very intense ASBs led to the
formation of cracks at the hit zone between the projectile
and the sample. The crack formation in the sample was
not detected outside the impact zone.
4 CONCLUSIONS
The main conclusions drawn from the experimental
results can be described as follows:
• A deformed band formation was observed in all the
steel specimens, which were tested with 7.62-mm AP
projectiles. However, transformed bands occurred in
the specimens having a hardness of 50 HRC or 60
HRC.
• The hardness of the bands was found to be remark-
ably higher than the matrix material.
• The impact action of the projectile led to the for-
mation of fine grains and precipitates in the trans-
formed bands.
• An increase in the resistance to the projectile motion
increased the transformed band formation.
• The effect of the projectile hit on the microstructure
of the AISI 4140 steel was found to be consistent
with that of the AISI 4340 steel 14.
Acknowledgement
This work was supported by the Research Fund of
TÜBÝTAK, Project # 106M211.
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T. DEMÝR, M. ÜBEYLÝ: A MICRO-DAMAGE INVESTIGATION ON A LOW-ALLOY STEEL ...
120 Materiali in tehnologije / Materials and technology 45 (2011) 2, 115–120
A. VESEL: ACTIVATION OF POLYMER POLYETHYLENE TEREPHTHALATE (PET) ...
ACTIVATION OF POLYMER POLYETHYLENE
TEREPHTHALATE (PET) BY EXPOSURE TO CO2 AND
O2 PLASMA
AKTIVACIJA POLIMERA POLIETILENTEREFTALATA (PET) S
CO2- ALI O2-PLAZMO
Alenka Vesel
Jo`ef Stefan Institute, Jamova 39, 1000 Ljubljana, Slovenia
Center of Excellence for Polymer Materials and Technologies, Tehnolo{ki park 24, 1000 Ljubljana, Slovenia
alenka.vesel@guest.arnes.si
Prejem rokopisa – received: 2010-11-04; sprejem za objavo – accepted for publication: 2011-01-26
A comparison of O2 and CO2 plasma treatment for the functionalization of PET (polyethylene terephthalate) is presented. The
plasma was created in a glass discharge chamber at a pressure of 75 Pa by an electrodeless RF discharge. The RF generator
operated at a frequency of 27.12 MHz and a power of about 200 W. The samples were cut into small pieces and exposed to
plasma for different periods. Immediately after the treatment the samples were characterized by high-resolution XPS. A
comparison of both survey and high-resolution C1s peaks revealed that the amount of the specific functional groups formed on
the surface during the plasma treatment was the same for the CO2 and O2 plasma. Within the limits of the experimental error the
concentration of hydroxyl groups was about 34 % and carboxyl groups was about 29 % for a sample treated in both plasmas for
30 s. The results were explained by the rapid dissociation of both molecules to neutral oxygen atoms that are fairly stable in
glass discharge tubes and readily react with the surface of polymer materials. Any effect of CO radicals is neglected since
oxygen atoms are chemically more reactive, so possible differences in the surface functionalization might have been observed
only for extremely short treatment times and/or orders of magnitude lower pressure.
Keywords: oxygen plasma, carbon dioxide plasma, PET polymer, surface functionalization, surface modification, XPS
Primerjali smo vpliv plazme, generirane v kisiku ali ogljikovem dioksidu, na funkcionalizacijo polimera PET (polietilen-
tereftalat). Plazmo smo ustvarili v stekleni razelektritveni cevi z RF-generatorjem s frekvenco 27,12 MHz in mo~jo 200 W. Tlak
plina v cevi je bil 75 Pa. Vzorce polimera smo razli~no dolgo izpostavili plazemski obdelavi, nato pa smo jih analizirali z
visokolo~ljivostnim XPS-spektrometrom. Iz primerjave preglednih in visoko lo~ljivih XPS-spektrov smo ugotovili, da je dele`
novo nastalih funkcionalnih skupin enak za obe plazmi: O2 in CO2. V obeh primerih je bila koncentracija hidroksilnih skupin
okoli 34-odstotna, koncentracija karboksilnih skupin pa okoli 29-odstotna. Podobnost obeh plazem smo razlo`ili z mo~no
disociacijo obeh molekul O2 in CO2 na nevtralne atome kisika, ki so v razelektritveni cevi dokaj stabilni in zaradi svoje
reaktivnosti reagirajo s povr{ino polimernega materiala. Vpliv radikala CO, ki nastaja pri disociaciji molekule CO2 smo
zanemarili, ker so kisikovi atomi veliko bolj reaktivni kot CO.
Klju~ne besede: kisikova plazma, plazma ogljikovega dioksida, polimer PET, funkcionalizacija povr{ine, povr{inska
modifikacija, XPS
1 INTRODUCTION
Low-pressure weakly ionized plasma is a popular
tool for the modification of the surface properties of
solid materials. It is often regarded as an ecologically
benign alternative to wet chemical processing. The major
effect of the plasma surface interaction is the potential
interaction within the excited plasma particles and the
atoms on the surface of the solid material. The result of
the interaction is either a reduction or oxidation of the
material surface.1–12 Accordingly, plasmas created in
different gases are used to obtain certain effects. For
reduction proposes, a hydrogen plasma is often applied.
The major technology based on the interaction of a
hydrogen plasma with solid materials is discharge
cleaning, i.e., the removal of oxidizing impurities from
the surface of the materials.1–5 A more modern tech-
nology based on the application of hydrogen plasma is
the synthesis and/or modification of nanomaterials.6–9
Oxidation, on the other hand, is performed by a
plasma created in various gases.10–12 Oxygen, nitrogen,
water vapour and carbon dioxide are gasses suitable for
the generation of plasma with oxidizing particles. The
technologies based on the application of oxidizing
plasma include discharge cleaning (in this case de-
greasing),13,14 plasma etching15–17, plasma sterilization18–19
and the plasma synthesis of metal oxide or nitride nano-
particles.20–23 Another technology of particular import-
ance is the surface activation of organic materials.24–31
Generally speaking, any organic material can be
functionalized by required functional groups using a
plasma created in an appropriate gas. In practise,
however, the type and concentration of specific
functional groups created on a surface of specific organic
material, is limited. In particular, the functionalization of
polymer materials using oxygen plasma often leads to
the appearance of at least three different functional
groups (like hydroxyl, carbonyl and carboxyl). Although
many attempts have been made to control the concen-
tration of each functional group on the surface of organic
materials, the results are far from being satisfactory. In
Materiali in tehnologije / Materials and technology 45 (2011) 2, 121–124 121
UDK 533:678 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 45(2)121(2011)
order to avoid the formation of all possible oxygen-rich
functional groups on the surface of the polymer, the
application of a plasma created in different gases was
suggested. For instance, several attempts have been made
to functionalize polymer PET (polyethylene tereph-
thalate) with carboxylic groups using a plasma created in
carbon dioxide instead of oxygen.32,33 The aim of this
paper is a comparison of the plasma created by the same
discharge in the same plasma vessel but in two different
gases: oxygen and carbon dioxide. The appearance of
different functional groups was monitored by high-
resolution XPS (X-ray photoelectron spectroscopy).
2 EXPERIMENTAL
2.1 Plasma treatment of polymer
Experiments were performed with a polyethylene
terephthalate (PET) foil from DuPont. The samples were
treated in the experimental system which was pumped
with a two-stage oil rotary pump with a pumping speed
of 4.4 · 10–3 m3 s–1. The discharge chamber was a Pyrex
cylinder with a length of 0.6 m and an inner diameter of
0.036 m. The plasma was created with an inductively
coupled RF generator, operating at a frequency of 27.12
MHz and an output power of about 200 W. Commer-
cially available oxygen or carbon dioxide was leaked
into the discharge chamber. The pressure was measured
by an absolute vacuum gauge. During our experiments,
the pressure was fixed at 75 Pa. The samples of PET foil
were treated in O2 or CO2 plasma for 10 s and 30 s.
2.2 X-ray photoelectron spectroscopy (XPS) characte-
rization
The surface of the plasma-treated PET samples was
analyzed with an XPS (X-ray Photoelectron Spectro-
meter) instrument TFA XPS Physical Electronics. The
base pressure in the XPS analysis chamber was about 6 ·
10–8 Pa. The samples were excited with X-rays over a
400-μm spot area with a monochromatic Al Kα1,2
radiation at 1486.6 eV. The photoelectrons were detected
with a hemispherical analyzer positioned at an angle of
45° with respect to the normal to the sample surface.
Survey-scan spectra were made at a pass energy of
187.85 eV and a 0.4-eV energy step, while for C1s
individual high-resolution spectra were taken at a pass
energy of 23.5 eV and a 0.1-eV energy step. Since the
samples are insulators, we used an additional electron
gun to allow for surface neutralization during the
measurements. The spectra were fitted using MultiPak
v7.3.1 software from Physical Electronics, which was
supplied with the spectrometer. The curves were fitted
with symmetrical Gauss-Lorentz functions. A Shirley-
type background subtraction was used. Both the relative
peak positions and the relative peak widths (FWHM)
were fixed in the curve-fitting process.
3 RESULTS AND DISCUSSION
Samples were exposed to oxygen or carbon dioxide
plasma for 10 s and 30 s. The time of 10 s is a typical
treatment time that ensures the saturation of the surface
with oxygen-rich functional groups. Figure 1 represents
the XPS survey spectra of an untreated sample and
samples treated with O2 and CO2 plasma. We can
observe qualitatively that the concentration of carbon is
decreased in favor of oxygen for both plasma-treated
samples. The quantitative results are summarized in
Table 1. The experiments were repeated several times in
order to minimize any statistical errors so the values
presented in Table 1 are averaged over several measure-
ments. The statistical error is within 1 %, indicating
fairly reproducible results. It is interesting that the
concentration of carbon and oxygen is practically the
same for samples treated with oxygen and carbon
dioxide plasma. Furthermore, the differences between
the 10 s and 30 s treatment times are minimal. This is
just another confirmation of the well-known fact that
saturation of the surface with functional groups occurs
before 10 s of treatment time.34,35
Table 1: Comparison of the surface composition of the PET polymer
treated in oxygen or carbon dioxide plasma.
Tabela 1: Primerjava povr{inske sestave polimera PET, obdelanega v
kisikovi plazmi ali plazmi ogljikovega dioksida
Sample C O
untreated sample 73.1 26.9
CO2 plasma – 10 s 55.9 44.1
O2 plasma – 10 s 56.5 43.5
CO2 plasma – 30 s 55.9 44.2
O2 plasma – 30 s 56.9 43.2
The survey spectra presented in Figure 1 give us
information about the concentration of elements in the
A. VESEL: ACTIVATION OF POLYMER POLYETHYLENE TEREPHTHALATE (PET) ...
122 Materiali in tehnologije / Materials and technology 45 (2011) 2, 121–124
Figure 1: Comparison of XPS survey spectra of untreated PET sample
and of PET sample treated in O2 or CO2 plasma for 10 s
Slika 1: Primerjava preglednega XPS-spektra neobdelanega vzorca
polimera PET ter vzorca polimera PET, obdelanega 10 s v O2- ali
CO2-plazmi
surface film but do not tell us anything about the
concentration and the type of each particular functional
group. In order to get an insight into the concentration of
the functional groups, we performed high-resolution
XPS measurements of the carbon C1s peak. Again, we
performed the analysis on several samples, but a typical
result is presented in Figure 2. The major peak at 285
eV corresponds to the –C=C bond, the peak at 286.3 eV
to the C-O(H) bond, while the well-pronounced peak at
288.8 eV corresponds to –COO- (carboxyl and ester
group). As expected, from the knowledge gained from
Figure 1 and Table 1, the concentration of hydroxyl
C-OH and carboxyl functional groups -COOH is
increased dramatically. Interesting, however, no diffe-
rence is observed between the sample treated in O2 and
CO2 plasma. Table 2 represents a quantification of the
results presented in Figure 2. Again, the values pre-
sented in Table 2 are averaged over several samples. It is
interesting that the concentration of different functional
groups on all the samples is practically the same, or
definitely within the limit of the experimental error.
From these results we can conclude that both O2 and CO2
plasma treatments lead to the formation of practically the
same functional groups.
The upper result is explained by taking into account
the characteristics of the oxygen and carbon dioxide
plasmas. Since the ionization fraction in both plasmas is
practically the same, and is of the order 10–6, the charged
particles play a minor role in the surface modification of
our sample. The dissociation fraction, on the other hand,
is at least five orders of magnitude larger than the
ionization fraction. Such a huge difference between the
ionization and dissociation fraction is explained by the
probabilities of neutralization and recombination. The
probability for the surface neutralization of charged
particles does not depend on the type of the material
facing plasma and is very close to 1. The surface
recombination probability, on the other hand, is very
sensitive to the surface properties and may be anything
between 10–6 and 1.36-37 In the case of glass discharge
chambers, the probability for surface recombination is
fairly low, since oxygen atoms do not chemisorb on the
glass surface. In both cases, the result of the dissociation
is atomic oxygen. Since the dissociation energies for
oxygen molecules and CO2 molecules are similar, it is
expected that the dissociation probability would be
practically the same in both plasma. Moreover, the
surface recombination probabilities are also similar for
both gases. Taking into account these considerations we
can explain the observed functionalization of the
polymer. In both O2 and CO2 plasmas the major reactants
are neutral oxygen atoms. There is no reason that the
atoms originating from oxygen molecules or carbon
dioxide molecules would act differently. Any differences
might have appeared at much lower treatment times, i.e.,
well before the saturation of the surface with functional
groups is observed. Such experiments, however, are not
possible in our labs since we do not have a pulsed
plasma generator.
4 CONCLUSION
Well-defined foils of polymer PET were exposed to
plasma created in oxygen and carbon dioxide gases in
order to study any possible differences in the surface
functional groups created during plasma processing.
High-resolution XPS was applied to study the type and
concentration of different functional groups. Within the
limits of the experimental error, we can clearly conclude
that there are no differences in the oxygen functional
groups between the treatment with oxygen and the
carbon dioxide plasmas. Our results were explained by
the formation of a large density of neutral oxygen atoms
in both plasmas, since these atoms are the major
reactance in low-pressure, weakly ionized, highly
reactive plasma.
ACKNOWLEDGEMENT
This project was supported by Slovenian Research
Agency, Project L2-2047. The authors also acknowledge
A. VESEL: ACTIVATION OF POLYMER POLYETHYLENE TEREPHTHALATE (PET) ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 121–124 123
Figure 2: Comparison of high-resolution XPS spectra of carbon C1s
of untreated PET sample and of PET sample treated in O2 or CO2
plasma for 10 s
Slika 2: Primerjava visoko lo~ljivega XPS-spektra ogljika C1s
neobdelanega vzorca polimera PET ter vzorca polimera PET,
obdelanega 10 s v O2- ali CO2-plazmi
Table 2: Comparison of the concentration of different functional
groups formed on PET polymer during treatment in oxygen or carbon
dioxide plasma.
Tabela 2: Primerjava koncentracije razli~nih funkcionalnih skupin,
nastalih na povr{ini polimera PET med obdelavo v kisikovi plazmi in
plazmi ogljikovega dioksida
Sample -C=C C-O O=C-O
untreated sample 75.8 13.0 11.2
CO2 plasma – 10 s 42.4 32.3 25.3
O2 plasma – 10 s 42.2 31.7 26.1
CO2 plasma – 30 s 38.7 33.1 28.3
O2 plasma – 30 s 36.5 34.3 29.2
the financial support from the Ministry of Higher
Education, Science and Technology of the Republic of
Slovenia through the contract No. 3211-10-000057
(Center of Excellence Polymer Materials and Techno-
logies).
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124 Materiali in tehnologije / Materials and technology 45 (2011) 2, 121–124
A. VESEL: ACTIVATION OF POLYMER POLYETHYLENE TEREPHTHALATE (PET) ...
C. B. FOKAM et al.: THE IMPACT ON RIGID PVC PIPES: A STUDY OF THE CORRELATION ...
THE IMPACT ON RIGID PVC PIPES: A STUDY OF THE
CORRELATION BETWEEN THE LENGTH OF THE
CRAZED ZONE AND THE AREA OF THE IMPACTED
REGION
UDAR TOGIH PVC-CEVI: [TUDIJA KORELACIJE MED DOL@INO
RAZPOKANE ZONE IN POVR[INO ZONE UDARA
Christian Bopda Fokam1, Mohamed Chergui1, Kalifa Mansouri2,
Mohamed El Ghorba1, Mohamed Mazouzi1
1Laboratoire de Contrôle et de Caractérisation des Matériaux, B.P 8118, Oasis-Route El Jadida – ENSEM / Casablanca, Maroc
2Ecole Normale Supérieure de l’Enseignement Technique, Bd hassan II, Mohamedia Maroc
fokam79@yahoo.fr
Prejem rokopisa – received: 2010-07-04; sprejem za objavo – accepted for publication: 2011-01-110
This paper presents a numerical methodology aimed at identifying quantitatively the crazing under quasi-static impact loading
of rigid PVC pipes. It also presents the correlation between the length of the crazed region and the area of the impact. Testing
low-speed impacts has provided a damaging loading before the rupture of the sample’s pipe section. The striker impact on the
pipes allows us to appreciate the damage caused. The craze in the impacted area was identified in our previous work 1, the
different criteria of the craze initiation are presented. The criterion of Steinstern and Myers, better adapted to rigid PVC 2, is
used in a numerical simulation on ANSYS for the problems of static contact. For tests carried out at low velocity, an equivalent
static load representing the dynamic impact loading can account for the stress states in the structure of the pipe 3. The analysis of
the stress fields in the structure of the pipe compared to the stress of the craze initiation depends on the first stress invariant and
allows us to deduce the length of the crazed region in the thickness of the pipe. Finally, we present the correlation between the
length of the crazed region and the area of the zone of impact (the area of the imprint of the striker on the pipe).
Keywords: damage, crazing, impact, PVC, amorphous polymer
Predstavljena sta numeri~na metodologija za kvantitativno identifikacijo razpokanja pri kvazi stati~ni udarni obremenitvi togih
PVC-cevi in korelacija med dol`ino z razpokane zone in povr{ino udara. Preizkusni udari z majhno hitrostjo so ustvarili
po{kodbeno obremenitev pred zlomom prereza cevi. Udarec cevi s kladivom omogo~a, da se oceni nastala po{kodba.
Razpokanje zone udara je identificirano v na{em prej{njem delu 1, zato so predstavljena merila za~etka razpokanja. Merilo
Steinsteina in Myersa, ki je bolj prilagojen za toge PVC-cevi 2, je uporabljen pri numeri~ni simulaciji ANSYS za problem
stati~nega stika. Za preizkuse pri majhni hitrosti lahko ekvivalentna stati~na obremenitev pomeni dinami~no obremenitev in
omogo~i, da se opredeli napetostno stanje v strukturi cevi 3. Analiza napetostnih polj v strukturi cevi, primerjana z napetostjo
za~etka razpokanja, je odvisna od invariante prve napetosti in mogo~a, da se dolo~i globina dol`ine zone razpoke v steni cevi.
Kon~no je predstavljena tudi korelacija med dol`ino razpokane zone in povr{ino zone udara ( povr{ina odtisa kladiva).
Klju~ne besede: po{kodbe, razpokanje, udar, PVC, amorfni polimer
1 INTRODUCTION
Plastic materials are a large proportion of pipe
water-supply systems primarily due to the ease of
installation and the relatively low cost of the materials.
Plastic pipes used for water canalization continues to be
the subject of many studies focusing mainly on the
behavioral nature of the material. During construction
works, pipes are often subjected to accidental impact
such as falling rocks, scratches etc. The ability to resist
pressure and the reliability becomes important to
engineers. The approach generally used to solve this
problem requires the identification and determination of
the length of the morphological degradation mechanism
of the material during impact.
Amongst the morphological mechanisms identified
as responsible for the degradation of amorphous
polymers, we find mostly crazes and shear bands 4.
The identification of a craze in the area of the impact
has already been demonstrated in a previous study 1. It is
important to define the critical length of the crazed
region, beyond which the material can be considered as
either obsolete or unusable.
This article focuses mainly on the presentation of a
numerical methodology for the determination of the
length of a crazed area. A correlation between the area of
the impact on the pipe (imprint) and the length of the
crazed region is also studied. To illustrate this study, we
first define a craze and the different initiation criteria.
Subsequently, the material and the dimensions of the test
specimens are presented, as well as the experimental
impact conditions. Finally, steps in the numerical
methodology and the analyses for the determination of
the length of the crazed region are exposed.
2 CRAZING
2.1 Presentation
Crazing is a defect resulting in the appearance of
cracks on the surface of a material. Crazes are super-
ficially similar to cracking.
Materiali in tehnologije / Materials and technology 45 (2011) 2, 125–130 125
UDK 532.5:628.1:539.4:678.743 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 45(2)125(2011)
However, there are differences between them. Crazes
are characterized by crack partitions linked together by
small parallel fibrils of the material 5 (see Figure 1a).
Crazes and cracks proceed in three stages. Initiation,
which is the enlargement and rupture of fibrils. Crazing
involves a process of stretching the bulk of the material
towards the area were the fibrils are located at the
beginning of the crack 8.
2.2 Criteria of craze initiation
In elastic and visco-elastic regimes or after the
development of a significant plastic deformation, the
corresponding formula of the initiation criteria is distinct
for each regime. In a previous publication 1, we detected
the first craze in elasto/visco-elastic regime for rigid
PVC. The initiation criteria presented below, correspond
to the craze appearances in elasto/visco-elasto regimes
and elaborate on different phenomenological bases.
