VSEBINA – CONTENTS
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
The effect of compositional variations on the fracture toughness of 7000 Al-alloys
Vpliv sprememb v sestavi na `ilavost loma aluminijeve zlitine vrste 7000
M. Vratnica, Z. Cvijovi}, N. Radovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 191
Wear mechanism of duplex-coated P/M Vanadis 6 ledeburitic steel
Mehanizem obrabe ledeburitnega jekla P/M Vanadis 6 z dupleksno prevleko
P. Jur~i, M. Hudáková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 197
Analiza toplotnih razpok na orodjih za tla~no litje aluminija
Analysis of thermal cracks on die casting dies
D. Klob~ar, J. Tu{ek, M. Pleterski, L. Kosec, M. Muhi~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 203
Lasersko reparaturno varjenje termorazpok na orodjih za tla~no litje aluminija
Laser repair welding of thermal cracks on aluminium die casting dies
M. Pleterski, J. Tu{ek, L. Kosec, D. Klob~ar, M. Muhi~, T. Muhi~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 211
Use of artificial neural networks in ball burnishing process for the prediction of surface roughness of AA 7075 aluminum alloy
Uporaba umetnih nevronskih mre` za napoved hrapavosti povr{ine pri krogelnem glajenju aluminijeve zlitine AA 7075
U. Esme, A. Sagbas, F. Kahraman, M. Kemal Kulekci . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 215
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Merjenje obrabne obstojnosti strukturne keramike Al2O3
Wear-resistance measurement of structural Al2O3 ceramics
M. Ambro`i~, S. Veskovi~ Bukudur, T. Kosma~, K. Krnel, D. Eterovi~, N. Petkovi~ Habe, I. Pribo{i~ . . . . . . . . . . . . . . . . . . . . . . . . . . 221
DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES . . . . . . . . . . . 227
1. MEDNARODNA KONFERENCA O MATERIALIH IN TEHNOLOGIJAH POD POKROVITELJSTVOM IUVSTA IN FEMS
13. – 15. oktober, 2008, Portoro`, Slovenija
1st INTERNATIONAL CONFERENCE ON MATERIALS AND TECHNOLOGY SPONSORED BY IUVSTA AND FEMS
13–15 October, 2008, Portoro`, Slovenia. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 234
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 42(5)189–233(2008)
MATER.
TEHNOL.
LETNIK
VOLUME 42
[TEV.
NO. 5
STR.
P. 189–233
LJUBLJANA
SLOVENIJA
SEP.–OCT.
2008

M. VRATNICA ET AL.: THE EFFECT OF COMPOSITIONAL VARIATIONS ON THE FRACTURE TOUGHNESS ...
THE EFFECT OF COMPOSITIONAL VARIATIONS ON
THE FRACTURE TOUGHNESS OF 7000 Al-ALLOYS
VPLIV SPREMEMB V SESTAVI NA @ILAVOST LOMA
ALUMINIJEVE ZLITINE VRSTE 7000
Maja Vratnica1, Zorica Cvijovi}2, Nenad Radovi}2
1 Faculty of Metallurgy and Technology, University of Montenegro, 81000 Podgorica, Cetinjski put b. b., Montenegro
2 Faculty of Technology and Metallurgy, University of Belgrade, 11000 Belgrade, Karnegijeva 4, Serbia
majav cg.ac.yu
Prejem rokopisa – received: 2008-02-12; sprejem za objavo – accepted for publication: 2008-04-03
To provide an understanding of how compositional variations affect the microstructural parameters associated with coarse
intermetallic (IM) particles and the fracture toughness in AA 7000 aluminum forgings, a microstructural and fractographic
analysis as well as mechanical tests were carried out on three industrially produced Al-Zn-Mg-Cu alloys with different contents
of impurities (Fe+Si). Light optical microscopy and image analysis were used to assess the volume fraction, size and
distribution of all the soluble and insoluble coarse (> 0.1 µm) IM particles identified in the corresponding R-C and L-R planes
for T73-type heat treatments by selective etching and energy-dispersive X-ray spectroscopy. These quantitative data were
correlated with the plain-strain fracture toughness, KIC, with the results being used to produce useful information on alloy design
and thermomechanical processing via microstructural control. The scanning electron microscope observation of fracture surface
features and an estimation of the area fractions of different fracture modes in the plastic zone segments of a test specimen
showed that multiple failure mechanisms occurred with coarse voiding at the intermetallics becoming more important as the
fraction of coarse IM particles increases. A quantitative assessment of the relevant microstructural and fractographic parameters
will be utilized for developing and verifying a multiple micromechanisms-based model for fracture toughness.