Following the observation of an incubation time for
the craze formation where the constant stress is less than
half of the yield stress (	y), Argon et al. 9 proposed a
sophisticated formula (1), as shown below
	 	
	1 2 13 2
− =
+
A
C I Qy/
(1)
where A and C represent material parameters. 	1 and 	2,
respectively, are the maximum and minimum principal
stress. I1 is the first stress invariant, Q = 0.0133 is a
factor controlling the dependency of the criterion to the
shear stress on I1 and 	y is the yield stress.
Sternstein and Myers 10 proposed a criterion based on
the mechanism of micro-void formation in a dilatational
stress field, and on the stabilization of micro voids
through a deviatoric stress (2). Recently, following the
same approach, Gearing 11 proposed a criterion based
only on the maximum principal stress (3):
	 	1 2
0
0
1
− ≥ +A
B
I
(2)
	1
1
≥ +A
B
I
(3)
With the same idea, Oxborough and Bowden sugge-
sted the definition of a criterion based on critical defor-
mation. The criterion has been reformulated in the stress
for an elastic material, with  standing for the Poisson’s
ratio and E for the Young’s modulus (X' = EX and Y’ =
EY are two material parameters depending on the
temperature) (4):
	 	 	1 2 3
1
− − ≥ +v v X
Y
I
'
'
(4)
To account for the sensitivity to hydrostatic pressure
observed for some polymers, Ishikawa 12 proposed the
following criterion (5):
I A
B
I1 1
≥ + (5)
3 EXPERIMENTAL PROGRAM
3.1 Material
The pipes used are mostly made up of polyvinyl
chloride (rigid PVC). They come from the same
C. B. FOKAM et al.: THE IMPACT ON RIGID PVC PIPES: A STUDY OF THE CORRELATION ...
126 Materiali in tehnologije / Materials and technology 45 (2011) 2, 125–130
Figure 2: Dimensions (in millimeter) of samples of the test tube for
the impact test
Slika 2: Dimenzije vzorcev (v milimetrih) preizkusne cevi za udarne
preizkuse
Table 1: Mechanical properties of rigid PVC
Tabela 1: Mehanske lastnosti togega PVC
Properties Values
Young’s modulus 3200 MPa
Tensile stress fracture 75 MPa
Elongation at fracture tensile 100 %
Tensile strength in compression 50–75 MPa
Figure 1: (a) Diagram of a craze with cracks characterized by the pre-
sence of fibrils 6 (b) Craze region with the result of fissure formation
7. Zone A represents the region were the glassy polymer is elastically
deformed. B, the area were initiation crazing occurs. C propagation or
growth of the crazed region. D Rupture of the fibrils leading to the
transformation of the craze into a crack of length a.
Slika 1: (a) Diagram po{kovanja, ki ga karakterizirajo fibrili 6. (b)
Zona risov z rezultati nastanka razpok 7. Zona A je podro~je, kjer je
bil steklast polimer elasti~no deformiran. B je povr{ina nastanka risov.
D je prelom fibrilov, ki rise spremeni v razpoko z dol`ino a.
molding. The pipe dimensions are 63 mm in diameter
and a wall with a 4.5 mm thickness (see Figure 2).
3.2 Presentation of the impact machine
This device allows us to drop a weight (striker) from
a selected height onto a pipe (see Figure 3a and 3b).
The section of the pipe is placed on a support "V". A
cylindrical rigid bar to help prevent the rebound of the
striker during and after the impact is introduced. A
weight of the striker equal to 16 kg and of radius R = 50
mm is used. (See Figure 3c). The height used ranges
between 0 m and 2 m. The mass, height and velocity are
used to calculate the kinetic energy of the impact.
After impact, the object dropped of mass 16 kg,
leaves an elliptical mark called the impacted area or the
area of impact at the point of impact (see Figure 4).
4 METHODOLOGY FOR DETERMINING THE
LENGTH OF THE CRAZED REGION
In a previous publication 1 we have identified the
crazing as the morphological mechanism of damage in
the impacted area (see Figure 4). In this section, we
propose a numerical methodology for determining the
length of the crazed region in the pipe wall thickness
below the impacted area (see Figure 4).
The craze initiating to the critical stress (in agree-
ment with the criterion of craze initiation considered)
and the methodology consist of delimiting areas in the
pipe were the stress levels are greater than the critical
stress (	  	 cr
am , 	 cr
am : critical stress initiation).
The different criteria for the craze initiation presented
above (Section1.2) are valid, depending on the type of
material. An analysis of these criteria showed that the
criterion of Sternstein and Myers (2) is the best adapted
for rigid PVC.
The determination the length of the crazed region
requires the calculation of stress fields in the wall
thickness of the pipe material at the time of impact.
These experiments are difficult. We have constructed a
numerical model based on Ansys for the simulation of
the problem of impact dynamics 13. As our test impacts
are at low velocities (V < 6 m/s), it is shown in the
literature that a static equivalence of a dynamic problem
is sufficient to account for the stress in the structure 3.
Therefore, the impact of the striker on the pipe is treated
as a static contact problem between two solid bodies.
The normal load applied on the upper surface of the
striker (see Figure 5) is obtained through the equations
of Hertz, linking the dimensions of the impacted area
(see Figure 4) to the static normal force. A visco-elastic
constitutive law is applied to the numerical model in
order to reduce the computation times of the problem,
which is already highly non-linear, without altering the
objectives of the study. Indeed, the craze initiating for
the case of rigid PVC in the elastic regime and the
critical stress initiation are expected to lower the yield
stress.
The high stress levels in the material (stress greater
than the resistance of the material studied) do not affect
the final results of the experiment.
C. B. FOKAM et al.: THE IMPACT ON RIGID PVC PIPES: A STUDY OF THE CORRELATION ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 125–130 127
Figure 5: Geometric representation of the numerical model: (a) full
actual geometry, (b) ¼ of the geometry implemented due to the
symmetry of the problem.
Slika 5: Geometri~na slika numeri~nega modela: (a) polna realna
geometrija, (b) ¼ geometrije, uporabljene zaradi simetrije problema
Figure 4: Photographs of the impacted area on sections of the pipe at
varying heights after the drop of the striker (mass)
Slika 4: Posnetka povr{ine udara na prerezih cevi po razli~ni vi{ini
padca udarnega kladiva
Figure 3: (a) Impact machine from the German TestSystems UTS, (b)
synoptic diagram of the impact machine, (c) dimensions of the hemi-
spherical striker
Slika 3: (a) Udarna naprava nem{kega proizvajalca TestSystems UTS,
(b) sinopti~en diagram udarne naprave, (c) dimenzije hemisferi~nega
kladiva
The tests impacts, being numerous, means that the
details of the methodology for determining the length of
the crazed region in the pipe wall thickness will be
presented for an impact using a striker of mass 16 kg
dropped from a height of 1.25 m (designated impact:
1.25 m/16 kg).
	S&M, the stress parameter of the pipe to be found,
corresponds to the criterion of Sternstein and Myers (2):
	 	 	 	S& M cr
am= − ≥ = +1 2
0
0
1
A
B
I
(6)
where 	1 > 	2 > 	3, A0 = 19.936 MPa and B0 = 1 203.384
MPa 2 for rigid PVC.
Figure 7, represents a cartograph stress state in the
wall thickness of the pipe. Around each isovalue of a
contour line with the stress represented by we observe
that from the given equation above, the critical stress of
Sternstein and Myers (	 cr
am , see formula (6)) varies
weakly with I1 (see Figure 6). Moreover, it seems
reasonable that the local critical stress for the craze
initiation decreases as I1 increases. Thus, the creation of
micro-voids before crazing would be easier and would
lead to a lower stress initiation. This reflects in a
decrease in the critical stress while "I1" increases.
In Figure 6 we see that the variation of is not very
significant and in all cases it does not fluctuate above
and below 10 % of its average value. Thus, the critical
stress 	 cr
am at all points on an isovalue contour of the
stress '	S&M' is considered constant.
Each isovalue contour of the stress in Figure 7b is a
possible limit between the crazed regions and others not
crazed in the thickness of the pipe. This limit is
identified when. We also notice that all the isovalue
contours cover the entire section of the thickness of the
pipe (see Figure 7b). Thus, whatever the limit of the
isovalue contour, crazing occupies the entire thickness of
the pipe, where the stress fields represent (with being
constant on an isovalue contour).
Considering as a constant on an isovalue contour, the
representation of the stress states compared to critical
stress initiation along an imaginary line (see Figure 7b)
is sufficient to determine the length of the crazed region,
which corresponds to .
Figure 8, represents the stress states compared to
along an imaginary line located at any position, crossing
the right section of the structure of the pipe (S1) just
below the contact point (see Figure 7).
In Figure 8, the intersection between the plot and
indicates the limit of the crazed region. Beyond this
intersection, the stress is less than the stress initiation.
The stress state is abnormally large and the structure can
be attributed to the elastic constitutive law applied to the
numerical model. Nevertheless, the intersection point is
C. B. FOKAM et al.: THE IMPACT ON RIGID PVC PIPES: A STUDY OF THE CORRELATION ...
128 Materiali in tehnologije / Materials and technology 45 (2011) 2, 125–130
Figure 8: A Graph of the critical stress for craze initiation against the
stress state of the material subjected to loading, along an imaginary
line selected in the pipe.
Slika 8: Grafikon kriti~ne napetosti za nastanek po{kodbe 	cram proti
napetosti 	S&M za obremenjeni material po imaginarni ~rti v cevi.
Figure 6: A graph of the critical stress of the craze initiation of the
Sternstein and Myers criterion 	cram = A0 + B0/I1 around an isovalue
contour in the thickness of the pipe.
Slika 6: Graf kriti~ne napetosti za~etka po{kodbe po Sterstein-
Myersovem merilu 	cram = A0 + B0/I1 za enake vrednosti po obodu
cevi
Figure 7: Cartograph of stress states 	S&M =	1 – 	2 corresponding
to the criterion of Sternstein and Myers. (a) ¼ of the geometry of the
pipe, (b) surface S1 of the thickness of the pipe
Slika 7: Kartografija napetostnega stanja 	S&M =	1 – 	2
where the stress 51 MPa 
 	y = 52 MPa (	y: yield
strength in tensile of the rigid PVC).
In Figure 8, the half-length "ac" of the crazed region,
of the stress state corresponding to the test impact using,
respectively, a height and mass of 1.25 m and 16kg is:
7.62 mm (ac = Lc / 2 : half the length of the crazed zone).
Table 2 summarizes the determined size of the half
lengths of the crazed regions corresponding to the
different impact tests performed (h  [0.2 m 1.7 m]).
Table 2: Half-length of the crazed region "ac" corresponding to the
impacts of mass 16kg dropped from different heights
Tabela 2: Polovica dol`ine po{kodovane zone "ac", ki ustreza udaru
mase 16 kg spu{~ene z razli~ne vi{ine
Designation
(m/kg)
h a/mm ac/mm
0.2 m/16 kg 0.2 10 2
0.5 m/16 kg 0.5 12.5 6.45
1 m/16 kg 1 13 7.22
1.25 m/16 kg 1.25 13.75 7.62
1.5 m/16 kg 1.5 14.5 8.49
1.7 m/16 kg 1.7 aMax = 15.5 acMax = 9.55
From Table 2, we notice that the dimensions of the
various areas of impact "a" are higher than the
corresponding crazed "ac". It is true that bleaching the
impacted area (see Figure 4a) allows the localization of
the morphological degrading mechanisms (damage
caused by the impact).However, it does not help to
determine its length. But the developments of these two
parameters show correlations (see Figure 9).
In the range of the impact tests performed (h  [0.2
m, 1.7 m]), the evolutions of "ac" and "a" show an almost
linear growth with the level of impact (see Figure 9a). In
Figure 9b, the correlative evolution shows a linear
dependence of these two quantities. The area of the
imprint (impacted area) may be a good indicator of the
length of the crazed region in a pipe subjected to an
impact.
5 CONCLUSION
A numerical methodology for determining the length
of the crazed region in a rigid PVC pipe subjected to a
quasi-static impact has been presented.
With crazing identified on the impacted area 1, deter-
mining its length with the methodology proposed
involves a comparison of the stress states in the structure
for the critical stress of the craze initiation. After
choosing the criterion of initiation of Sternstein and
Myers (6) for rigid PVC 2 we notice that its dependence
on the first stress invariant I1 induces some distinct
critical values. However, on an isovalue contour of
stress, the criterion does not deviate more than ±10 %
from the average value (see Figure 6). From this small
variation we considered the value of the critical initiation
as a constant on an isovalue contour. Thus, the
comparative representation of the stress in the wall
thickness of the pipe 	S&M and the critical stress initiation
craze 	 cr
am (see Figure 8) along a line crossing all the
boundaries of all the isovalue contours (see Figure 7b)
allow us to deduce the length of the crazed region
corresponding to 	S&M  	 cr
am . Depending on the level of
the impact load, we see a linear growth of the length of
the crazed region (see Figure 9a), which is linearly
correlated to the size of the corresponding impacted area
(see Figure 9b).
6 REFERENCES
1 Fokam Bopda C., Chergui M., Ghorba M., K. Mansouri Impact des
tubes en PVC rigide : tenu mécanique et mécanisme de détérioration,
Communication, Congrès Français de Mécanique (Marseille), ref:
N°1313- S18 /CFM’2009 (2009)
2 Fokam Bopda C., Chergui M., Ghorba M., K. Mansouri Discri-
mination des critères d’amorçage de craquelures: PVC rigide (non
plastifié), Communication, RNJCP 2009, Faculté, Mbem Sik
(Casablanca) (2009)
3 Y. Aminanda, B. Castanié, Modélisation de l’indentation des
structures sandwichs à peaux métalliques, Mécanique & Industries 6
(2005), 487–498 DOI: 10.1051/meca:2005061
4 Van der Giessen E, Lai J Proc. 10th Int. Conf. On Deformation, Yield
and Fracture of Polymers (1997)
5 Tijssens MGA, Van der Giessen E, Sluys LJ (2000) Mech Mat 32:19
6 Kausch, H.-H., Polymer fracture, 2nd ed., Springer-verlag, Berlin
(1987)
7 C. G’Sell, J. M. Haudin introduction à la mécanique des polymères,
Editors INPL, Nancy, (1995), vol. 1.
C. B. FOKAM et al.: THE IMPACT ON RIGID PVC PIPES: A STUDY OF THE CORRELATION ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 125–130 129
Figure 9: (a) Evolution comparison of crazed region "ac" and im-
pacted area "a" depending on the level of impact. (b) Evolution of the
corresponding non-dimensional "ac/acMax" by to "a/amax"
Slika 9: (a) Evolucija primerjanega po{kodovanega podro~ja "ac'" in
podro~ja udara v odvisnosti od velikosti udara; (b) ustrezne
brezdimenzijske vrednosti "ac/acmax" in "a/amax"
8 S. K. M. Ting, J. G. Williams, A. Ivankovic, Engineering Fracture
Mechanics, Fracture and Damage Mechanics, 71 (2004) 4–6,
657–668
9 Argon AS, Hannoosh JG, Phil Mag 36 (1977), 1195
10 Sternstein SS, Myers FA, J Macromol Sci –Phys B 8 (1973), 539
11 Gearing B.P., Anand L. On modelling the deformation and fracture
reponse of glassy polymers due to shear yielding and crazing. Int J
Sol. Struct. 41 (2004), 3125–3150
12 Ishikawa M., Sato Y., Higuchi H. Polymer, 37 (1996) 7, 1177–1181
13 Fokam Bopda Christian, Chergui M., Ghorba M., K. Mansouri
(2009) Impact tube en PVC rigide, Démarche de validation d’un
modèle numérique revue ClicEdition, livre: Les avancées en
Mécanique, Tendance de la Modélisation ET Nouvelles techniques
expérimentales (Algérie, Bikskra) Ref: 0165-C04/ CAM2009
C. B. FOKAM et al.: THE IMPACT ON RIGID PVC PIPES: A STUDY OF THE CORRELATION ...
130 Materiali in tehnologije / Materials and technology 45 (2011) 2, 125–130
D. TOLMAC et al.: A COMPARATIVE ANALYSIS OF THEORETICAL MODELS AND EXPERIMENTAL RESEARCH ...
A COMPARATIVE ANALYSIS OF THEORETICAL
MODELS AND EXPERIMENTAL RESEARCH FOR
SPRAY DRYING
PRIMERJALNA ANALIZA TEORETI^NIH MODELOV IN
EKSPERIMENTALNA RAZISKAVA RAZPR[ILNEGA SU[ENJA
Dragisa Tolmac, Slavica Prvulovic, Dragana Dimitrijevic, Jasna Tolmac
University of Novi Sad, Technical faculty "Mihajlo Pupin", Djure Djakovica bb, 23000 Zrenjanin, Serbia
dragisat@gmail.com
Prejem rokopisa – received: 2010-09-14; sprejem za objavo – accepted for publication: 2011-02-24
Spraying is very often applied in the chemical and food-processing industries. It includes the spraying of a suspension –
spraying with a content of about 30 % to 65 % of a dry substance. The drying time is in most cases only several seconds. In this
work the convective drying principle in warm air is presented.
A mathematical model of momentum, heat-and-mass transfer in the atomization-swirler zone was proposed. An uneven
distribution of particles and entrainment effects were taken into account in the model. To verify the model an extensive
experimental investigation was performed on water evaporation at different initial air temperatures, feed rates, flow rates of the
drying agent for different parameters of the atomization and initial particle size distribution. Changes in the air temperature
inside the stream of sprayed material, material temperature, evaporation rate, and changes in the Sauter diameter and the
distribution of particle diameter vs. their distance from the atomizer were determined experimentally. Good agreement between
the experimental results and the theoretical data was achieved.
In the paper an attempt was made to apply the model to the calculation of spray drying in a pilot-plant dryer. The air supplied to
the dryer tangentially to the axis was characterized by a high swirl. In order to consider the swirl air flow pattern in the dryer,
changes in heat-and-mass balance were made for the particle-drying agent system. The model was verified experimentally on
the basis of results of investigations on the drying of a 55 % mass solution of starch and water.
It was found that our model described the process of drying in the analyzed system. An extensive literature survey on spray
drying, including unsteady-state phenomena, is presented in the paper. The main sources of errors occurring in the modeling of
the spray drying process were discussed. A critical estimation of the existing mathematical models of the process was made and
some of them were described in detail.
Key words: spray dryer, model, evaporation, temperature distribution
Razpr{evanje je pogosto uporabljeno v prehrambni industriji in obsega razpr{evanje emulzije s 30–65 % suhe snovi. ^as su{enja
je najve~krat nekaj sekund. V tem delu obravnavamo princip konvektivnega su{enja v toplem zraku. Predlagamo model, ki
obsega moment, transfer mase in toplote v coni atomizacije in vrtin~enja, neenakomerno porazdelitev delcev in efekte srkanja.
Za verifikacijo modela je bila izvr{ena obse`na eksperimentalna raziskava izparevanja vode z zrakom z razli~nimi
temperaturami zraka, hitrostjo napajanja, hitrostjo su{enja su{ilnega sredstva, parametri atomizacije in z za~etno porazdelitvijo
delcev. Spremembe temperature zraka v toku razpr{enega materiala, temperatura materiala, hitrost izparevanja, spremembe
premera Sauter in porazdelitev premera delcev v odvisnosti od razdalje od atomizatorja so bile dolo~ene eksperimentalno.
Dose`eno je bilo dobro ujemanje eksperimentalnih in teoreti~nih rezultatov. Obravnavan je bil tudi poskus uporabe modela za
izra~un razpr{ilnega su{enja v pilotnem su{ilniku. Za zrak, ki je bil tangencialno vpihovan v su{ilnik, je zna~ilno mo~no
vrtin~enje. Da bi se upo{teval vpliv tega vrtin~enja, sta bili spremenjeni bilanci toplote in mase sistema delec – su{ilno sredstvo.
Model je bil eksperimentalno preverjen na podlagi rezultatov su{enja 55-odstotne vodne raztopine {kroba. Ugotovljeno je, da
model dobro opisuje proces su{enja v analiziranem sistemu. V ~lanku je tudi pregled literature o razpr{ilnem su{enju in
obravnavane so najva`nej{e napake pri modeliranju. Kriti~no so ocenjeni dosedanji matemati~ni modeli procesa. Nekateri pa so
detaljno opisani.
Klju~ne besede: razpr{ilni su{ilnik, model, izparevanje, porazdelitev temperature
1 INTRODUCTION
Different suspensions of certain products react with
uneven drying, which depends on the properties of the
material that is dried. It is very difficult to recommend a
unique technique for spray driers. In comparison with
other systems, spray drying requires a relatively small
storage space. The service and maintenance are very
simple and can be performed with little labor. These
advantages and an economical heat transfer provide the
optimal cost of the final dried product. The drying time
is in most cases only a few seconds and sensitive
products, such as vitamins, chemical and food products
and others, who better submit to a high temperature in a
short period of time, but lower temperatures for a longer
period of time could be dried.
For substances and materials with over 30 % of
moisture, spray drying is a good solution to the problem
of drying.
Thanks to the principle of convective drying, the
dried material and contact drying agent, the system
receives intensive exchange of heat and mass and is very
difficult to give an adequate mathematical model of the
process, given the complex processes of heat-and-mass
transfer involved.
There have been many attempts to formulate a
mathematical description of spray drying since the
Materiali in tehnologije / Materials and technology 45 (2011) 2, 131–138 131
UDK 66.047:519.68 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 45(2)131(2011)
1950s. In spray drying there are many phenomena that
are difficult to represent in the form of a mathematical
model. For many years the polydispersity of spray,
entrainment effects, or problems of internal heat-and-
mass transfer in the disperse phase have been
inadequately considered in models. In practise, models
of the process that have been created since the beginning
of the 1970s were concerned only with the flow of the
drying agent.
A significant advancement was achieved in the
1970s, when the principles of momentum, heat-and-mass
balances between the continuous and disperse phases 1
(for monodisperse spray and plug flow of the air), and
especially when was presented a model in which the
entrainment effects and non-uniformity of atomization
were accountable were formulated 2. The solution of the
model of spray drying taking into account axial and
tangential velocity distributions of air was presented,
also. All the aforementioned relations (the entrainment
rate and the hydrodynamics of the drying agent) were
determined empirically. In further publications the
researchers presented many successful applications of
their model both in the modelling and the scaling-up of
the process. The simple model was recognized by many
researchers and studies based on this model have been
published 3.
However, the biggest change occurred when the
so-called particle source in the cell (PSI-cell) model for
gas-droplet flows was proposed 4. A method for solving
the Navier-Stokes equations and the continuity equation
was developed where the droplets were treated as
sources of mass, momentum and energy to the gaseous
phase. In the model, the gas flow field and droplet
trajectory are recalculated sequentially until the flow
field fails to change with repeated interactions. In the
model proposed originally 1 many simplified assump-
tions were made that referred mainly to the calculation
of trajectories and the temperature of the particles.
Further extensive studies resulted in developing more
sophisticated versions of the PSI-cell model, also
including commercial packages such as FLOEW3D and
PHOENIX.
The PSI-cell-type models were used successfully to
predict such subtle phenomena as the analysis of
low-frequency oscillations in the flow field inside the
spray drying chamber 5, a determination of the condi-
tions in which recirculation zones appear in the chamber
6 or the effects of inlet air geometry and cone angle on
the wall deposition rate 7.
It should be stressed, however, that a practical
application of the PSI-cell-type model requires profound
knowledge of spray drying and some experience in the
programming and calculation of systems with distributed
parameters. One problem, which is still unsolved, is the
choice of a turbulence model that should be applied in
the calculations 8 in such a way that the theoretical
results obtained are as near as possible to the experi-
mental values.
Therefore, there is still a considerable demand for
simple models that can be used to optimize the process,
to minimize its costs, and to evaluate product quality 9.
However, there are many publications of a classical
construction representing new concepts of mathematical
models of spray drying and offering an extensive
description of the whole process 10.
The quality of a spray-drying simulation depends not
only on the on the quality of the applied mathematical
model but also, and sometimes mainly, on the accuracy
of the determination of the initial parameters of the
process, and, in particular, the parameters of atomization
(mainly the initial particle size distribution and the spray
cone angles). An error made in this stage is carried
forward and may cause the results obtained to disqualify
the modelling process fully, irrespective of the type of
model applied.
In the literature, several interesting attempts have
been undertaken to overcome this problem. It was
proposed to determine experimentally the air
temperature distributions in the dryer chamber and then
to apply the results directly as the input in the integral
equations for the simulation of heat-and-mass transfer of
the process 11. Unfortunately, the experimental
determination of the temperature profiles in the dryer is
as difficult as the determination of atomization
parameters. It also proposed to perform calculations for a
randomly generated initial particle size distribution. The
model requires an assumption of the air temperature
distributions in the dryer 12,13. The authors state that the
model can be developed in such a way that the
temperature can be calculated iteratively. However, so far
neither has its developed a version of the model been
proposed nor has its experimental verification been
given.
In this paper, an attempt was made to formulate our
own mathematical model for spray drying, which, based
on the classical methods for balancing the transport
processes between the continuous and disperse phases,
would make it possible to analyze the phenomena on the
level accessible only to PSI-cell-type models. The
proposed model is a development of the concept
presenting a theoretical method for the calculation of the
entrainment rate 14. One of the aims of the study was to
apply the model and to verify it experimentally in two
different drying systems (flat air velocity profile,
pressure nozzle, evaporation of water and high swirl
air-flow pattern in a chamber, two-fluid nozzle and
drying of solution containing dissolved solids.).
2 MODEL OF THE PROCESS
The spray envelope was divided into two regions: the
spray boundary region and the spray core region. The
spray boundary region is determined by the location of
the particles of the atomized material. A change in the
location of the particles towards the radial direction also
means an increase in the spray diameter and the amount
D. TOLMAC et al.: A COMPARATIVE ANALYSIS OF THEORETICAL MODELS AND EXPERIMENTAL RESEARCH ...
132 Materiali in tehnologije / Materials and technology 45 (2011) 2, 131–138
of air inside the spray which comes into contact with
particles. The amount of air which controls the spray can
be determined by defining the position of the particles as
a function of the distance to the atomizer – swirler. The
main assumption made when developing the model is the
following: if spraying is monodispersive, or if only one
fraction of particles is considered, then in the cross-
section of the spray the temperature and air humidity are
distributed evenly. The temperature and humidity
gradients appear when the atomization is polydispersive.
This distribution can be determined by solving the
subsequent heat-and-mass transfer equations for parti-
cular fractions.
If the air velocity profile is flat, the total amount of
air which should be taken into account in the heat-
and-mass balance will be the sum of the air flowing
through the main core of the spray and the air dwg that
entered the spray, determined by the change in the
position of the particles (Figure 1).
The position of the particle as a function of the
distance to the atomizer can be calculated from the
generally known transfer equations that have the
following form for three dimensions:
d
d
px g
p
d
p p g
p p
U
h
g C
U U U g
d U
x x
= −
⎛
⎝
⎜
⎞
⎠
⎟ −
−⎡
⎣
⎢
⎤
⎦
⎥1
3
4
1



( )
px
(1)
d
d
pr pt
D
p pr gr
p p p
U
h
U
r
C
U U U g
d U x
= −
−⎡
⎣
⎢
⎤
⎦
⎥
2 3
4
1( )

(2)
d
d
pt pt
d
p pt gt
p p p
U
h
U
r
C
U U U g
d U x
= − −
−⎡
⎣⎢
⎤
⎦⎥
3
4
3
4
1( )

(3)
where Up is the relative particle velocity calculated from
Up = [(Upx – Ugx)
2 + (Upr –Ugr)
2 + (Upt – Ugt)
2]1/2 (4)
In the case of flat velocity profiles in the systems,
Eqs. (1)–(4) will be simplified, as Ugr = Ugt = 0. The
tangential and radial locations of participles will be
determined from
d
d
pr
p
r
h
U
U x
= (5)
d
d
t pr
p
x
h
U
U x
= (6)
If at first the mass balances for a solid-gas system,
assuming that atomization is monodispersive or that only
one fraction of the dispersed material is to consider and
assume that in the stream cross-section, there is neither a
temperature nor an air humidity gradient and that
continuous and dispersed phases are concurrent.
The mass balance has the form in Figure 1 and it is
given by equation 7.
wgY + w1X + dwgYo =
= (Wg + dwg) (Y + dY) + w1(X + dx) (7)
With the transformation we obtain where Y0 is the
bulk humidity (outside the stream).
d
d
d
d
d
d
g
g g
Y
h
w
h
Y Y
W
w
w
X
h
=
−
−0 2 (8)
Similarly, a heat balance for both phases can be
derived (Figure 1):
wgig + w1i1 + dwgig0 =
= (wg + dwg)(ig + dig) + w1(i1 + di) (9)
After the transformation we obtain:
dwgigo = wg dig + dwgig +w1di1 (10)
Humid air and material enthalpies are described by
the equations
ig = cgs Tg + (cvTg + H)Y (11)
and
i1 = (Xcw + cs)T1 (12)
through the calculation of the derivatives of air and
material enthalpies from
dig = cH dTg + (cvTg+H)dY (13)
and
di1 = (Xcw + cs) dT1 + cwT1 dX (14)
and subsequent substitution into Eq. (10), we obtain the
equation that can be used to determine the gas tempe-
rature in the stream cross-section taking into account
entrainment effects
d
d
d
d
d
d
g
g h
g g
g
g v g
T
h w C
i i
w
h
w c t H
Y
ho
= − − +
⎡
⎣⎢
−
1
( ) ( )Δ
– w Xc c
T
h
w c T
x
h1
1
1 1( )w s w
d
d
d
d
+ − ⎤⎦⎥ (15)
The entrainment rate dwg/dh can be determined on
the basis of the increment of the stream cross-section
(Figure 1):
Materiali in tehnologije / Materials and technology 45 (2011) 2, 131–138 133
D. TOLMAC et al.: A COMPARATIVE ANALYSIS OF THEORETICAL MODELS AND EXPERIMENTAL RESEARCH ...
Figure 1: Heat-and-mass transfer balances for spray-air system
Slika 1: Bilance prenosa toplote in mase za sistem razpr{evanje – zrak
[ ]d g
g g
w
U
Y
r dr r
x
=
+
+ −

1
2 2π π( ) (16)
therefore, the entrainment rate can be determined from
d
d
d
d
g pr gw
h
U
Y
r
g
=
+
2
1
π

(17)
Where the increment of the stream diameter dr/dh is
calculated from Eq. (5).
A complete description of the process requires a
determination of the humidity and material temperature
in the process.
To evaporate water from a solution containing solid
material of moisture X, the mass balance has the form
d
d
D
g p p
x
h
W x
V U x
=
+( )1 1

(18)
where wD is the evaporation rate given by
WD = ƒNkSPrg (Y* – Y) (19)
and ƒ is the relative drying rate defined by
x x
x X
−
−∫
0
cr e
(20)
where the subscript e refers to the equilibrium between
the gas and the solids and the subscript to the critical
point, above which drying is unhindered and the drying
rate depends solely on the rate of heat transfer to the
material.
In the case of drying of a solution containing solid
material, the drying is unhindered until a crust forms on
the droplets. The moisture content during crust formation
is the critical moisture content. For a moisture content
higher than the critical content or to evaporate the
solution that does not contain solid material, the value of
ƒ should be substituted as unity.
The heat balance for a particle material enables a
determination of the particle temperature as a function of
the distance to the atomizer. The heat delivered to the
particle will be used to evaporate the particle:
d(mpcpTp) = [khSp(Tg – Tp) – WD/N(H + cvTg)]dt (21)
Assuming that changes in the drying process concern
mainly the moisture content and temperature, the
left-hand side of Eq. (21) can be transformed to:
d(mpcpTp) = d[rpvp(cs + cwX) Tp] =
= rpVp[(Cs + CwX) dTp + cwTpdx] (22)
Substituting the right-hand side of Eq. (22) into (21),
we obtain after rearrangements
[ ]d
d
p g p
p p s w
T
h
a T T W N H C T s
d c c X
D v g p
=
− − +
−
+
6 ( ) ( )
( )
Δ

+
U c T
C c U
x
x
p w p
s wX p+
1
(23)
In this process, a continual change in the particle
diameters takes place, which has a significant effect on
heat-and-mass transfer coefficients and on the drag
coefficient. The actual particle diameter can be deter-
mined from 1:
d d
P P
P P
=
−
−
⎛
⎝
⎜
⎞
⎠
⎟p0
sw w
sw w
0
(24)
For a moisture content lower than the critical.
The particle diameter no longer decreases and thus
Eq. (24) ceases to be valid.
In all calculations, heat-transfer coefficients were
determined from the Ranz-Marshall equation and mass-
transfer coefficients from the Chilton-Colburn analogy,
where the psychometric coefficient and humidity
potential coefficient were taken as unity.
The model described above can be solved in the
following way:
1. For monodispersive atomization, the solution of Eqs.
(1)–(24) is used to calculate all the process
parameters taking into account the entrainment
effects.
2. For polydispersive atomization assuming constant air
temperatures and humidities in the cross-section of
the spray, the evaporation rate (Eq. (19) should be
calculated as a sum of the evaporations from
particular fractions i:
W N k S Y Y
i
n
d i i pi g= − −
−
∑ ∫1
1
 ( * ) (25)
For the entire stream of atomized material, the
entrainment rate (Eq. (17)) can be calculated on the basis
of the motion of the largest diameter particles.
Heat-and-mass balance equations for particles solved
separately for particular fractions are used to calculate
temperatures, moisture contents and the position of the
particular fraction in the dryer.
For polydispersive atomization when the air humidity
and temperature distributions in the spray cross-section,
the system of model equations should be solved
D. TOLMAC et al.: A COMPARATIVE ANALYSIS OF THEORETICAL MODELS AND EXPERIMENTAL RESEARCH ...
134 Materiali in tehnologije / Materials and technology 45 (2011) 2, 131–138
Figure 2: Calculating of humidity and temperature distribution in
spray
Slika 2: Izra~un porazdelitve vlage in temperature v curku
individually for each fraction starting with the particles
of the largest diameter dpn. A uniform temperature and
air humidity field Tgr and Yng will be obtained in the
stream cross-section. The files should be stored in the
computer’s memory because the motion of the next
fraction of particles of diameter dpn–1 will take place in
the field of gas humidity Yn and temperature Tgn
determined by the fraction of diameter dpn (Figure 2).
Proceeding in this way, we will obtain the tempera-
ture and air humidity distributions in the stream cross-
sections: the more fractions that are considered the more
accurate is the final calculated result.
3 EXPERIMENTAL DETAILS
To test the proposed mathematical model, experimen-
tal investigations were carried out in a dry 55 % solution
of starch and water in spray drying. Figure 3 shows a
schematic of the experimental equipment of the spray
dryer in which the experiments took place. In Table 1
are given the basic characteristics of the spray dryer.
Table 1: The basic characteristics of the spray-drying chambers
Name Units Dimensions
Diameter chamber spray dryer
Elevation spray drying
chamber
(mm)
(mm)
5000
5500
Cyclone diameter (mm) 1400
Power fan
Air flow
(kW)
(m3/h)
18.5
9750–14400
Atomizer (swirler)
Electric motor r/min
(kW)
(min–1)
0.75
3800
Inlet air drying is done by fans (3). Air over the
absorptive heater (4), with the help of fans (3). Thermal
power heater (4), can be regulated in the range 40–65
kW. The material is drying leads through the entrance
(6). Material for drying (55 % mass solution of starch
and water), empties into the atomizer (7) with the axle
(8) using electric propulsion. Spray nozzle rotates with
an rpm of n = 3800 min–1 and disperses the liquid
solution in small droplets. Aerosols are derived in the
form of a rotational disc with the diameter of 320 mm
and with a paddle. The extensive spray speed is 63.5
ms–1.
In contact with hot gas – air drying dispersed droplets
one obtains and the intensive exchange of heat and mass
is achieved. Air flow to the fan (3), can be regulated, and
the drying air velocity V = (0.17 to 0.24) ms–1 can be
obtained.
The dried material using the rotary feeder is
transported as a finished product (9). Air with dust
particles are drains with the piping (5) in the cyclone (2)
and the completion separation is achieved and dis-
charged in the atmosphere using fans (3).
4 MEASURING METHODS
The main part of the paper is concerned with the
measurement of the air and material temperature in the
spray, the evaporation intensity, and the particle size
distribution for various air-flow velocities, feed rates,
inlet air temperatures and spray cone angles. The tempe-
rature profiles of the air and atomized material as a
function of the distance to the nozzle were determined
by means of type T thermocouples and an EMT-02
temperature meter. To determine the temperature of the
gas flowing inside the spray, the thermocouples were
placed in special Teflon casings with an l8 mm external
diameter. The Teflon casing diameter was selected on the
basis of numerous tests as it achieved a stable tempera-
ture read-out. A similar measuring technique was used
by 17. To measure the evaporation intensity, a specially
designed measuring plate was used. This plate was
inserted into the funnel through the hatches in the
measuring section where atomized material was
collected on the whole cross-sectional area of the tunnel,
which made it possible to calculate the evaporation
intensity.
The change in the particle size distribution of the
atomized material as a function of the distance to the
spray nozzle was determined using the method of image
analysis. A sample of atomized material was placed in a
Petri dish where the surface was covered with a layer of
non-volatile silicone oil. The image obtained from a
microscope and camera was recorded on a video
cassette, and then processed by a microcomputer image
analyzer (IMAL 2561128). To minimize the statistical
error for every analyzed sample, 600 particle size
distributions of the atomized material were calculated
15,16. The ambient air temperature and humidity were also
measured with a THERM 2285-2 meter produced by
Alhborn with the probe FH 9626 -11.
D. TOLMAC et al.: A COMPARATIVE ANALYSIS OF THEORETICAL MODELS AND EXPERIMENTAL RESEARCH ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 131–138 135
Figure 3: Schematic of the experimental plant spray dryer: 1 –
chamber spray dryer, 2 – cyclone, 3 – fan, 4 – heater, 5 – air outlet, 6 –
supply of material, 7 – swirler (atomizer), 8 – drive axle, 9 – rotary
feeder, 10 – measuring points
Slika 3: Shema eksperimentalnega su{ilnika: 1 – komora su{ilnika, 2
– ciklon, 3 – vetrilo, 4 – grelnik, 5 – dovod zraka, 6 – dovod materiala,
7 – razpr{evalnik (atomizer), 8 – gonilna os, 9 – vrte~i se napajalnik,
10 – merilne to~ke
5 EXPERIMENTAL RESULTS AND DISCUSSION
One of the main sources of error in the spray drying
simulation is an imprecise specification of the initial
process parameters, for example, the spray cone angle
and particle size distribution. The parameters were
carefully determined for each swirl insert and for each
feed rate. Experimental investigations of the water
evaporation were carried out for three feed rates (98,
135, 157) kg h–1, and three values of air temperature (80,
100, 130) °C. An example of the initial particle size
distribution obtained at feed rate of 98 kg h–1 is shown in
Figure 4.
Both in this case and for other values of the initial
atomization parameters a characteristic log-normal
particle size distribution was obtained. The experimental
investigations show that the character of the distribution
does not change as a function of the distance to the
atomizer-swirler in all experimental runs.
All theoretical calculations were made with the
assumption that the sign of negligible changes in the air
and material temperature in the cross-section of ato-
mized material spray which was confirmed experimen-
tally.
The increase of the temperatures in the cross-section
spray dryer is only slight, Figure 5. The absence of an
air temperature gradient in the spray cross-section can be
explained with the design of the experimental set-up and
by process parameters. The air-velocity profile in the
tunnel of the dryer is flat and the turbulence is low. In the
study is up 18 a temperature distribution along the spray
dryer radius was observed and similar results were
obtained.
As simulation calculations show that in agreement
with experimental observations(in all experimental trials
no wetting of the tunnel wall was reported) the particle
of diameter corresponding to a 95 % cumulative percen-
tage undersize did not contact the walls of the tunnel
dryer, Figure 5 shows a comparison between the
theoretical and experimentally calculated changes in the
air temperature, material and evaporation rate as a
function of the distance from the atomizer for the same
feed rate and air-flow velocity. A theoretically calculated
spray radius is also plotted on the same graph. A rapid
drop in the air temperature and then its characteristic rise
are related to entrainment effects. An increase in the
amount of air in the spray due to its expansion causes the
D. TOLMAC et al.: A COMPARATIVE ANALYSIS OF THEORETICAL MODELS AND EXPERIMENTAL RESEARCH ...
136 Materiali in tehnologije / Materials and technology 45 (2011) 2, 131–138
Figure 6: Evaporation at different air temperatures
Slika 6: Izhlapevanje pri razli~nih temperaturah zraka
Figure 4: Initial particle size distribution
Slika 4: Za~etna porazdelitev velikosti delcev
Figure 5: Comparison of theoretical and experimental results
Slika 5: Primerjava teoreti~nih in eksperimentalnih rezultatov
temperature to rise despite intensive evaporation, until
the moment when the balance occurs between the heat
used for evaporation and the heat supplied into the spray
by the entraining air is achieved. After this time a drop in
the air temperature as a function of the distance to the
atomizer is observed. Similar profiles of particular
functions were obtained in many studies 17.
The experimental results shown in Figure 6 illustrate
the changes in the evaporation rate along the tunnel of
the dryer for different temperatures of the drying agent.
Air temperature has a decisive effect on the evaporation
rate. No significant influence of the drying agent velocity
on the shape of the curves discussed was observed.
Particles achieve a zero relative velocity after a short
period of time, which limits the convective heat transfer
between the continuous and disperse phases.
Lack of air circulation zones in the tunnel causes the
spray to compact and move parallel to the tunnel of the
dryer walls. The very good agreement between the
theoretical and experimental results is worthy of note.
The final figure discussed (Figure 7) shows the effect
of the initial particle size distribution on the evaporation
rate at the same air temperature and feed rates. The
analyses of the results show that the character of the
initial particle size distribution has a remarkable
influence on the whole evaporation process. Particles
with a smaller Sauter mean diameter are evaporated
much faster than particles having larger mean diameters.
In the case when the initial particle size distributions are
similar, the feed rate is the parameter that determines the
evaporation rate. Good agreement between theoretical
and experimental results was also obtained in this case.
6 CONCLUSIONS
A mathematical model of momentum, heat-and-mass
transfer during spray drying was proposed and verified.
The model of a classical construction reflects the
majority of phenomena occurring between the con-
clusions and dispersed phases during atomization. The
main difference between the proposed model and the
models described in the literature is the balancing of the
transport processes between the continuous and disperse
phases. It is possible to describe the trajectories of the
particle fractions and take into account the entrainment
effects. The model was verified both for the system of
simple as well as complex drying agents and for the
solution containing solid particles. The model presented
in this paper can be used to predict many subtle
phenomena that occur during spray drying, in particular
during complex air flow. In all theoretical simulations of
the drying process of evaporation, reasonable agreement
with the experimental data was achieved.
7 REFERENCES
1 Bellinghausen, R., Esper G. J., Gehrmann, D., Scale-up of spray
driers on the basis of laboratory scale experiments, Meet. of the
EFChE Drying Working Party, Utrecht, 1993
2 Crowe, C. T., Sharma, M. P, Stock, D. E., The particle-source- in-cell
(PSI-cell) model for gas-droplet flows, J. Fluid Eng., 99 (1977),
325–334
3 Gauvin, W. H., Katta, S., Basic concepts of spray drying design,
AIChE J., 22 (1976) 4, 713–724
4 Kuts, P. S., Samsonyuk, V. K., Akulich, P. V., Liquid phase
distribution in a spray formed by pneumatic nozzle in the spray dryer
chamber, in A. S. Mujumdar (ed.), Drying’92, Part A, Elsevier,
Amsterdam, (1992), 543–550
5 Langrish, T. A. G. and Zbiciñski, I., The effects of air inlet geometry
and spray cone angle on the wall deposition rate in spray dryers,
Trans. Inst. Chem. Eng. A, 72 (1994), 420–430
6 Larrgrish, T. A. G., Oakley, D. E., Keey, R. B., Bahu, R. E., Hutchin-
son, C, A., Time-dependent flow patterns in spray dryers, Trans. Inst.
Chem. Eng., 71 (1993), 355–360
7 Lefebvre, A. H, Atomization and sprays, Taylor & Francis, Bristol,
USA, 1989
8 Livesley, D. M., Oakley, D. E., Gillespie, R. F., Elhaus, B., Ranpuria,
C. K., Taylor, T., Wood, W., Yeoman, M. L., Developments and
validation of a computational model for spray-gas mixing in spray
dryers, in A. S. Mujumdar (ed.), Drying’92, Part A, Elsevier,
Amsterdam, (1992), 407–416
9 Miura, T., Othani, S., Maeda, S., Heat and mass transfer to and from
sprays, in A. S. Mujumdar (ed.), Drying’80, Vol. 1, Hemisphere,
McGraw Hill, New York, NY, 1 (1980), 351–356
10 Oakley, D. E., Bahu, R. E., Spray/gas mixing behaviour with spray
dryers, in A. S. Mujumdar and I. Filkova (eds.), Drying 9.1, Elsevier,
Amsterdam, (1990), 303–313
11 Papadakis, S. E., King, C. J., Air temperature and humidity profiles
in spray drying. 1. Features predicted by the particle source in cell
model, 2. Experimental measurements, Ind. Eng. Chem. Res.,
27(1988) 11, 2111–2123
12 Parti, M., Palancz, B., Mathematical model for spray drying, Chem.
Eng Sci., 29 (1974), 355–362
13 Prvulovic, S., Tolmac, D., Blagojevic, Z., Experimental research on
energetic charakteristics of starch dryer, FME Transactions, 37(2009)
1, 47–52
D. TOLMAC et al.: A COMPARATIVE ANALYSIS OF THEORETICAL MODELS AND EXPERIMENTAL RESEARCH ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 131–138 137
Figure 7: Effect of sauter mean on evaporation rate
Slika 7: Vpliv povpre~ja Sauter na hitrost izhlapevanja
14 Satija, S., A scale-up study of nozzle tower spray dryers, Drying
Technol., (1987) 1, 63–85
15 Som, S. K., Mitra, A. K., Sengupta, S. P., Evaporation, drop Som, S.
K., Mitra A. K., Sengupta, S. P., Evaporation, drop Iiquid spray
containing dissolved solids in a convective medium, Drying
Technol., 8 (1990) 3, 571–591
16 Tolmac, D., Introduction to the Theory of Drying with Examples
from Practice, Technical Faculty, Zrenjanin, Serbia, 2007
17 Zbiciriski, I., Grabowski, S., Strumillo, C., Kiraly, L., Krzanowski,
W., Mathematical modelling of spray drying, Comput. Chem Eng.,
12 (1988) 2/3, 209–214
18 Zhelev, J. B., Experimental investigation of flow pattern in a spray
dryer, Drying Technology 7 (1989) 3, 74–85
D. TOLMAC et al.: A COMPARATIVE ANALYSIS OF THEORETICAL MODELS AND EXPERIMENTAL RESEARCH ...
138 Materiali in tehnologije / Materials and technology 45 (2011) 2, 131–138
F. VODOPIVEC et al.: CREEP RESISTANCE OF MICROSTRUCTURE OF WELDS OF CREEP RESISTANT STEELS
CREEP RESISTANCE OF MICROSTRUCTURE OF
WELDS OF CREEP RESISTANT STEELS
ODPORNOST PROTI LEZENJU PRI MIKROSTRUKTURI ZVAROV
JEKEL, ODPORNIH PROTI LEZENJU
Franc Vodopivec, Monika Jenko, Roman Celin, Borut @u`ek, Danijela A. Skobir
Institute of Metals and Technology, Lepi pot 11, SI-1000 Ljubljana, Slovenia
franc.vodopivec@imt.si
Prejem rokopisa – received: 2011-03-02; sprejem za objavo – accepted for publication: 2011-03-28
Welds are essential parts of tube systems in high temperature power works producing electrical energy and have a heterogeneus
microstructure due to the welding gradient of temperature and cooling rate. A short summary is given of findings related to the
creep resistance of welds of creep resistant steels. In the majority of references it was found that creep resistance was lower for
the intercritical part of heat affected zone (HAZ). The effect of potential changes of intercritical microstructure on creep rate at
exploitation is discussed.
Key words: creep resistant steel,welds, creep rate, HAZ, intercritical zone
Zvari so bistvena komponenta cevnih sistemov v termoelektrarnah, njihova mikrostruktura pa je heterogena zaradi gradienta
temperature in ohlajanja pri varjenju. Pripravljen je bil kratek pregled ugotovitev v zvezi z odpornostjo razli~nih delov cone
toplotnega vpliva varjenja (TVC) proti lezenju v dosegljivih virih. Ve~ina virov navaja, da je odpornost najni`ja v podro~ju, kjer
je bilo med varjenjem jeklo segreto v interkriti~no podro~je CTV. Analizirane so mogo~e spremembe interkriti~ne
mikrostrukture in od njih odvisne spremembe hitrosti lezenja.
Klju~ne besede: jeklo odporno proti lezenju, zvari, hitrost lezenja, TVC, interkriti~na cona
1 INTRODUCTION
Welds are an essential parts of steels tube systems in
thermal power works. The heating cycle of welding, the
temperature gradient and cooling rate give to the weld
part of tubes, the heat affected zone (HAZ) a hetero-
geneous microstructure of products of transformation of
austenite at cooling. This microstructure depends on the
temperature at welding and the cooling rate. It consists
of a slightly affected microstructure of basic steel, the
products of transformation of cooling from intercritical
range, the range of temperature of stability ferrite +
austenite and the transformation at cooling of austenite
from a temperature up to melting point on the boundary
to the deposed metal. Also the HAZ microstructure may
be affected by the cooling rate and influenced
additionally by reheating at deposition of following
welding passes. The heating time is short, however, the
local temperature is sufficient for inducing changes of
grains size and may even change the quantity and
distribution of carbide particles hindering the movement
of dislocations and affecting the creep rate. For this
reason, the creep resistance of welds of creep resistant
steels has received considerable attention. In this work,
an abbreviated review of the most significant findings in
accessible references is given. Conclusions are proposed
with accent on more reliable findings and some so far
insufficiently clear facts are mentioned, also. The effect
of possible changes of microstructure in intercritical
zone on creep rate is examined, also.
2 REVIEW
For a 21Cr steel it was established that the weld
metal had a greater creep rate than the base metal due to
agglomerate of carbide particles facilitating the
cavitation.1 The creep activation energy of 337.5 kJ/mol
was deduced, a value in acceptable agreement with the
creep activation energy for -iron.2 For the steels
9Cr1Mo and 2.25Cr1Mo, the dissimilar weld joint
showed lower creep strength than base steels and the
creep deformation was for both steels concentrated in the
intercritical (ICZ, + range) zone of HAZ.3 Creep
cavities were generated mostly at grain boundaries in
ICZ, their number increased with creep deformation that
was self-accelerating.4 Also, it was established that at
creep tests at 650 °C and 675 °C, the change of hardness
was small for ICZ and great for weld metal, coarse
grained HAZ and base metal. Creep failure of welds
occurred with IV type cracking in ICZ with the highest
generation of voids. HAZ modelling showed for ICZ the
higher equivalent stress and strain.5
The higher ICZ susceptibility to cracking in welds of
a 1.25Cr0.5Mo steel in service may be due to the hetero-
geneous distribution of precipitates. Close to 90 % of
voids are found at grains of size lower than 10 μm.
However, 200 to 300 voids per mm2 did not shorten
significantly the steel rupture time.6 Calculation showed
that ICZ has a higher equivalent strain and high
hydrostatic pressure and than the base material submits
ICZ to a strong constraint.7 The distribution of voids and
Materiali in tehnologije / Materials and technology 45 (2011) 2, 139–143 139
UDK 669.14.018.8:621.791.05:620.18 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 45(2)139(2011)
equivalent stress were similar for all HAZ zones, thus
independent on microstructure. Also, it is suggested that
hydrostatic pressure may accelerate the coalescence of
voids to cracks.7 The greater crack susceptibility of ICZ
may be due to a heterogeneous distribution of precipi-
tates and the greater change of particles morphology
during creep.8
The grain boundaries sliding hinders the generation
of cavities at accelerated creep rate tests.9 For this
reason, the density of cavities could not be an absolute
parameter of creep damage effect in ICZ. From strain
rate measurements the rupture life can be predicted with
reasonable accuracy.9 The most effective factors reducing
the creep rupture strength of welds of steel P 91 are the
finer austenite grains that accelerate the growth rate of
subgrains from martensite laths and the softer martensite
matrix that both increase the rate of softening and creep
cavitation.10
The intrinsically higher ICZ creep rate is due to the
effect of stress triaxiality.11 Stress relief cracking could
occur in coarse grained HAZ because of sulphur and
phosphorus segregations.12 In the investigation of the
effect of triaxiality on creep with notched specimens, it
was found that the equivalent stress played a major role
in the multiaxial rupture life of a 1.25Cr0.5Mo steel.13
For a similar steel, it was deduced that multiaxial stress
could play a key role in the type IV failure of welds and
that this failure may be caused by grain boundaries
sliding of fine grains produced by the partial
transformation at welding.14 With small punch tests it
was found that the creep rupture time of HAZ was by
low stress levels shorter than for the base and deposited
metal15. Also, it was assumed that a film like carbide
phase at grain boundaries and coarsening of M23C6
particles may affect the creep rupture strength of HAZ.15
For the 9Cr1MoVNbN and 12Cr2MoWVTiB steels
the creep damage was greater for undermatched than for
equal and overmatched welds and for overmatched welds
the failure tendency was higher than for equal matched
welds.16 After normalising at 1050 °C and tempering at
780 °C, the creep rupture strength for HAZ was equal as
for the parent steel17. For proper understanding and
modelling of creep behaviour of welds, it is necessary to
know the individual creep strength of different weld
regions and the effect of stress multiaxiality that is
changed permanently because of creep and relaxation.18
The creep rupture strength of cross weld joints is
usually about 20 % to 30 % lower than that of the base
steel.19 Also, it was found that in steel P 91 the particles
spacing of secondary phases was higher in HAZ than in
base steel. Creep strength is for welds identical to that of
the base steels up to 575 °C and the creep resistance may
be influenced by post-weld heat treatment.20 Creep
fracture occurs in ICZ and the density of creep voids is
greater in the interior of the specimen section, where also
the multiaxial stress is larger.21
The creep strength of the 9 Ce to 12 Cr steels was
improved and the formation of small grains in ICZ
suppressed with addition of 0.1 % B because of the grain
strengthening effect of this element22. In23 the effect of
boron addition is confirmed and is established that no
coarsening of carbide particles occurred by 10 000 h
creep test at 650 °C. Also, the cross weld stress of boron
steel is up to 10 000 h of creep time higher for the weld
than for the base steel. The stress concentration in the
softened zone of ICZ explains it creep propensity only if
the smaller grain size is considered, also.24
Up to 575 °C the weld creep strength is in the scatter
band of ±20 % of the base steel, while, at higher
temperature the weld creep strength is below this range.25
In HAZ the recovery is faster because of the faster
deterioration of laths structure and the decrease of dislo-
cation density that facilitate a faster creep deformation.26
With small punch tests the localization of creep fracture
in ICZ was confirmed especially by small load and a
simple relation was established between the punch load
and the equivalent stress.27 Creep damage from prelimi-
nary uniaxial tests shortens the time to rupture by small
punch test.28
The following conclusions and remarks are derived
from this review:
– in welds of creep resistant steels the creep resistance
of the HAZ intercritical zone (ICZ) is the lowest and
the creep rate increased appropriately;
– creep failure occurs with coalescence of grain
boundary voids to type IV cracks;
– in operation, creep relaxation occurs that affects the
equivalent stress and the creep resistance;
– boron presence in steel increases the creep resistance
of the ICZ zone of HAZ;
– reliable weld creep modelling should be based on the
effect of operation time and temperature on creep
rate for individual constituents of HAZ and on stress
multiaxiality;
– data on the accurate creep rate of typical HAZ
microstructure constituents are insufficient;
– plastic deformation increases the number of defects
in metal lattice and accelerates diffusion and related
processes, also diffusion and creep. It is not clear
how these processes can be accelerated by stressing
without plastic deformation.
3 EFFECT OF CHANGES OF INTERCRITICAL
HAZ MICROSTRUCTURE ON CREEP
RESISTENCE
In creep resistant steels the creep rate depends on
spacing of particles, mostly of carbides of chromium,
vanadium, niobium and other metals added in the steel.
Two theoretical equations were developed for the depen-
dence of creep rate and particles spacing :29,30
 ( )
 	= b /k TG D2 2B (1)
and in the detachment concept of dislocations over-
coming of non coherent precipitates31,32
F. VODOPIVEC et al.: CREEP RESISTANCE OF MICROSTRUCTURE OF WELDS OF CREEP RESISTANT STEELS
140 Materiali in tehnologije / Materials and technology 45 (2011) 2, 139–143
 ( ) exp( )= ⋅ −6 /k TG E/k TB B (2)
With  – creep rate, b – Burgers vector, kB – Boltz-
mann constant, T – temperature in K, G – shear modulus,
 – acting stress, D – diffusion coefficient,  – density of
mobile dislocations and E – creep activation energy.
By constant volume share of carbides f, the particles
spacing and carbide particles size d are related with:30
 = 4 1 3d/ fπ / (3)
By constant volume share of carbide, the creep rate
increases with particles spacing which, by constant
volume share of carbide, depends of particles size, thus
of particles coarsening rate, also. This rate depends on
volume diffusion rate of the main metal bound in carbide
phase, fi.: chromium in M23C6 and vanadium and nio-
bium in MC carbides.
For the coarsening of M23C6 carbide particles the
experimental coarsening rate was determined33,34
d k t tce,1073K ce,1073K
3 26148 10= = ⋅ −. (4)
The Lifshitz-Slyouzov-Wagner35 (LSW) equation for
coarsening of particles is:
 d S Dt/ k TB
3 8 9= W (5)
With d3 – increase of particles size in time t, S –
equilibrium content of carbide constituting metal in
solution in the matrix,  – carbide particle-matrix inter-
face energy, W – volume of diffusing atoms, kB –
Boltzmann constant, D – diffusion coefficient and T –
temperature in K.
Introducing the parameters for the steel X20 the
coarsening kinetics for M23C6 particles in steel X20
tempered at 1073 K the relation was deduced:34
d tce,1073K
3 26131 10= ⋅ −. (6)
This relation is in acceptable agreement with the
experimental equation (4).
As shown in Figure 1, by equal particle size the
creep rate depends strongly on the share of particles in
stringers37.
The number of stringers per unity of surface
decreases with tempering time several times more
rapidly that particles coarsening. The explanation is the
faster growth of particles at grain and subgrain
boundaries38,39,40 very probably because the boundary
diffusion is faster then volume diffusion. The kinetics of
decrease of the number of stringers of particles the
relation was deduced:
n n k t tt t n
x= − = ⋅ − ⋅178 10 0 706 108 2. . (7)
with nt – density of stringers at the time t, ni – density of
stringers at t = 0, kn – rate of decrease of stringers
density (m–2 s–1) and t – tempering time (s).
After approximately 350 h of tempering at 800 °C the
stringers density was diminished below the critical value
of about 0.60 · 108 m–2, the creep rate increased by
approximately 8.5 times,37 while a much lower increase
would be attained from the increase of particles spacing
if the rate was calculated using equation (1) and the
particle spacing calculated using equation (3) based on
the experimentally assessed average particles size. These
findings suggest that the creep resistance of the inter-
critical part of HAZ depends to a significant extent of the
number if stringers of carbide particles and their
stability.
Investigations of the coarsening of carbide particles
in a 12Cr steel have shown41 that in specimens crept for
up to 16 000 h at 650 °C by the stress of 80 MPa the
coarsening rate of MX particles was 3.73 times greater
that in grip part of tested specimens.34
Parts of HAZ are heated to a different temperature
and a different microstructure is obtained. In some parts
of HAZ the temperature is theoretically sufficient for
significant change to steel creep resistance, however,
because of the short heating time only partial solution of
carbide particles is achieved with the greatest solution
for VC, lower for NbC and the lowest for M23C6.
Accordingly, it is logical to conclude that, especially the
particles distribution in parts of HAZ, heated to the
theoretical temperature of complete solution of VC and
NbC particles and the HAZ part heated at higher
temperature, but to low for a sufficient solution of M23C6
particles, the creep rate is higher mostly because of the
diminished effect of particles in stringers.
Equations (1) and (2) show that creep rate increases
with the density of mobile dislocations which is related
to the creep stress (), Burgers vector (b) and shear
modulus (G)42 and the constant  = 0.4.
   	= ( / MGb 2 (8)
For accelerated creep rate tests by  = 170 MPa,
G853K = 57.6 · 103 MPa 32, M = 3, b = 2.5 · 10–10 nm, the
dislocation density of  = 0.978 · 1014 m–2 is obtained.37
Let us examine how processes of transformation of
austenite to martensite and ferrite transformation in HAZ
F. VODOPIVEC et al.: CREEP RESISTANCE OF MICROSTRUCTURE OF WELDS OF CREEP RESISTANT STEELS
Materiali in tehnologije / Materials and technology 45 (2011) 2, 139–143 141
Figure 1: Dependence of accelerated creep rate on the number of
stringers of carbide particles in martensite in 0.18C11.5Cr1.08
Mo0.29V steel.36
Slika 1: Odvisnost med hitrostjo pospe{enega lezenja in {tevilom
nizov martenzitnih izlo~kov v jeklu 0.18C11.5Cr1.08 Mo0.29V36
may affect the mobile dislocation density. These dis-
location are generated by plastic deformation. Accord-
ingly, no significant effect could be expected from the
generation of elastic stresses by austenite to martensite
transformation and their relaxation. By transformation of
the mixture of austenite and ferrite, thus by cooling from
intercritical temperature, these elastic stresses could
relax with slight plastic deformation of ferrite and its
mobile dislocation density slightly increased. Thus, the
density of mobile dislocations could be slightly
increased in parts of HAZ with incomplete ferrite to
austenite transformation at heating. Carbide stringers
consist for a greater part of M23C6 particles. These
particles grow faster because of boundary diffusion39,40
and the rate of decrease of the density of stringers is
much greater than for the average coarsening rate in
stringers and the rate of stringers decmposition. For this
reason, it is expected that by heating at welding the creep
resistance is diminished much less by increase of
particles spacing than by the decrease of stringers
density and it is logical to conclude that the increased
creep rate for the intercritical part of HAZ is due mostly
to the decrease of the effect of stringers of carbide
particles on creep. After heating at higher temperature, a
great part of carbide particles is dissolved in austenite
and stringers form again at grain and subgrain
boundaries of martensite and their effect on creep rate is
restored. This interpretation of processes in HAZ, rsp. at
temperature of ferrite + austenite range, would be for the
creep resistant steel X20 the logical explanation for two
experimental findings: 1) compared to as delivered steel
X20, the accelerated creep rate at 580 °C and stress of
170 MPa the creep rate was increased for about one
order of magnitude after cooling from + range, while
it was slightly decreased after cooling from the
temperature of total carbide particles solution in
austenite and 2) the difference of creep rate was
decreased after longer tempering.43
4 CONCLUSION
On the base of the review of findings in references, of
the analysis of the effects of change of carbide particles
size and distribution, it is concluded that the decrease of
creep rate for the intercritical part of HAZ of creep
resistant steels is due to a wide extent to the decrease of
the effect of particles in stringers on the creep process.
This conclusion is supported by so far unpublished
experimental results on the effect of annealing
temperature on accelerated creep rate for the steel X20.
This work was supported by the company TE [o{tanj
and the Slovenian Research Agency (ARRS).
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F. VODOPIVEC et al.: CREEP RESISTANCE OF MICROSTRUCTURE OF WELDS OF CREEP RESISTANT STEELS
Materiali in tehnologije / Materials and technology 45 (2011) 2, 139–143 143

O. TOPÇU et al.: EFFECT OF THE MARTENSITE VOLUME FRACTION ...
EFFECT OF THE MARTENSITE VOLUME FRACTION
ON THE MACHINING OF A DUAL-PHASE STEEL
USING A MILLING OPERATION
VPLIV VOLUMENSKEGA DELE@A MARTENZITA NA OBDELAVO
DVOFAZNEGA DUALNEGA JEKLA Z REZKANJEM
Okan Topçu1, Mustafa Übeyli2, Adem Acir3
1TOBB University of Economics and Technology, Mechanical Engineering, 06560 Ankara - Turkey
2Osmaniye Korkut Ata University, Mechanical Engineering, 80000 Osmaniye – Turkey
3Gazi University, Mechanical Education, Teknikokullar, 06560 Ankara - Turkey
mubeyli@etu.edu.tr
Prejem rokopisa – received: 2010-05-17; sprejem za objavo – accepted for publication: 2010-12-27
Dual-phase steels are attractive because of their good combination of strength and ductility. However, an increase in the
martensite content will decrease the steel’s machinability significantly. Therefore, optimum cutting parameters in the machining
of such materials should be determined. In this paper, the machinability of a low-alloy steel having various martensite volume
fractions was investigated. The four groups of samples, three of which were intercritically annealed at three different
temperatures and one of which was normalized, were used. Then, the machining experiments were carried out using a
face-milling operation with two different tools. After that, the flank wear on the tools with respect to the chip-volume or
flank-wear limit was determined for each case. Next, the surface roughness of the samples was measured. Finally, the wear
behavior of the tools was examined with a scanning electron microscope (SEM). The experimental results indicated that the
flank wear of the tools increased dramatically with the increasing martensite content of the dual-phase steels.
Keywords: dual phase steel; machinability; tool wear; milling.
Dvofazno dualno jeklo se odlikuje po dobri kombinaciji trdnosti in duktilnosti. Pri pove~anju vsebnosti martenzita se
obdelovalnost pomembno zmanj{a, zato je treba za jeklo dolo~iti optimalne pogoje rezanja. V tem delu je bila raziskana
obdelovalnost malolegiranega jekla z razli~nim volumenskim dele`em martenzita. Uporabljene so bile {tiri skupine vzorcev, tri
izmed njih so bile interkriti~no `arjene pri treh razli~nih temperaturah, ena pa je bila samo normalizirana. Obdelovalni preizkusi
so bili izvr{eni s ~elnim rezkanjem z dvema razli~nima orodjema. Nato je bila za vsak primer dolo~ena ~elna obraba orodij
glede na prostornino ostru`kov ali mejna vrednost ~elne obrabe in tudi hrapavost povr{ine obdelancev. Obrabno vedenje orodij
je bilo raziskano z vrsti~nim elektronskim mikroskopom (SEM). Poskusi so pokazali, da ~elna obraba orodij dramati~no raste z
vsebnostjo martenzita v dualnem jeklu.
Klju~ne besede: dvofazno dualno jeklo, obdelovalnost, obraba orodij, rezkanje
1 INTRODUCTION
Making a dual-phase microstructure in a low-carbon
and/or alloyed steel provides not only a high strength but
also a good formability and yielding without serrations
1–7. In this type of microstructure, soft ferrite and hard
martensite phases are present and maintain the ductility
and the strength of the steel. The improvement in the
mechanical properties via forming a dual-phase micro-
structure in the steel supplies the benefit of reducing the
weight of systems. The studies on dual-phase steels have
been focused generally on the microstructural and
mechanical characterizations after intercritical heat-treat-
ment applications 1–17. On the other hand, the machina-
bility of these materials is also very important, particu-
larly with respect to the martensite fraction. Optimum
cutting parameters in machining should be determined
for dual-phase steels with respect to the martensite
fraction and the tool types. This knowledge would be
very helpful for accelerating the machining process and
extending the tool life. In the literature, the studies on the
machinability of dual-phase steels are very scarce 18–20.
Therefore, the machinability of these materials still
needs to be investigated in detail. El-Gizawy 18 carried
out a study on the machining characteristics of a
high-strength, low-alloy steel with 80 % martensite and
20 % ferrite using a face-milling operation in which
high-speed steel tools were used. It was concluded that
the steel with ferrite and martensite showed superior chip
disposability compared to that with ferrite and pearlite 18.
In addition, Sueyoshi and Tanaka 19 examined the machi-
nability of a tri-phase steel having the microstructure of
ferrite, martensite and graphite and compared the results
with a dual-phase steel containing ferrite and martensite
phases. They found that the drillability of the tri-phase
steel was better than the dual-phase one. In another study
20, tool wear and chip disposability in the machining of a
tri-phase steel were investigated and the results com-
pared with those of the dual-phase steel. It was stated
that the tri-phase steel provided better chip-disposal
behavior due to the fine graphite nodules. Furthermore,
the flank wear for the tri-phase steel was found to be
lower than that for the dual-phase steel at low cutting
speeds with high-speed steel tool.
Materiali in tehnologije / Materials and technology 45 (2011) 2, 145–150 145
UDK 669.14.018.2:620.1/.2:621 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 45(2)145(2011)
In this paper, the machinability of a low-carbon,
low-alloyed steel in a milling operation was presented.
The main aim was to determine the effect of the marten-
site volume fraction and the cutting speed on the flank
wear of coated and uncoated tools.
2 MATERIALS AND METHODS
The steel was supplied as a hot-rolled billet from a
private steel company in Turkey. Its composition was
0.28 % C, 1.45 % Mn, 0.21 % Cr, 0.20 % Si, 0.13 % V,
0.01 % Nb and (bal) Fe in terms of mass fractions (%).
The specimens were prepared as square shapes of size
(100 × 100 × 10) mm. Next, the intercritical heat treat-
ments at (737, 754 and 779) °C were applied to these
steel specimens to obtain three different martensite
volume fractions. These temperatures were utilized in
our previous study 16 and corresponded to low, medium
and high martensite volume fractions (approximately
30 %, 50 % and 80 %) in the steel. In addition, the
normalizing treatment at 900 °C was also performed for
one group of specimens to obtain a fine ferrite and
pearlite microstructure to compare with the intercritically
treated specimens. The details of the heat-treatment
procedures of this steel can be found in Ref. 16. In total,
four different specimen groups, according to their micro-
structures, were formed. After finishing the thermal
treatment procedures, the martensite volume fractions in
the intercritically heat-treated samples were checked
using an image-processing computer program. Further-
more, the macro-hardness of all the samples was
measured by applying the Rockwell C test 21. Finally, the
face-milling experiments were carried out on a CNC
vertical machining center (Johnford VMC-550 Fanuc
Series O-M) in dry conditions (Figure 1). Two different
types of tools, of which the properties are given in Table
1, were used in the face milling of the steel specimens.
The tool geometry and the cutting conditions are also
described in Table 2. The three different cutting speeds
(100, 160 and 220) m/min were used in the operations
while the feed was kept constant (0.2 mm per tooth). The
milling operations were performed until a chip volume
of 8000 mm3 or a flank wear limit (VB) of 0.3 mm was
reached. Furthermore, the surface roughness of the
machined samples was measured using a surface-
roughness measurement device, i.e., a Perthometer M1
(Mahr). The microscopic observations on the tool inserts
were made with the help of a scanning electron
microscope (SEM) and an optical microscope to see their
wear behavior clearly.
Table 1: Cutting-tool type and characteristics used in the experiment
Tabela 1: Tip rezalnega orodja, ki je bilo uporabljeno pri poskusih
Tool name Cutting tooldesignation
Substrate and
coating
Coating
technique
Depth of cut
(DOC) h/mm
Tool A TPMN160308 K20 (nocoating) Uncoated 1.0
Tool B TPMN160308 F620 + TiNcoating CVD 1.0
Table 2: Tool geometry and cutting conditions
Tabela 2: Geometrija orodja in pogoji rezanja
Cutting conditions Parameters
Clamping type of tool holder Collet
Tool holder diameter [mm] 32
Rake angle 0°
Relief angle 11°
Insert angle 60°
Cutting length, l/mm 100
Number of tooth Single
Approach angle 90°
Nose radius, r/mm 0.8
Cutting speed, v/(m/min) 100, 160, 220
Feed per tooth [mm per tooth] 0.2
Axial depth of cut, ha/mm 1
Radius depth of cut, hr/mm 10
3 RESULTS AND DISCUSSION
Figure 2 shows typical microstructures of the
samples with different thermal treatment histories. The
martensite volume fractions of samples, treated inter-
critically at (737, 754 and 779) °C, were approximately
(30, 50 and 80) %, respectively. The hardness was
recorded to be (30, 35 and 40) HRC for the same
samples, successively. The lowest hardness value
belonged to the normalized specimen, which was 20
HRC. Figure 3 depicts the change in the flank wear of
both types of tools at the cutting speed of 100/min in the
milling of four group specimens. It can be seen from this
figure that an increase in the hardness of the dual-phase
samples accelerated the wear of both types of tools.
Although the performances of both tools were similar for
the milling of the high martensitic samples (40 HRC),
the uncoated tool had the lower flank-wear values for the
milling of the low and medium martensitic samples. On
O. TOPÇU et al.: EFFECT OF THE MARTENSITE VOLUME FRACTION ...
146 Materiali in tehnologije / Materials and technology 45 (2011) 2, 145–150
Figure 1: Schematic view of the face-milling operation. In one pass
the tool moves from one end to the other end of sample at a certain
cutting speed.
Slika 1: Shemati~en pogled na proces ~elnega rezkanja. V enem
koraku se orodje pri dolo~eni hitrosti rezanja premakne iz enega na
drugi konec vzorca.
the other hand, the normalized and the low martensitic
dual-phase samples led to a similar flank wear on Tool A
(uncoated). In addition, the Tool B (coated) failed at the
beginning of the milling of the normalized specimen (20
HRC) due to the excessive formation of a built-up edge
(BUE).
The variation in the flank wear for the tools with
respect to the chip volume at the cutting speed of 160
m/min is illustrated in Figure 4. When the high marten-
sitic sample was machined, the flank wear increased very
rapidly on both tools. It exceeded the limit when the chip
volume of 3000 mm3 was removed from the samples. It
is interesting to note that in the machining of the low and
medium martensitic samples, the Tool A had lower
flank-wear values than the Tool B. Tool B (coated) was
fractured at the beginning of the milling process of the
normalized sample again so the data was missed for this
case. Furthermore, it also failed in the milling of the
dual-phase sample with 35 HRC after a chip volume of
7000 mm3 was removed.
O. TOPÇU et al.: EFFECT OF THE MARTENSITE VOLUME FRACTION ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 145–150 147
Figure 5: Flank wear on the tools at the cutting speed of 220 m/min in
the milling of the investigated samples.
Slika 5: ^elna obraba orodij pri hitrosti rezanja 220 m/min pri
rezkanju preiskanih vzorcev
Figure 3: Flank wear on the tools at the cutting speed of 100 m/min in
the milling of the investigated samples
Slika 3: ^elna obraba orodij pri hitrosti rezanja 100 m/min pri
rezkanju preiskanih vzorcev
Figure 2: Typical microstructures of the steel samples a) normalized,
b) intercritically annealed at 737 °C, c) intercritically annealed at 754
°C, d) intercritically annealed at 779 °C
Slika 2: Tipi~ne mikrostrukture vzorcev jekla: a) normalizirano, b)
interkriti~no `arjeno pri 737 °C, c) interkriti~no `arjeno pri 754 °C, d)
interkriti~no `arjeno pri 779 °C
Figure 4: Flank wear on the tools at the cutting speed of 160 m/min in
the milling of the investigated samples.
Slika 4: ^elna obraba orodij pri hitrosti rezanja 160 m/min pri
rezkanju preiskanih vzorcev
Figure 5 presents the tool-wear behavior versus chip
volume during the milling of the investigated samples at
a cutting speed of 220 m/min. The effect of the sample
hardness on the flank wear of the tools was clearly seen.
Increasing the cutting speed of the tool caused faster tool
wear in all cases. Tool B exceeded the flank-wear limit
for the dual-phase samples with 35 HRC and 40 HRC
before removing the chip volume of 2000 mm3. Further-
more, Tool B fractured after having seven passes over the
dual-phase sample having 30 HRC. At the cutting speed
of 220 m/min, the normalized specimen was machined
by both tools without causing any fracture. Figures 6
and 7 show a typical SEM view of the flank wear
observed on the tools after machining. An increase in the
cutting speed of the process increases the temperature of
the cutting zone significantly. This leads to a softening of
the chip and the formation of a BUE. For this reason, the
removal of the BUE from the cutting tool becomes much
easier. On the other hand, in the intercritically annealed
samples, the hardness values are much higher than the
normalized one. So that BUE formation is not so severe
in these samples.
Even though the performances of both tools are close
to each other at all cutting speeds in the machining of
high martensitic dual-phase steel, Tool A exhibited a
better performance than Tool B in the machining of other
types of samples. The adhesion of the chips to the coated
tool was found to a much greater extent, especially for
the cutting speeds of 100 m/min and 160 m/min. Further-
more, the coating material of TiN was easily removed
from the tool by an abrasive wear mechanism for the
harder samples. Therefore, its protection ability was
O. TOPÇU et al.: EFFECT OF THE MARTENSITE VOLUME FRACTION ...
148 Materiali in tehnologije / Materials and technology 45 (2011) 2, 145–150
Figure 7: View of Tool B after machining of the low martensitic
dual-phase steel a) X100, b) X200
Slika 7: Videz orodja B po obdelavi jekla z malo martenzita; a)
pove~ava 100-kratna, b) 200-kratna
Figure 8: A typical failure of the coated tool after milling: a) the
normalized specimen, b) the dual-phase specimen with 30 HRC, c) the
dual-phase specimen with 40 HRC (100-times)
Slika 8: Zna~ilna oblika po{kodb prekritega orodja po rezkanju: a)
normalizirani vzorec, b) dualno jeklo s trdoto 30 HRC, c) dualno jeklo
s trdoto 40 HRC (pov. 100-kratna)
Figure 6: Typical SEM view of the flank wear on Tool A after
machining of the normalized specimen at the cutting speed of 160
m/min.
Slika 6: Tipi~na SEM-slika ~elne obrabe orodja A po obdelavi nor-
maliziranega vzorca pri hitrosti rezanja 160 m/min
largely destroyed by the hard martensite phase in the
dual-phase steels. It is generally expected that the
coating material reduces the friction between the tool
and the workpiece and causes lower temperatures and
longer tool lifetimes. However, lower temperatures
prevent the softening and effective removal of the BUE
from the cutting zone. As mentioned before, higher the
BUE accumulation takes place in the cutting zone, the
higher the mechanical stresses form on the tool. Figure 8
illustrates the obtained wear of Tool B after the milling
of various samples. It is clear that the mechanical cracks
dominate the failure mechanism for this tool. The
abrasive wear caused by the chip on the rake face can be
seen. On the other hand, thermal cracks were widely
observed in Tool A, especially at the cutting speeds of
160 m/min and 220 m/min (Figure 9). The crater wear
accelerated the fracture of Tool A significantly.
The surface roughness for the machined samples with
Tools A and B is presented in Figure 10. It changed
between 0.5 μm and 1.3 μm, depending on the cutting
speed and the type of workpiece material. When the
coated tool was used, the highest surface-roughness
values were recorded for the normalized samples. In
addition to that, there seemed to be a low tendency to
increase the surface roughness of the dual-phase samples
with increasing the cutting speed for the milling with
either Tool A or B. The surface roughness was mainly
influenced by the BUE formation.
4 CONCLUSIONS
The results of the machinability experiments for
dual-phase steel using a milling operation led to the
following conclusions:
• The machinability of the dual-phase steel decreased
substantially with an increase in the martensite
content.
• An increase in the cutting speed led to faster
tool-flank wear on both the uncoated and coated
tools.
• The highest BUE formation was observed on the
coated tool (Tool B) in the milling of the normalized
samples.
• Thermal cracks dominated the failure mechanism of
the uncoated tool, whereas the mechanical cracks
governed that of the coated tool.
• The surface roughness of the dual-phase samples had
a tendency to slowly increase with an increasing
cutting speed.
For future studies, the effect of the feed and the depth
of the cut on the tool wear can be investigated to
O. TOPÇU et al.: EFFECT OF THE MARTENSITE VOLUME FRACTION ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 145–150 149
Figure 10: Surface roughness of the machined samples
Slika 10: Hrapavost povr{ine obdelanih vzorcev
Figure 9: A typical view of the cracks on the uncoated tool after
milling the dual-phase sample with: a) 30 HRC, b) 35 HRC, c) 40
HRC (100-times)
Slika 9: Zna~ilen videz razpok na nepokritem orodju po rezkanju
dualnih vzorcev s: a) 30 HRC, b) 35 HRC, c) 40 HRC (pov.
100-kratna)
understand the machinability of the dual-phase micro-
structure in details.
5 REFERENCES
1 J. Y. Koo, G. Thomas, Thermal cycling treatments and microstruc-
tures for improved properties of Fe-0.12 % C-0.5 % Mn steels, Mater
Sci. Eng., 24 (1976), 187
2 Y. Tomota, K. Kuroki, T. Mori, I. Tamura, Tensile deformation of
Two-Ductile-Phase Alloys: Flow curves of α-γ Fe-Cr-Ni alloys, Mat.
Sci. Eng., 24 (1976), 85
3 R. G. Davies, The mechanical prooperties of zero-carbon ferrite-
plus-martensite structures, Met. Trans., 9A (1978), 451
4 H. K. D. H. Bhadeshia, D. V. Edmonds, Analysis of the Mechanical
Properties and Microstructure of a High-Silicon Dual -Phase Steel,
Metal Sci., 14 (1980), 41
5 N. K. Balliger, T. Gladman, Work hardening of dual phase steels,
Metal Sci. 15 (1981), 95
6 M. Sudo, M. Higashi, H. Hori, T. Iwai, S. Kambe, Z. Shibata, Effects
of microstructures on the mechanical properties of multi-phase sheet
steels. Trans. ISIJ, 21 (1981), 820
7 N. Matsumura, M. Tokizane, Microstructure and mechanical pro-
perties of dual-phase steel produced by intercritical annealing of lath
martensite, Trans. ISIJ, 24 (1984), 648
8 T. Hüper, S. Endo, N. Ishikawa, K. Osawa, Effect of volume fraction
of constituent phases on the stress-strain relationship of dual phase
steels, ISIJ Int., 39 (1999), 288
9 S. S. M. Tavares, P. D. Pedroza, J. R. Teodosio, T. Gurova, Mecha-
nical properties of a quenched and tempered dual phase steel, Scr.
Mater., 40 (1999), 887
10 S. Sun, M. Pugh, Properties of thermomechanically processed dual-
phase steels containing fibrous martensite, Mater. Sci. & Eng., A 335
(2002), 298
11 J. Lis, A. K. Lis, C. Kolan, Processing and properties of C-Mn steel
with dual-phase microstructure, J. Mater. Process Technol., 162–163
(2005), 350
12 M. Erdogan, S. Tekeli, The effect of martensite volume fraction and
particle size on tensile properties of a surface-carburized AISI 8620
steel with a dual-phase core microstructure, Mater Charact, 49
(2003), 445
13 X. Liang, J. Li, Y-H. Peng, Effect of water quench process on
mechanical properties of cold rolled dual phase steel microalloyed
with niobium, Mater Lett, 62 (2008), 327
14 Y. J. Chao, Jr. JD. Ward, R. G. Sands, Charpy impact energy, fracture
toughness and ductile-brittle transition temperature of dual-phase
590 steel, Mater Des., 28 (2007), 551
15 M. Sarwar, T. Manzoor, E. Ahmad, N. Hussain, The role of connec-
tivity of martensite on the tensile properties of a low alloy steel,
Mater. Des., 28 (2007), 1928
16 O. Topçu, M. Übeyli, On the microstructure and mechanical
characterisations of a low carbon and micro-alloyed steel, Mater.
Des., 30 (2009) 8, 3274
17 O. Topçu, M. Übeyli, T. Demir, On the hardenability of an intercri-
tically heat treated microalloyed steel, Instr. Sci. and Tech., 38
(2010), 178
18 A. S. El-Gizawy, Machining characteristics of dual phase steel,
Materials and Manufacturing Processes, 5 (1990) 3, 377
19 H. Sueyoshi, R. Tanaka, Heat-treatment and machinability of the
tri-phase steel composed of ferrite, martensite, and graphite, Nippon
Kinzoku Gakkai-si, 54 (1990) 2, 231
20 H. Sueyoshi, H. Tanaka, K. Suenaga, R. Tanaka R., Tool wear and
chip-disposability in the cutting of the tri-phase steel composed of
ferrite, martensite, and graphite, Nippon Kinzoku Gakkai-si, 54
(1990) 5, 589
21 ASTM Standards, Designation E 18-93, Standard test methods for
Rockwell hardness and Rockwell superficial hardness of metallic
materials; 1993
O. TOPÇU et al.: EFFECT OF THE MARTENSITE VOLUME FRACTION ...
150 Materiali in tehnologije / Materials and technology 45 (2011) 2, 145–150
R. CELIN, F. TEHOVNIK: DEGRADATION OF A Ni-Cr-Fe ALLOY IN A PRESSURISED-WATER ...
DEGRADATION OF A Ni-Cr-Fe ALLOY IN A
PRESSURISED-WATER NUCLEAR POWER PLANT
DEGRADACIJA ZLITIN Ni-Cr-Fe V TLA^NOVODNIH JEDRSKIH
ELEKTRARNAH
Roman Celin, Franc Tehovnik
Institute of metals and technology, Lepi pot 11, 1000 Ljubljana, Slovenia
roman.celin@imt.si
Prejem rokopisa – received: 2011-01-03; sprejem za objavo – accepted for publication: 2011-01-31
In the early days of pressurized-water nuclear-power-plant design Ni-based alloys were selected because of their good
mechanical properties and corrosion resistance. Alloy 600 was used for some reactor-coolant pressure-boundary components
and Alloy 82/182 was used for welds. Industrial experience in the past three decades has shown that Alloy 600 components and
Alloy 82/182 welds are susceptible to primary-water stress-corrosion cracking (PWSCC). PWSCC is the intergranular or
transgranular cracking due to the combined action of stresses, temperatures and components in contact with the primary water
(reactor coolant). PWSCC leaks and cracks were detected on the reactor-coolant pressure-boundary components. In this work
some characteristics of the Alloy 600 and Alloy 82/182 welds and their PWSCC degradation are presented.
Key words: Ni-Cr-Fe alloy, stress corrosion cracking, nuclear power plant, repair
V za~etni fazi konstruiranja tla~novodnih jedrskih elektrarn so bile zaradi dobrih mehanskih lastnosti in korozijske odpornosti
izbrane zlitine na osnovi Ni. Za nekatere komponente na tla~ni meji reaktorskega hladila je bila izbrana zlitina z oznako Alloy
600, za zavarjene spoje pa zlitini z oznako Alloy 82/182. Izku{nje industrije z uporabo zlitine Alloy 600 in Alloy 82/182 v
zadnjih treh desetletjih pa so pokazale, da so le-te ob~utljive za napetostno korozijsko pokanje. Napetostno korozijsko pokanje
je pojav interkristalnih ali transkristalnih razpok, ki nastanejo zaradi skupnega vpliva napetosti, temperature in stika
komponente z reaktorskim hladilom. Zaradi napetostnega korozijskega pokanja so bile na komponentah tla~ne meje
reaktorskega hladila odkrite razpoke in netesna mesta. Namen prispevka je predstavitev lastnosti zlitin Alloy 600 in Alloy
82/182, njihova uporaba v jedrskih elektrarnah ter napetostno korozijsko pokanje.
Klju~ne besede: zlitina Ni-Cr-Fe, napetostno korozijsko pokanje, jedrska elektrarna, popravilo
1 INTRODUCTION
Alloy 600 was originally selected for use in smaller-
diameter piping penetrations in pressurized-water-reactor
(PWR) nuclear plants because of its good corrosion
resistance and a coefficient of thermal expansion similar
to that of low-alloy steel vessels and piping material.
Most of these penetrations are attached to the vessel or
piping with Alloy 82 or 182 (nickel-chromium-iron)
J-groove welds. Alloy 82/182 weld materials have also
been used for field-butt welds.
Primary-water stress corrosion cracking (PWSCC) of
Alloy 600 nozzles and Alloy 82/182 weld metal became
a source of concern in non-steam generator tubing in the
mid-1980s. Some significant PWSSC events in the past
twenty years are listed below1:
• 1991 – A through wall crack leak to the top of the
reactor vessel was detected during a pressure test at
the nuclear power plant Bugey 3 in France.
• 2000 – Cracks were discovered in Alloy 182 welds
joining the low-alloy-steel reactor vessel hot-leg
nozzles to stainless-steel pipes at Ringhals 4
(Sweden) and VC Summer (United States).
• 2002 – The most severe event was the NPP Davis-
Besse case. PWSSC and the significant boric acid
corrosion of carbon steel material caused a serious
degradation of the reactor vessel closure head.
• 2003 – A small leak was discovered on a pressurizer
relief nozzle at Tsuruga 2 (Japan) due to an axial
crack in the Alloy182/82 butt weld between the
low-alloy-steel nozzle and the stainless-steel relief-
valve pipeline.
• 2005 – Calvert Cliffs NPP (United States) identified
crack indications in an Alloy 182/82 dissimilar metal
weld on a 2-in. (51-mm) diameter hot-leg drain
nozzle. Two axial crack indications were contained
entirely within the weld and butter area.
The list of components that are known to contain
Alloy 600/82/182 in at least some nuclear-power-plant
designs includes:
• reactor vessel heads,
• reactor vessel hot-leg and cold-leg nozzles,
• reactor vessel bottom-mounted instrument pene-
trations,
• steam generator primary nozzles and tubes,
• pressurizer and heat exchangers,
• reactor coolant loop-pipe branch connections.
Some of the generic locations of Alloy 600/82/182
are shown in Figure 1.
The main types of affected welds are:
• J-groove welds of the control rod drive mechanism
(CRDM) penetrations (Figure 2), reactor vessel
Materiali in tehnologije / Materials and technology 45 (2011) 2, 151–157 151
UDK 621.311.25:620.193:669.14 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 45(2)151(2011)
bottom-mounted instrumentation penetrations (BMI),
steam generator (SG) drain lines, pressurizer
instrument nozzles and hot-leg instrument nozzles;
• butt welds or full-penetration dissimilar-metal welds
of reactor pressure vessels and pressurizer nozzles
(Figure 3).
Alloy 600
Inconel Alloy 600 with UNS N06600 or W. No.
2.4816 is a standard nickel-chromium-iron engineering
material for heavy-duty applications. The limiting
chemical composition2 of the Alloy 600 is shown in
Table 1.
Table 1: Alloy 600 chemical composition in mass fractions (w/%)
Tabela 1: Kemijska sestava Alloy 600 v masnih dele`ih (w/%)
Ni Cr Fe C Mn S Si Cu
>72 14–17 6–10 
0.15 
1.0 
0.015 
0.5 
0.5
Alloy 600 is a stable, austenitic solid-solution mate-
rial. The high nickel content gives the alloy a good
corrosion resistance in many organic and inorganic
compounds. Chromium provides the resistance to
sulphur compounds and the resistance in oxidizing con-
ditions at high temperatures. The alloy can be hardened
and strengthened only by cold work.
The phases precipitating in the microstructure are
titanium nitrides, titanium carbides and chromium
carbides. The nitride particles are stable at all tempera-
tures below the melting point and are unaffected by any
heat treatment. At temperatures between 540 °C and 980
°C the chromium carbide precipitates from the solid
solution at the grain boundaries and in the matrix.
Figure 4 shows the austenitic microstructure of Alloy
600 with carbide particles visible in the polished and
etched metallographic specimen.
Alloy 600 components can be fabricated by press
forging, hammer forging, hot rolling, forming and
machining from a bar product (ASME II SB-166) or cold
drawing and hot finishing from a pipe and tube product
(ASME II SB-167).2 The typical mechanical properties3
for various forms and conditions are listed, for infor-
mation only, in Table 2.
The alloy’s strength and oxidation resistance at high
temperatures make it useful for many applications for
R. CELIN, F. TEHOVNIK: DEGRADATION OF A Ni-Cr-Fe ALLOY IN A PRESSURISED-WATER ...
152 Materiali in tehnologije / Materials and technology 45 (2011) 2, 151–157
Figure 4: Alloy 600 austenitic microstructure (a) with chromium
carbide particles at the grain boundaries (b)
Slika 4: Avstenitna mikrostruktura zlitine 600 (ao) z zrni kromovega
karbida na mejah kristalnih zrn (b)
Figure 2: J-groove weld on a reactor-vessel head
Slika 2: J-zvar na glavi reaktorske posode
Figure 3: Nozzle-to-pipe dissimilar-metal weld
Slika 3: Zvar med cevnim nastavkom in cevovodom iz razli~nih
materialovFigure 1: Generic locations of Alloy 600 components and Alloy
182/82 welds
Slika 1: Splo{ne lokacije komponent iz zlitine 600 in zvarov zlitin
82/182
a)
b)
furnace components in the heat-treatment industry. In the
aeronautical industry Alloy 600 is used for a variety of
engine and airframe components that must withstand
high temperatures, for example, exhaust liners and
turbine seals.
In pressurized-water reactor (PWR) nuclear power
plants, Alloy 600 has been used for steam-generator
tubes, control rod drive mechanism (CRDM) nozzles,
reactor vessel bottom mounted instrument (BMI) pene-
trations, pressurizer heater sleeves and other pressure-
retaining components.
The Alloy 600 was selected for use in nuclear power
plants because of:
• Its good mechanical properties, similar to those of
austenitic stainless steels.
• Its good general corrosion resistance in high-tempe-
rature water environments and resistance to caustic
stress-corrosion cracking better than austenitic stain-
less steels.
• It can be welded to carbon, low-alloy and austenitic
stainless steels.
• It is a single-phase alloy requiring no post-weld heat
treatment, also when submitted to post-weld heat
treatments required for low-alloy steel parts to which
it is welded. The resulting sensitization (decreased
chromium levels at grain boundaries associated with
the precipitation of chromium carbides at the
boundaries) does not result in a high susceptibility to
chloride attack exhibited by austenitic stainless steels
exposed to such heat treatments.
• Its thermal expansion properties, between those of
carbon/low-alloy steels and austenitic stainless steels,
make Alloy 600 a good transition metal between
these steels.
2 PRIMARY-WATER STRESS-CORROSION
CRACKING (PWSCC)
PWSCC is the initiation and propagation of inter-
granular cracks through the material in a seemingly
brittle manner, with little or no plastic deformation of the
bulk material and without the need for cyclic loading.
Generally, it occurs at stress levels close to the yield
strength of the bulk material, but does not involve
significant material yielding.
An analysis of damaged components showed that
PWSCC has a tendency to occur most quickly in parts
which were1:
• fabricated from more susceptible or higher-strength
materials,
• machined or cold worked prior to welding,
• installed using methods that can produce high
residual stresses, such as welding or roll expansion,
• operating at high temperatures,
• not stress relieved after installation.
Over the past three decades a significant number of
laboratory studies and industry events reported on the
PWSCC failure mechanism for components from Alloy
600 with Alloy 182/82 welds. The results of the tests and
data from industry have shown that the occurrence of the
PWSCC depends on the simultaneous contribution of:
• susceptible material,
• corrosive environment (primary water (reactor
coolant)),
• high stresses including residual stress and operating
stress.
2.1 Material
The fabrication process, heat treatments, and che-
mical compositions affect the formation of the material’s
microstructure and are the main contributors to the
PWSCC susceptibility. Alloy 600 bars, rods, plates,
pipes and stripes are usually heat treated to reduce the
yield strength and increase the material toughness to an
acceptable level. A higher heat-treatment temperature
and a longer duration result in a lower yield strength.
The fabrication process, heat treatment, and chemical
compositions affect the formation of the material’s
R. CELIN, F. TEHOVNIK: DEGRADATION OF A Ni-Cr-Fe ALLOY IN A PRESSURISED-WATER ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 151–157 153
Table 2: Mechanical properties of Alloy 600 (for information only)
Tabela 2: Mehanske lastnosti zlitine 600 (samo za informacijo)
Product Form Condition Tensile strength (MPa) Yield strength (MPa) Elongation (%)
Rod and Bar
Cold drawn Annealed 550–690 170–345 55–35
As drawn 724–1034 550–860 30–10
Hot finished Annealed 552–689 205–345 55–35
Hot finished 586–827 240–620 50–30
Plate
Hot rolled Annealed 55–725 205–345 55–35
As rolled 580–760 240–450 50–30
Tube and Pipe
Hot finished Hot finished 520–690 170–345 55–35
Annealed 520–690 170–345 55–35
Cold drawn Annealed 550–690 170–345 55–35
microstructure and are the main contributors to the
PWSCC susceptibility.
The heat-treatment temperature, annealing time and
the carbon content are interrelated and affect the alloy’s
microstructure. A carbide-precipitation diagram for the
Alloy 600 material shown in Figure 5 could be used to
assess the effect of the heat-treatment temperature and
time on the microstructure.4,5 The kinetics of precipi-
tation depends on the velocity of the diffusion processes,
which is greater at high temperature. The amount of
carbide precipitates depends on the annealing time.
A high-temperature (1066 °C) heat treatment for a
sufficient time (zone D in Figure 5) could form an inter-
granular carbide-particles network without the de-chro-
mization of the area of the adjacent grain boundary.
Stress-relief annealing in the range 700–800 °C in zone
C in Figure 5 also limits the intragranular carbide
formation and improves the material’s resistance to
PWSCC. A post-fabrication heat treatment in zones B
and A will result in PWSCC-susceptible material.
The mass fraction of carbon-content (w(C)/%) is an
important factor in the microstructure formation of Alloy
600 due to heat treatment. The lower w(C) moves the
x-curve of the carbide-precipitation diagram in Figure 5
to the right and improves the conditions for intergranular
carbide formation.
Generally, in Alloy 600 a good grain-boundary
carbide network increases the PWSCC resistance.
However, the susceptibility of Alloy 600 also depends on
the surface cold work due to machining, grinding and
reaming. A surface layer with a high cold work of the
highly susceptible material, for example, components
that were machined from bar stock and with weld root
grinding, is considered to be highly PWSCC susceptible.
2.2 Environment
There are several environmental parameters that are
influential on the PWSCC initiation and growth. The
most significant is the primary-water temperature. The
results of laboratory tests indicated that PWSCC is a
thermally activated process and that the crack initiation
and growth rate are strongly temperature dependent.6
Like many temperature-dependent processes, the
correlation between the PWSCC growth rate and the
temperature can be expressed with the Arrhenius
equation7:
 expa C
Q
RT
= −⎛⎝
⎜ ⎞
⎠
⎟ (1)
where:
a = growth rate
Q = activation energy for the crack growth phase
R = ideal gas constant
T = temperature (K)
C = constant
The other influential environmental parameters are
chemical additives to the primary water. The chemistry
of the primary water (reactor coolant) is maintained by
the chemical and volume control system, which is
designed to allow the operators to regulate the water’s
chemical composition. The major use of this system is to
control the primary-water boron content as a function of
the nuclear reactor’s power level. With the addition or
removal of lithium hydroxide the reactor coolant’s pH
value is controlled. The system is also designed to allow
the addition of hydrogen during normal operation.
Hydrogen gas is dissolved in the reactor coolant to
scavenge all the dissolved oxygen, which may be present
in the primary water. The effect of the lithium
concentration and pH value on the PWSSC is minimal.8
Tests using crack growth rate (CGR) specimens have
shown that crack growth tends to be faster when the
water’s electrochemical potential, depending on
hydrogen concentration, is close to the potential where
the Ni/NiO phase reaction occurs.8,9 Higher or lower
values of the hydrogen concentration decrease the crack
growth rates (Figure 6).
R. CELIN, F. TEHOVNIK: DEGRADATION OF A Ni-Cr-Fe ALLOY IN A PRESSURISED-WATER ...
154 Materiali in tehnologije / Materials and technology 45 (2011) 2, 151–157
Note: 1 mil/d is 1.0 · 10–6 m/h or 29 · 10–10 m/s
Figure 6: Influence of hydrogen concentration on the crack growth
rate of PWSCC in Alloy 600 at 338 °C
Slika 6: Vpliv koncentracije vodika na hitrost napetostnega korozij-
skega pokanja za zlitino 600 pri temperaturi 338 °C
Figure 5: Carbide-precipitation diagram for Alloy 600
Slika 5: Diagram izlo~anja karbidov za Alloy 600
2.3 Stress
Allowable stresses for a nuclear power plant
component are specified in ASME Boiler and Pressure
Vessel Code Section III. These requirements apply to
operating loadings, such as internal pressure, differential
thermal expansion, dead weight, and seismic loading. On
the other hand, the industry design standards do not
typically address residual stresses that can be induced in
the parts during fabrication. These residual stresses are
often much higher than the stresses in operation. In most
cases the residual stresses are ignored by the standards
since they are considered as secondary and self-reliev-
ing. However it is the combination of operating-con-
dition stresses and residual stresses that lead to
PWSCC.10 For the case of penetrations attached to the
reactor-vessel heads by partial penetration J-groove
welds, high residual stresses are caused by the surface
machining prior to installation. This machining causes a
thin, strongly deformed layer on the surface, increasing
the material yield and the tensile strength near the
machined surface.
The second source of residual stresses in the J-groove
weld is shrinkage, which occurs when welding the
nozzle into the high-restraint vessel shell, pulls the
nozzle wall outward. This creates yield-strength level
residual-hoop stresses in the nozzle base metal and
higher-strength cold-worked surface layers. These high
residual-hoop stresses contribute to the initiation of axial
PWSCC cracks on the cold-worked surface layer and to
the subsequent growth of these axial cracks in the
lower-strength nozzle base material. Residual stresses in
the nozzles and welds can lead to crack initiation from
the inside surface of the nozzle, opposite from the weld
and from the outside surface of the nozzle near or from
the surface of the J-groove weld.
Based on the industrial experience and laboratory
tests data, the crack growth rate model taking into
consideration the temperature dependence and stresses
was developed.11–16 The recommended crack growth rate
model of detected PWSCC flaws in thick-walled Alloy
600 components exposed to the primary water is:
 exp ( )a
Q
R T T
K K= − −
⎛
⎝
⎜
⎞
⎠
⎟
⎡
⎣⎢
⎤
⎦⎥
−
g
ref
th
1 1
  (2)
where:
a = crack growth rate at temperature T/(m/s)
Qg = 130 kJ/mole activation energy for crack growth
R = universal gas constant 8.314 · 10–3 kJ/(mole K)
T = absolute operating temperature at the location of the
crack
Tref = absolute reference temperature used to normalize
data 598.15 K
 = crack growth amplitude 2.67 × 10–12 at 325 °C
K = crack tip stress-intensity factor (MPa m1/2)
Kth = crack tip stress-intensity factor threshold 9 MPa m1/2
 = exponent 1,16
2.4 Dissimilar metal welds
The conventional welding processes can be used to
produce nickel alloy joints. Some of the characteristics
of nickel alloys require the use of slightly different
welding techniques than normally used for stainless-steel
welds.
Weld Alloys 82 and 182 have been commonly used
to weld Alloy 600 to itself and to other materials. These
alloys are also used for nickel-based alloy weld deposits
(buttering) on weld preparations and for cladding on
areas such as the insides of reactor-vessel nozzles and
steam-generator tube sheets. Alloy 82 with UNS N06082
or W. Nr. 2.4806 is a bare electrode material and is used
for gas tungsten arc welding (GTAW), also known as
tungsten inert gas (TIG) welding. Alloy 182 with UNS
W86182 or W. Nr. 2.4807 is a coated electrode material
and is used in shielded metal arc welding (SMAW). The
compositions of the two alloys are different2, leading to
different susceptibilities to PWSCC. Alloy 182 has a
lower chromium content (13–17 %) than Alloy 82
(18–22 %) and has a higher susceptibility to PWSCC,
probably as a result of the lower chromium content. A
comparison between these weld metals is shown in
Table 3.2
Table 3: Comparison between Alloy 182 and Alloy 82 weld metals
(w/%)
Tabela 3: Primerjava kemijske sestave Alloy 82 in Alloy 182 (w/%)
Ni Cr Fe Ti Ni + Ta C
Alloy 182 59 13–17 
10 
1.0 1.0–2.5 
0.1
Alloy 82 67 18–22 <10 
0.75 2.0–3.0 
0.1
Mn S Si Cu P Co
Alloy 182 5–9.5 
0.015 
1.0 0.5 
0.03 
0.12
Alloy 82 2.5–3.5 
0.015 
0.50 0.5 
0.03 
0.10
The location of dissimilar-metal welds between low
carbon and austenitic steel tubing and piping are shown
in Figure 1. Such transition joints are necessary because
of the corrosion resistance of stainless steel, while
low-carbon steels are commercially more appropriate.17
For example, the reactor pressure vessel and steam
R. CELIN, F. TEHOVNIK: DEGRADATION OF A Ni-Cr-Fe ALLOY IN A PRESSURISED-WATER ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 151–157 155
Figure 7: Dissimilar-metal weld cross-section
Slika 7: Pre~ni prerez zvara med razli~nima kovinskima materialoma
generators are made of low-carbon steels, whereas the
primary piping is made of stainless steel. Therefore, to
join the low-carbon steel components to stainless-steel
piping Alloy 182/82 welding consumables were used
(Figure 3). The macrostructure of a typical multiple-
passes dissimilar-metal weld is shown in Figure 7.
A distinguished columnar pattern of dendrites can be
seen through many weld passes. The dendrites, growing
in the opposite direction to the heat flow, tend to be per-
pendicular to the base material at the weld–base-material
boundary and tend to become vertical (root-to-crown
direction) as the weld thickness increases. The dendrites
in the centre of the weld are mainly vertical.
The boundaries between these similarly oriented
dendrites are called solidification subgrain boundaries
(SSGBs) and tend to have low angular mismatches, as
well as low energy, and are believed to form paths
relatively infrequently for PWSCC only. Where different
sheaves of dendrites intersect or overlap, larger angular
mismatches often occur between the grains. In this case,
the resulting grain boundaries, termed solidification
grain boundaries, can be high energy and are believed to
be more common paths for PWSCC.
The EPRI study18 concluded that PWSCC crack
growth rates for the alloy 82/182 weld metal behave in
accordance with the following relationship:
 expa
Q
R T T
f f K= − −
⎛
⎝
⎜
⎞
⎠
⎟
⎡
⎣⎢
⎤
⎦⎥
g
ref
alloy orient
1 1
  (3)
where:
a = crack growth rate at temperature T in m/s
Qg = 130 kJ/mole activation energy for crack growth
R = universal gas constant 8.314×10–3 kJ/(mole K)
T = absolute operating temperature at location of crack,
K
Tref = absolute reference temperature used to normalize
data 598.15 K
 = power law constant 1.5 × 10–12 at 325 °C (598,15 K)
falloy = 1.0 for Alloy 182 and 0.385 for Alloy 82
forient = 1.0 except 0.5 for crack propagation that is clearly
perpendicular to the dendrite solidification direction
K = crack tip stress-intensity factor (MPa m1/2)
 = exponent 1.6
3 MITIGATION, REPAIR AND REPLACEMENT
3.1 Mitigation
Since PWSCC became a serious issue a number of
techniques have been evaluated to delay or mitigate the
occurrence of degradation processes. These techniques
can be divided into three categories19:
1. Mechanical surface enhancement (MSE),
2. Environmental barriers or coatings,
3. Chemical or electrochemical corrosion potential
(ECP) control.
MSE techniques represent processes that reduce
surface tensile residual stresses or induce compressive
surface stresses on a component or weld surface.
Examples of MSE techniques are shot peening and
electro-polishing. Environmental barrier or coating
techniques represent processes that protect the material
surface in aggressive environments. Coating examples
include nickel plating and weld-deposit overlays.
Chemical or ECP control techniques represent changes
to the environment that alter the corrosion process or
produce corrosion potentials outside the critical range for
PWSCC. Examples of chemical or ECP control include
zinc additions to the primary water and modified pri-
mary-water chemistry (e.g., dissolved hydrogen levels,
lithium concentrations, and boron concentrations). In
some nuclear power plants a component temperature
reduction has also been applied.
3.2 Repair and replacement
Pressure boundary-components repair or replacement
is the alternative to mitigation techniques, especially
when a leakage is detected. ASME Code XI specifies
that the flaws, detected during in-service inspection,
must be removed or reduced to an acceptable size in
accordance with Code-accepted procedures.20 For
PWSCC in Alloy 600/182/82 components, several
approaches have been used, such as flaw removal, flaw
embedment and weld overlay.
For relatively shallow or minor cracking, the flaws
may be removed by grinding. This approach is going to
eliminate flaws and return the component to ASME
Code compliance. However, there will still be sus-
ceptible material exposed to the PWR environment that
originally caused the cracking. Simple flaw removal is
thus not meant to be a permanent solution, unless in
future component replacement is planned.
One of the approaches to repair is to embed the flaw
under a PWSCC-resistant material, typically Alloy 52
weld metal deposited over the susceptible Alloy 182/82
R. CELIN, F. TEHOVNIK: DEGRADATION OF A Ni-Cr-Fe ALLOY IN A PRESSURISED-WATER ...
156 Materiali in tehnologije / Materials and technology 45 (2011) 2, 151–157
Figure 8: J groove weld-flaw embedment repair
Slika 8: Popravilo J-zvara s prekritjem napake
weld surface. The embedment must satisfy all ASME
BPVC Section XI flaw-evaluation requirements.
Another form of repair that has been used extensively
to repair cracked and leaking pipe welds is the weld
overlay (Figure 8). The weld-overlay procedure has
been recognized as a Code-acceptable repair in ASME
Section XI.
In many cases mitigation techniques and/or repairs
were successful. To satisfy Code and regulatory require-
ments, maintenance and inspection programs were spe-
cially developed and implemented for PWSSC-suscep-
tible components. Additional maintenance and
inspection activities mean more inspection costs for
improved techniques to detect PWSCC and extra costs
due to outage-time extension. The replacement of the
components offers many advantages when considering
the long-term operation of nuclear power plant. The new
components are designed to eliminate all the susceptible
materials, replacing them with PWSSC-resistant
materials, such as Alloy 690 and the associated weld
metals Alloys 52 and 152. The main components that
have been replaced are steam generators, reactor vessel
heads and pressurizers.
4 CONCLUSION
Alloy 600 component items were used in pressurized
water reactors (PWRs) due to the material’s inherent
resistance to general corrosion in a number of aggressive
environments and because it has a coefficient of thermal
expansion very close to that of low-alloy steel. Over the
past thirty years, primary-water stress-corrosion cracking
(PWSCC) has been observed in Alloy 600 components
and in Alloy 182/82 welds.
In some cases PWSSC cracks caused relatively quick
and simple plugging of the leaking, small-diameter tubes
in steam generators. On the other hand, many times
PWSSC degradation resulted in long plant outages to
replace leaking pressurizer heater sleeves, a leaking
reactor-vessel hot-leg outlet nozzle weld and several
CRDM nozzles. Many plants around the world have
replaced their steam generators and reactor-vessel
closure heads. The replacements were performed
because of the high cost of repairs, the risk of leakage or
the high cost of inspections necessary to ensure a
satisfactory level of safety and reliability.
5 REFERENCES
1 Materials Reliability Program PWSCC of Alloy 600 Type Materials
in Non-Steam Generator Tubing Applications – Survey report
through June 2002: Part 1: PWSCC in Components Other Than
CRDM/CEDM Penetrations (MRP-87), EPRI, Palo Alto, CA: 2003.
1007832
2 ASME Boiler and pressure vessel Code, Section II: Materials
3 SMC-027, Specials Metals Corporation, 2008, www.specilasmetals.
com/documents
4 Y. S. Garud, T. L. Gerber: Intergranular stress corrosion cracking of
Alloy 600 tubes in PWR primary water – Review and assessment for
model development, 1983, EPRI-NP-3057
5 G. Frederic, P. Hernalsteen, J. Stubbe, Qualification of remedial
methods to prevent primary-side stress corrosion cracking of steam
generator tubing: Volume 3, Global heat treatment, Final report,
1987, EPRI-NP-5249-Vol. 3
6 R. Bandy, D. van Rooyen, Corrosion, 40 (1983), 425–430
7 P. H. Hoang, A. Gangadharan, S. C. Ramalingam, Nuclear Engi-
neering and Design, 181 (1998), 209–219
8 C. Soustelle, M. Foucault, P. Combrade: PWSCC of alloy 600: a
parametric study of surface film effects, Proc. of the 9th Inter. Symp.
on Environmental Degradation of Materials in Nuclear Power
System – Water Reactors, TMS, 1999, p. 105
9 S. A. Attanasio, J. V. Mullen, J. W. Wuthrich, W. Wilkening, D. S.
Morton: Stress corrosion crack growth rates (SCCGRs) for Alloy
182 and 82 Welds, Proc. of the Conf. on Vessel Penetration
Inspection, Crack Growth and Repair, US NRC NUREG/CP-0191,
2005
10 Companion Guide to the ASME Boiler & Pressure Vessel Code,
Third Edition, Ch. 44, Jeff Gorman, Steve Hunt, Pete Riccardella,
Glenn A. White, PWR reactor vessel Alloy 600 issues, 2007
11 W. H. Bamford, J. P. Foster, Crack Growth and Microstructural
Characterization of Alloy 600 PWR Vessel Head Penetration
Materials, EPRI, Palo Alto, CA: 1997. TR-109136
12 T. Cassagne, D. Caron, J. Daret, Y. Lefëvre, Stress Corrosion Crack
Growth Rate Measurements in Alloys 600 and 182 in Primary Water
Loops Under Constant Load, Ninth Inter. Symp. on Environmental
Degradation of Materials in Nuclear Power Systems—Water
Reactors Newport Beach, CA, August 1–5, 1999
13 D. Gómez-Briceño, J. Lapeña, F. Blázquez, Crack Growth Rates in
Vessel Head Penetration Materials, Proc. of the Inter. Symp.
Fontevraud III: Contribution of Materials Investigation to the
Resolution of Problems Encountered in Pressurized Water Reactors,
Chinon, France, September 12–16, French Nuclear Energy Society,
Paris, 1994, p. 209–214
14 M. L. Castaño, D. Gómez-Briceño, M. Alvarez-de-Lara, F. Blázquez,
M. S. Garcia, F. Hernández, A. Lagares, Effect of cationic resin
intrusions on IGA/SCC of Alloy 600 under primary water conditions,
Proc. of the Inter. Symp. Fontevraud IV: Contribution of Materials
Investigation to the Resolution of Problems Encountered in
Pressurized Water Reactors (France, September 14–18, 1998),
French Nuclear Energy Society, Paris, 2 (1998), 925–937
15 G. A. White, J. Hickling, L. K. Mathews, Crack growth rates for
evaluating PWSCC of thick-wall Alloy 600 material, Proc. of the 11th
Inter. Symp. on Environmental Degradation of Materials in Nuclear
Power Systems-Water Reactors, NACE International, (2003),
166–179
16 G. A. White, J. Hickling, C. Harrington, MRP Development of crack
growth rate disposition curves for primary water stress corrosion
cracking (PWSCC) of thick-section Alloy 600 components and Alloy
82, 182 and 132 Weldments, 2005 EPRI International PWSCC of
Alloy 600 Conference, Santa Ana Pueblo, NM, March, 7–10, 2005
17 M. Sireesha, V. Shankar, S. K. Albert, S. Sundaresan, Materials
science and Engineering, A292 (2000), 74–82
18 Materials Reliability Program, Crack growth rates for evaluating
primary water stress corrosion cracking (PWSCC) of Alloy 82, 182,
and 132 Welds (MRP-115NP), EPRI, Palo Alto, CA: 2004. 1006696
19 S. Fyfitch, Alloy 600 PWSCC Mitigation: past, present, and future,
Proc. of the Conf. on Vessel Penetration Inspection, Crack Growth
and Repair, US NRC NUREG/CP-0191, 2005
20 ASME Boiler and Pressure Vessel Code, Section XI: Rules for in
service inspection of nuclear power plant components
R. CELIN, F. TEHOVNIK: DEGRADATION OF A Ni-Cr-Fe ALLOY IN A PRESSURISED-WATER ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 151–157 157

B. CHOKKALINGAM et al.: INVESTIGATION INTO THE MECHANICAL PROPERTIES OF MICRO-ALLOYED ...
INVESTIGATION INTO THE MECHANICAL
PROPERTIES OF MICRO-ALLOYED AS-CAST STEEL
RAZISKAVE MEHANSKIH LASTNOSTI MIKROLEGIRANIH
JEKEL
Bommannan Chokkalingam1, S. S. Mohamed Nazirudeen2, Sukaswami S. Ramakrishnan3
1, 2Department of Metallurgical Engineering, PSG College of Technology, Coimbatore- 641 004. Tamil Nadu, India
3Mechanical Engineering, Karunya University, Coimbatore- 641 114. Tamil Nadu, India
bchokkalingam@gmail.com
Prejem rokopisa – received: 2010-12-22; sprejem za objavo – accepted for publication: 2011-02-23
The effects of micro-alloying on the hardness, tensile strength, room-temperature impact energy and elongation of low-carbon
as-cast steel have been investigated and compared with non-micro-alloyed as-cast steel. Elements such as vanadium, niobium,
titanium and zirconium were added in combination as micro-alloys. The results show that the addition of these elements in the
cast steel increases the hardness and tensile strength by 31 % and 10.3 %, respectively, while the elongation and impact energy
decrease to 47.5 % and 43.7 %, respectively. This method of alloying refines the microstructure of the cast steel in the as-cast
form.
Key words: cast steel, micro-alloying, tensile strength, hardness, toughness
Raziskan je vpliv mikrolegiranja na trdoto, raztr`no trdnost, `ilavost pri sobni temperaturi in razteznost ogljikove jeklene litine
in primerjan z vplivom na litino brez mikrolegiranja. Kombinacije elementov: vanadija, niobija, titana in cirkonija, so bile
dodane v litino. Rezultati ka`ejo, da dodatek teh elementov pove~a trdoto in trdnost do 31 % oz. do 10,6 % in zmanj{a raztezek
za 47,5 %, `ilavost pa za 43,7 %. Mikrolegiranje napravi mikrostrukturo jeklene zlitine bolj drobnozrnato.
Klju~ne besede: jeklena litina, mikrolegiranje, raztr`na trdnost, trdota, `ilavost
1 INTRODUCTION
In general, based on the method of production, steels
can be classified as either wrought or cast steels.
Wrought steel products are shaped by plastic deforma-
tion, whereas cast steel products are produced by casting
methods. Normally, both these processes are used to
produce steel products. In particular, the production of
intricate internal shapes and complicated designs is
easier in cast steel. Moreover, economic production of
fewer quantities is an added advantage of the casting
process over wrought steel processes. However, the
mechanical properties of cast steels are lower when
compared with wrought steels. Hence, the improvement
in mechanical properties is further required for cast steel
due to its wide use in applications like railways,
chemical plants, automobiles and petroleum refineries,
etc.
Generally, the mechanical properties are improved by
alloying. The different types of alloying methods are
low-level, medium-level, high-level and micro-level
alloying. Among these alloying techniques, micro-
alloying1,2 is one of the methods in which the alloying
elements such as vanadium, niobium, titanium and
zirconium are added individually up to 0.10 %. The total
micro-alloying additions of all these alloying elements
combined to be within 0.20 %. This method is widely
practiced in wrought steel processing, whereas it is less
used in cast steel. A study of the combined addition of
vanadium, niobium, titanium and zirconium in cast steel
was not found in the existing literature. Hence in this
work, the effect of a combined addition of all these
alloying elements together in the cast steel was investi-
gated. The microstructure, hardness, tensile strength,
impact energy and elongation of micro-alloyed cast steel
were investigated and compared with those of
non-micro-alloyed cast steel.
2 EXPERIMENTAL PROCEDURE
The melting was performed using a medium-
frequency coreless induction furnace with a holding
capacity of 50 kg and a basic lining. The non-micro-
alloyed and micro-alloyed cast steels’ chemical
compositions are given in Table 1. The mild steel scrap
was melted initially and the alloys ferro-manganese,
ferro-silicon were added into the melt. The micro-
alloyed cast steel composition was achieved by the
addition of the necessary alloys after testing the melt
using a vacuum spectrometer. The vanadium level was
maintained at 0.10 %, while the niobium, titanium and
zirconium additions were maintained at 0.05 %, 0.025 %
and 0.025 %, respectively. The micro-alloying elements
vanadium, niobium and titanium were added into the
molten metal in forms of ferro-vanadium, ferro-niobium
and ferro-titanium, while ferro-zirconium was added in
the ladle during the tapping of the molten metal from the
furnace at 1610 °C. The de-oxidisers aluminium and
Materiali in tehnologije / Materials and technology 45 (2011) 2, 159–162 159
UDK 669.14.018:620.17 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 45(2)159(2011)
calcium silicide were added in the ladle during tapping
the molten metal from the furnace. Both molten steels
were poured into standard Y-block sand moulds.
The metallographic examinations were made with a
light microscope using broken specimens from Charpy
impact tests that were also used in the SEM and EDS
analyses. These analyses were performed using a
scanning electron microscope, JEOL JSM 6360, with an
energy-dispersive spectroscopy (EDS) analyzer (LN2
type detector). The accelerating voltage at the time of the
EDS was 25 kV, the working distance was 17mm and the
probe current was 1.0 nA. The hardness was measured
by using a Zwick hardness tester with an indentation
load of 10 kg. The impact energy was measured with the
Charpy method using standard test pieces (55 mm
length, 10 mm square with V notch (45°, 2 mm deep
with 0.25 mm radius at the base notch). The tensile test
was performed in a Microtek tensile testing machine.
3 RESULTS AND DISCUSSION
3.1 Effects of micro-alloying elements
The purposes of adding micro-alloying elements to
the cast steel are either to get fine grains or to form
precipitates. The various factors contributing to the
increase in hardness of the micro-alloyed cast steel are
variations in the pearlite content, the ferrite grain size
and the formation of fine carbonitride precipitates.
Among these factors, the increase in hardness and
strength is mainly due to carbonitride precipitates. Based
on earlier studies3–6, carbonitride precipitates form in the
matrix, precipitate at the gamma/alpha interphase and
with random precipitation.
The micro-alloying elements titanium and niobium
precipitate at elevated temperatures6–12 and generally,
both elements form a precipitate in the matrix. The
second and third type of carbonitrides formed when there
are sufficient amounts of titanium and niobium present
in the form of a solid solution in austenite during the
transformation stage at the upper critical temperature.
The micro-alloying element vanadium precipitates as
vanadium carbide at lower temperatures on the
advancing gamma/alpha interface. It forms either as an
interphase precipitate or random precipitate in the super-
saturate ferritic matrix. Fine carbonitride precipitates
distributed mainly in ferrite grains are very effective at
strengthening. Earlier microhardness studies on ferrite
grains5–8 revealed that the increase in the strength/
hardness of micro-alloyed cast steels is mainly due to
precipitation hardening, and the effect of solid-solution
strengthening is lower. The solid solution of micro-
alloying atoms is negligible due to the lower concen-
tration of micro-alloying atoms after the formation of
carbonitrides. The increase in the strength and hardness
of vanadium micro-alloyed cast steel is due to the
formation of vanadium carbonitrides V (CN)5,6,10,12. The
increase in the amount of micro-alloying elements
increases the volume fraction of the precipitates.
Zirconium is a nitride and also a strong oxide
forming element. In the solid-state zirconium nitride11,12
has a high melting point and a high hardness. The
zirconium carbide forms at a higher temperature,
whereas the zirconium nitride forms at a lower tempe-
rature. It was found that the zirconium addition improves
the hardness and does not decrease the impact toughness.
The zirconium addition causes sulfide inclusions to
shape as globular rather than elongated, which improves
the toughness and ductility of the micro-alloyed cast
steel.
3.2 Microstructures
The microstructure of non-micro-alloyed and micro-
alloyed as-cast steels is shown in Figure 1 (a) and (b)
B. CHOKKALINGAM et al.: INVESTIGATION INTO THE MECHANICAL PROPERTIES OF MICRO-ALLOYED ...
160 Materiali in tehnologije / Materials and technology 45 (2011) 2, 159–162
Table 1: Chemical composition of experimental cast steels in mass fractions, (w/%)
Tabela 1: Kemi~na sestava preizkusnih jeklenih litin v masnih dele`ih, (w/%)
Experimental Steels
Elements in mass fractions, w/%
C Si Mn P S V Nb Ti Zr
Non-micro-alloyed steel (NMA)
Micro-alloyed steel (MA)
0.23
0.23
0.42
0.42
1.20
1.20
0.02
0.02
0.03
0.03
0.0
0.10
0.0
0.05
0.0
0.025
0.0
0.025
Figure 1: Microstructures of cast steel a) Non-micro-alloyed (NMA),
b) Micro-alloyed (MA)
Slika 1: Mikrostruktura jeklene litine a) brez mikrolegiranja (NMA), z
mikrolegiranjem (MA)
respectively. The microstructure in Figure 1b clearly
shows a finer grain size in the micro-alloyed than in the
non-micro-alloyed cast steel.
The EDS analysis of the micro-alloyed cast steel is
shown in Figure 2. This analysis confirms the presence
of the elements V Nb, Ti and Zr in the micro-alloyed cast
steel.
3.3 Fracture surface
The fractured surfaces of all the broken impact
specimens were examined using scanning electron
microscopy with the aim to understand the mechanisms
of the fracture of both. The fracture surfaces are shown
in Figure 3 (a) and (b). Changes were observed in the
fracture with respect to the size, shape, and depth of the
micro-voids, and the size and shape of the fracture
facets.
The fracture appearance in Figure 3a is an inter-
granular brittle fracture characterized by coarse facets
with numerous very small solidification defects, while
Figure 3b shows a normal brittle fracture with fine
cleavage facets without solidification defects.
The causes of the intergranular fracture in the cast
steel are sulphur distribution, cooling rate, larger grain
size, aluminium and nitrogen contents. The various
methods to control these defects are controlling the
aluminium and nitrogen contents in the liquid steel, the
use of vacuum melting to remove nitrogen to a low level,
the addition of cerium to form stable nitrides, and grain
refinement by the addition of alloying elements such as
titanium, zirconium, vanadium, niobium and boron.
Grain refinement is an effective method among these to
reduce the intergranular fracture and solidification
defects. A finer grain size was obtained by the addition
of micro-alloying elements and solidification defects
were also not found in the fractured surface of the
micro-alloyed cast steel specimen.
3.4 Mechanical properties
The mechanical properties hardness, tensile strength,
room-temperature impact energy and percentage elon-
gation given in Table 2 show the average properties of
various test specimens at different positions. The results
indicate that the addition of micro-alloying increases
significantly the tensile strength and the hardness,
whereas it decreases the percentage elongation and the
room temperature impact toughness considerably. The
increase of hardness and tensile strength to 31 % and
10.3 %, respectively, was observed in the V, Nb, Ti and
Zr micro-alloyed as-cast steel compared to the non-
micro-alloyed as-cast steel. The elongation and
room-temperature impact toughness of the micro-alloyed
as-cast steel decreases to 47.5 % and 43.7 %, respec-
tively, compared to the non-micro-alloyed as-cast steel.
Table 2: Mechanical properties of experimental cast steels
Tabela 2: Mehanske lastnosti preizkusnih jeklenih litin
Experimental steels Hardness,HV
UTS,
MPa
Elonga-
tion, %
Impact
tough-
ness, J
Non-micro-alloyed
steel (NMA) 168 580 17 16
Micro-alloyed steel
(MA 220 640 9 9
4 CONCLUSIONS
The effect of steel micro-alloying with vanadium,
niobium, titanium and zirconium as micro-alloying
elements on the hardness, tensile strength, percentage
B. CHOKKALINGAM et al.: INVESTIGATION INTO THE MECHANICAL PROPERTIES OF MICRO-ALLOYED ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 159–162 161
Figure 3: Fracture analysis of cast steel a) Non-micro-alloyed
(NMA), b) Micro-alloyed (MA)
Slika 3: Povr{ina preloma jeklene litine a) brez mikrolegiranja
(NMA), b) z mikrolegiranjem (MA)
Figure 2: EDS analysis of micro-alloyed cast steel
Slika 2: EDS-analiza mikrolegirane jeklene litine
elongation and room-temperature impact toughness were
investigated in as-cast steel. The results obtained are: as
follows
An increase in hardness to 31 % was obtained for
micro-alloyed cast steel compared with non-micro-
alloyed cast steel and the highest hardness of 220 HV
was obtained for micro-alloyed cast steel.
The tensile strength was increased by 10.3 % in
micro-alloyed cast steel when compared to non-micro-
alloyed cast steel with the highest strength of 640 MPa
obtained.
The elongation and the impact toughness of the
micro-alloyed as-cast steel showed a decrease of 47.5 %
and 43.7 %, respectively, when compared to those of the
non-micro-alloyed as-cast steel.
A smaller grain size was obtained in the micro-
alloyed as-cast steel.
5 REFERENCES
1 S. Seshan: Studies on microalloyed cast steels; Proceedings of the
sixth Asian foundry congress, Calcutta, (1999), 181–187
2 V. Raghavan: Physical metallurgy principles and practice, 2nd ed;
Prentice hall of India limited, New delhi, (2006), ISBN – 81-203 –
3012 – 9, 106, 193
3 B. D. Jana, A. K.Chakrabarti, K. K. Ray: Study of cast microalloyed
steels; Materials science and technology, 19 (1993), 80–86
4 B. Kalandyk, H. Matysiak, J. Glownia: Microstructure strength rela-
tionship in microalloyed cast steels; Reviews on advanced material
science, 8 (2004), 44–48
5 H. Najafi, J. Rassizadehghani, A. Halvaaee: Mechanical properties of
as cast microalloyed steels containing V, Nb and Ti; Material science
and technology, 23 (2007), 699–705
6 J. Rassizadehghani, H. Najafi, M. Emamy, Eslami Saeen: Mecha-
nical properties of V, Nb and Ti bearing as cast micro alloyed steels;
Journal of material science and technology, 23 (2007), 779–784
7 Guang Xu, Xiaolong Gan, Guojun Ma, Feng Luo, Hang Zoe: The
development of Ti alloyed high strength microalloy steel; Materials
and Design, 31 (2010), 2891–2896
8 D. Rasouli, Sh. Kamanah Asl, G. H. Akbar Zadeh, Daneshi: Opti-
mization of mechanical properties of a micro alloyed steel; Materials
and Design, 30 (2008), 2167–2172
9 L. Bejar Gomez, A. Medina Flores, H. Carreon, I. Alfonsa, J. Bernal
Ponce, J. A. Ascencio: Production and characterization of niobium
and titanium microalloyed steels; Revista Maexicana De Fisica , 55
(2009), 110–113
10 Hamidreza Najafi, Jafar Rassizadehghani, Siroos Asgari: As cast
mechanical properties of vanadium/niobium microalloyed steels;
Materials Science and Engineering A, (1–2) (2008), 1–7
11 H. A. Akbarzadeh, M. Tamizifar, Sh. Mirdamadi, A. Abdolhossini:
Mechanical properties and micro structures of Zr – microalloyed cast
steel; ISIJ International, 45 (2005), 1201–1204
12 D. A. Skobir, M. Godec, M. Balcar, M. Jenko: The influence of the
microalloying elements of HSLA steel on the microstructure and
mechanical properties; Mater. Tehnol., 44 (2010), 343–347
B. CHOKKALINGAM et al.: INVESTIGATION INTO THE MECHANICAL PROPERTIES OF MICRO-ALLOYED ...
162 Materiali in tehnologije / Materials and technology 45 (2011) 2, 159–162
K. STRANSKY et al.: THE EFFECT OF ELECTROMAGNETIC STIRRING ON THE CRYSTALLIZATION ...
THE EFFECT OF ELECTROMAGNETIC STIRRING ON
THE CRYSTALLIZATION OF CONCAST BILLETS
KRISTALIZACIJA KONTINUIRNO ULITIH GREDIC V
ELEKTROMAGNETNEM POLJU
Karel Stransky1, Frantisek Kavicka1, Bohumil Sekanina1, Josef Stetina1,Vasilij
Gontarev2, Jana Dobrovska3
1Faculty of Mechanical Engineering, Brno University of Technology,Technicka 2, 616 69 Brno, Czech Republic
2University of Ljubljana, A{ker~eva 12, 1000 Ljubljana, Slovenia
3VSB – Technical University of Ostrava, Tr. 17. listopadu, 708 33 Ostrava, Czech Republic
stransky@fme.vutbr.cz
Prejem rokopisa – received: 2010-10-18; sprejem za objavo – accepted for publication: 2011-02-28
Electromagnetic stirring (EMS) applied on a steel caster (concasting machine) is basically a magneto-hydraulic process that
influences the crystallisation processes and the solidification of billet steel. From the viewpoint of physics and chemistry, the
course of the process is co-determined by a number of relevant parameters, the physical and thermokinetic characteristics of the
concast steel and also the electrical and magnetic quantities. EMS suppresses the growth of columnar crystals of billets and
reduces the tendency to crack during casting and at low temperatures. A caster was used for the testing of two induction stirrers
– one on the actual mould and the other beneath the mould – to determine the effect of EMS on the formation of the structure of
non-alloyed steel. As part of these tests, certain parts of the billets had been cast without the use of stirrers, while other parts
underwent alternate switching on and off of the stirrers for as many as nine combinations of modes. Samples were taken from
the sections of these billets, fine-ground and etched to make the dendritic structure visible. The mode with the highest efficiency
was when both stirrers ran simultaneously. The growth of the columnar crystals, which pointed inward, was limited to ¼-to-1 3 of
the width of the billet when there was no stirring. Experimental research was also confronted with the results acquired from the
application of the models of the temperature field and chemical heterogeneity and the physical-similarity theory.
Keywords: concasting, electromagnetic stirring, dendritic structure, quality of billets, defects
Elektromagnetno me{anje (EMS), ki se uporablja v kontinuirnem livnem stroju, je osnovni magnetno- hidravli~ni proces pri
spremembi kristalizacije in pri strjevanju jeklenih gredic. S stali{~a fizike in kemije proces dolo~ajo ustrezni parametri in
materialne, fizikalne in termokineti~ne zna~ilnosti kontinuirno ulitega jekla in tudi elektri~ni in magnetni vplivi. EMS
prepre~uje rast stebri~astih kristalov v gredicah in zmanj{uje nagnjenje za tvorbo razpok med ulivanjem pri nizkih temperaturah.
Livni stroj je bil uporabljen za preizkus dveh indukcijskih me{al (enega dejansko v kokili in drugega pod njo) za dolo~itev
vpliva EMS na tvorbo strukture nelegiranega jekla. Kot del teh preizkusov so bili dolo~eni deli gredic uliti brez uporabe me{al,,
drugi pa so bili izpostavljeni izmeni~nemu priklopu in izklopu me{ala za ve~ kot devet kombinacij. Vzorci so bili odvzeti po
prerezu gredic, dobro zbru{eni in jedkani, tako da je bila dendritna struktura vidna. Najbolj{i u~inek je bil dose`en pri
isto~asnem delovanju obeh me{a. Rast stebri~astih kristalov, ki so bili usmerjeni v notranjost, je bila omejena na ¼ do 1 3 {irine
gredice, kjer ni bilo me{anja. Eksperimentalne raziskave so bile tudi primerjane z rezultati, ki so bili dose`eni z uporabo
modelov temperaturnega polja, kemijske heterogenosti in teorije fizikalnih podobnosti.
Klju~ne besede: kontinuirno ulivanje, elektromagnetno me{anje, dendritna struktura, kakovost gredic, napake
1 INTRODUCTION
Currently, casters use rotating stators of electro-
magnetic melt-stirring systems. These stators create a
rotating magnetic induction field with an induction of B,
which induces eddy-current J with velocity v in a
direction perpendicular to B. The induction B and the
current J create an electromagnetic force, which works
on every unit of volume of steel and brings about a
stirring motion in the melt. The vector product (v × B)
demonstrates a connection between the electromagnetic
field and the flow of the melt. The speeds of the liquid
steel caused by the EMS are somewhere from 0.1 m/s to
1.0 m/s. The stirring parameters are within a broad range
of values, depending on the construction and techno-
logical application of the stirrer. The power output is
mostly between 100 kW and 800 kW, the electric current
between 300 A and 1000 A, the voltage up to 400 V and
with billet casting the frequency is from 5 Hz to 50 Hz.
The EMS applied on the steel caster is basically a
magneto-hydraulic process affecting the crystallisation
processes and solidification of billet steel. The com-
plexity of the entire process is enhanced further by the
fact that the temperatures are higher than the casting
temperatures of the concast steel. The temperature of the
billet gradually decreases as it passes through the caster
down to a temperature lying far below the solidus
temperature. From the viewpoint of physics and
chemistry, the course of the process is co-determined by
a number of relevant material, physical and thermo-
kinetic characteristics of the concast steel and also the
electrical and magnetic quantities. There is also a wide
range of construction and function parameters pertaining
to the caster and EMS as well as the parameters relating
to their mutual arrangement and synchronisation.
Numerous works from recent years relate that the exact
mathematical modelling of EMS on a caster is still
unsolvable 1–3. The basic EMS experiment was conduc-
Materiali in tehnologije / Materials and technology 45 (2011) 2, 163–166 163
UDK 669.046:620.19 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 45(2)163(2011)
ted on a CONCAST billet caster where two individual
mixers were working, as in Figure 1. The first stirrer,
entitled MEMS (Mould Electromagnetic Stirring), is
mounted directly on the mould and the second stirrer,
entitled SEMS (Strand Electromagnetic Stirring), is
mounted at the beginning of the flow directly after the
first cooling zones but in the secondary-cooling zone.
Here the outer structure of the billet is already created by
a compact layer of crystallites; however, in the centre of
the billet there is still a significant amount of melt that is
mixed by the SEMS.
2 CONDITIONS OF THE EXPERIMENT
The first stirrer (MEMS) stirs the melt still in the
mould while the billet is undergoing crystallization and
solidification. The second stirrer (SEMS) works at a time
when the melt is already enclosed by a shell of cry-
stallites around the perimeter of the billet and inside the
billet there is less melt than above in the active zone of
the first stirrer.
When both stirrers were switched off, the crystalli-
sation and solidification continued in the normal way,
i.e., the solidifying melt did not undergo a forced
rotational movement.
Samples were taken throughout the course of the
experiment – from parts of the billet cast using the
MEMS and SEMS and without and also using either one.
The samples were taken in the form of cross-sections
(i.e., perpendicular to the billet axis). The samples were
fine-ground and etched with the aim of making visible
the dendritic structure which is characteristic for
individual variants of the solidification of the billet.
The verification of the influence of MEMS and
SEMS on the macrostructure of the billet was carried out
on two melts of almost the same chemical composition
(Table 1).
Table 1: Chemical composition of experimental melts [mass fractions,
w/%]
Tabela 1: Kemijska sestava eksperimentalne taline [masni dele`i,
w/%]
Melt C Mn Si P S Cu Cr Ni Al Ti
A 0.14 0.31 0.22 0.014 0.009 0.03 0.05 0.02 0.02 0.002
B 0.13 0.32 0.22 0.018 0.012 0.09 0.06 0.04 0.02 0.002
Table 2: The billet concasting modes and sampling
Tabela 2: Na~ini litja gredic in vzor~evanje
Melt
Concasting
mode
–sampling
Superheating of
steel above
liquidus °C
MEMS
stirring
A
SEMS
stirring
A
Fig.
A
1A 37 210 0
2A 31 0 0 Fig. 5
3A 33 0 29
4A 30 210 57 Fig. 6
B
1B 35 210 0
2B 30 0 0
3B 27 0 57
4B 24 210 57
5B 24 210 29
(Note: Detailed records of the experimental verification of the effects
of MEMS and SEMS during concasting on the relevant device pertain
to Table 2. The data are appended with a time history of the MEMS
and SEMS connection and with information relating to the lengths of
individual billets and the points from which the actual samples had
been taken (i.e., the cross-sections from which the dendritic structures
had been created).
The timing of the concasting process of the billets –
without the involvement of the stirrers and with the
working of the EMS of individual variants of stirrers
(MEMS and SEMS) – is given in Table 2. The speed of
the concasting (i.e., the movement, the proceeding of the
billet through the mould) of the billet was maintained
constant during the experimentation at a value of
K. STRANSKY et al.: THE EFFECT OF ELECTROMAGNETIC STIRRING ON THE CRYSTALLIZATION ...
164 Materiali in tehnologije / Materials and technology 45 (2011) 2, 163–166
Figure 2: Dendrite growth in the concasting structure without EMS –
mode 2A
Slika 2: Rast dendritov v strukturi pri kontinuirnem ulivanju brez
EMS – na~in 2A
Figure 1: The positions of the MEMS and SEMS stirrers
Slika 1: Lega me{al MEMS in SEMS
2.7 m/min. Table 2 shows that as many as nine con-
casting variants were verified. The lengths of individual
experimental billets – from which samples had been
taken – were always a multiple of the metallurgical
length. The average superheating of the steel above the
liquidus was (32.8 ± 3.1) °C in melt A and
(28.0 ± 4.6) °C in melt B, which lies within the standard
deviation of the temperature measurements.
3 EVALUATING EXPERIMENTS
Evaluation of all nine variants of concasting
(Table 2) indicates that the arrangement of the dendrites
in the cross-section follows the same tendency in the first
phase of crystallization. The structure is created by
columnar crystals – dendrites – perpendicular to the
walls of the billet (Figure 2). In the billets that were not
stirred the dendrites gradually touch one another on the
diagonals of the cross-section. Here their growth either
ceases, or the dendrites bend in the directions of the
diagonals and their growth continues all the way to the
centre of the billet. The columnar dendrites that grow
from the middle part of the surface maintain their basic
orientation – perpendicular to the surface – almost all the
way to the centre of the billet. In the central part of the
cross-section there is an obvious hollow on all nine
macroscopic images. This is most probably a shrinkage.
The above-described mechanism of dendrite growth
during concasting without stirring is frequently the
object of interest (Figure 2).
Inside the billets, when using the MEMS stirrer (or
both MEMS and SEMS), the kinetics of solidification
and dendrite growth is initially the same as without
stirring. This also creates columnar dendrites that touch
along the diagonals; however, soon their growth ceases
close to the surface. Dendrites, which are called
equiaxed dendrites, continue to grow – their orientation
is more random and only partly directed towards the
centre of the billet (Figure 6).
It appears that this dendrite growth mechanism
manifests itself the most when both stirrers are working
simultaneously (Table 2: 4A, 4B and 5B). If MEMS and
SEMS are working simultaneously, the stirring effect
significantly prevents the formation of columnar crystals.
If only MEMS is working and SEMS is switched off (1A
and 1B), then the prevention of columnar crystals is less
evident. The working mode of SEMS alone (modes 3A
and 3B) cannot be clearly differentiated from the
changes in the dendritic structure in relation to the
structure formed without stirring (2A and 2B).
Figure 3 (the macro-ground dendritic structure)
shows the depth of the columnar band of dendrites in the
direction away from the surface of the billet (Figure 3 –
see arrows) and its value, which (with the simultaneous
stirring of MEMS and SEMS) is (23.4 ± 1.8) mm. The
same qualified guess was made for the ordinary billet
casting (i.e., without stirring). Here, the depth of
dendrites can be guessed almost all the way to the central
shrinkage at 70 mm (Figure 2 – see arrows). It is known
that additives and impurities during solidification are
often concentrated in the points of contact of the growing
dendrites, where the maximum of segregated additives
and impurities and the greatest probability of technolo-
gical defects occurs.
In the given case, this undesirable effect can be
expected along the diagonals, which have a length of up
to 100-to-103 mm towards the central shrinkage. This
point of contact of the dendrites during the simultaneous
working of SEMS and MEMS is only (29.8 ± 1.9) mm,
i.e., 3.4-times less. The central area of the billet con-
taining a hollow as a result of a shrinkage is then filled
with dendrites growing into a vacuum (i.e., under-
pressure) (Figure 4).
4 DISCUSSION
Under the assumption that the maximum of defects
(i.e., pores, impurities, additives and micro-shrinkages)
K. STRANSKY et al.: THE EFFECT OF ELECTROMAGNETIC STIRRING ON THE CRYSTALLIZATION ...
Materiali in tehnologije / Materials and technology 45 (2011) 2, 163–166 165
Figure 4: Dendrites in the centre of the billet
Slika 4: Dendriti v centru gredice
Figure 3: The growth of dendrites in the billet structure using the
MEMS and SEMS – mode 4A
Slika 3: Rast dendritov v strukturi gredice pri uporabi me{al MEMS
in SEMS – na~in 4A
are formed along the diagonals it is possible to expect
that in the areas of the corners – specifically on the edges
– the nucleation of cracks will be higher than on the
walls of the billet. If the first approximation of the
fracture toughness of the relevant billet made from
low-carbon steel is KIC  75.0 MPa m1/2, then in the
ordinary concasting process it can be assumed that the
length of the contact of columnar dendrites along the
diagonal will be approximately lnormal = 101.5 mm
(Figure 2). On the other hand, if both electromagnetic
stirrers (MEMS and SEMS) are engaged simultaneously,
the contact length of the columnar dendrites along the
diagonal decreases to lel.mg = 29.8 mm (Figure 3).
Along these lengths (i.e., the areas) it could be expected
that during concasting the concentration of the primary
defects will increase where according to the mechanical
fracture theory the following equations should apply for
the preservation of the continuity of the surface:
KIC  	normal πΔ Δl l wnormal normal( / ),
KIC  	el.magn. πΔ Δl l wel.magn. el.magn.( / ).
The first equation applies to normal concasting
without EMS and the second to billet casting with both
MEMS and SEMS engaged simultaneously. The
component (l/w) is the shape factor, which in the first
approximation could be the same in both equations, thus
making it possible to estimate the stress and strain at the
peaks of the dendrites touching each other along the
diagonals.
	 normal
IC
normal
=
⋅
K
lπ πΔ
75
0 0298.
= 245.1 MPa (1)
which is the limit stress and strain for normal concast
billets without EMS, i.e.
	 el.magn.
IC
el.magn.
=
⋅
K
lπ πΔ
75
01015.
= 132.8 MPa (2)
which is the limit stress and strain in the area of the
edges of the billets during concasting if both, the
MEMS and SEMS stirrer, are engaged. A comparison of
both limit stresses and strains indicates that the billets
(otherwise cast under the same conditions) cast without
stirring are almost twice as susceptible to cracking
along the edges as billets cast using both stirrers.
A similar assumption can be made even in the case of
assessing the effect of columnar dendrites in the central
part of the surface of the billet where, without stirring,
their length grows from the surface of the wall all the
way to the central shrinkage (Figure 2), while with the
stirrers the dendrites are significantly shorter. The
boundaries of the dendrites are, however, much less
damaged by technological defects (vacancies, etc.) than
the areas of their touching – of the peaks along the
diagonals. Long-term statistical monitoring of the quality
of (150 × 150) mm billets and the chemical composition
has proven that the application of EMS has significantly
reduced the occurrence of defects (in this case, cracks).
5 CONCLUSIONS
This paper introduces the results of a very demanding
experimental verification of the effect of electromagnetic
stirring (EMS) on the dendritic structure of steel during
the concasting of (150 × 150) mm billets. Nine different
variants of concasting were verified during ordinary
concasting without EMS, with MEMS mounted on the
mould, with SEMS mounted in a secondary cooling zone
and using both MEMS and SEMS.
Macroscopic grinding was conducted on the samples
taken from cross-sections of individual billets in order to
make the dendritic structure visible and evaluate it. The
greatest effect of the EMS was experimentally observed
during the mixing using both MEMS and SEMS
simultaneously. The area of the columnar dendrites
oriented perpendicular to the surfaces of the walls has a
thickness limited to ¼-to-1 3 of the billet thickness. In the
remaining central part of billets stirred in this way the
structure which dominates is the equiaxed dendrite
structure. Long-term statistical monitoring of the quality
of billets has proven that the application of EMS reduces
the occurrence of defects.
Acknowledgments
This analysis was conducted using a program devised
within the framework of the GA CR projects No. 106/
08/0606, 106/09/0940,106/09/0969 and P107/11/1566.
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166 Materiali in tehnologije / Materials and technology 45 (2011) 2, 163–166
K. STRANSKY et al.: THE EFFECT OF ELECTROMAGNETIC STIRRING ON THE CRYSTALLIZATION ...