Key words : 7000 Al- alloys, chemical composition, microstructure, fracture toughness
S ciljem, da se ugotovi, kako spremembe v sestavi vplivajo na parametre mikrostrukture, odvisne od velikih delcev
intermetalnih spojin (IM), in na `ilavost loma v izkovkih iz zlitine AA 7000, so bile izvr{ene mikrostrukturne,
mikrofraktografske in mehanske preiskave pri treh industrijskih zlitinah Al-Zn-Mg-Cu z razli~no vsebnostjo ne~isto~ (Fe + Si).
Opti~na mikroskopija in analiza slike sta bili uporabljeni za dolo~itev volumenskega dele`a, velikosti in porazdelitve topnih in
netopnih velikih (>0,1 µm) IM-zrn v ustreznih R-C- in L-R-ploskvah po toplotni obdelavi T73 s selektivnim jedkanjem in z
disperzivno spektroskopijo rentgenskih `arkov. Ti kvantititivni podatki so bili korelirani z `ilavostjo loma KIC, dobljeni pa so
bili tudi podatki, ki so koristna informacija za na~rtovanje zlitin in termomehansko obdelavo s kontrolo mikrostrukture. Analiza
prelomnih povr{in z vrsti~nim mikroskopom in dolo~itev povr{ine razli~nih dele`ev preloma v plasti~ni zoni preizku{ancev je
pokazala ve~ vrst mehanizmov preloma in nastanek tem ve~ velikih jamic ob IM-zrnih, ~im ve~ji je bil dele` teh velikih zrn.
Kvantitativna ocena relevantnih mikrostrukturnih in mikrofraktografskih parametrov bo uporabljena za razvoj in za verifikacijo
modela `ilavosti loma na podlagi ve~ mahanizmov preloma.
Klju~ne besede: aluminijeva 7000 zlitina, kemi~na sestava, mikrostruktura, `ilavost loma
1 INTRODUCTION
High-strength aluminum alloys of the AA 7000
(Al-Zn-Mg-Cu) series are widely used for structural
applications due to their good combination of specific
strength and fracture toughness1. However, the critical
fracture toughness properties, especially in the short
transverse direction, may be seen as questionable, since
the fracture resistance is influenced by a number of
parameters, including a range of microstructural features
that are controlled by the chemistry and processing1,2,3,4.
Furthermore, the microstructural anisotropy associated
with wrought materials may influence the failure mode
depending on the load and crack orientation2.
It is now recognized that the coarse particles of
intermetallic (IM) phases are generally detrimental to the
fracture properties. This is associated with the fact that
although the fracture processes in precipitation-hardened
AA 7000 alloy products involve multiple micromecha-
nisms, the decohesion and fracture of these particles,
which are brittle and have weak interface bonding, is the
first step in a sequence of events that lead to the overall
material fracture1,2,5. The remaining fracture path is
partitioned between intergranular fracture and micro-
void-induced transgranular fracture.
The undesirable coarse particles with sizes in the
range of 1 µm to 20 µm are IM phases of two types: (a)
insoluble Fe- and Si-bearing phases formed during the
solidification process, and (b) normally soluble phases
containing alloying elements that do not completely
dissolve during the homogenization and solution
treatment1. In order to improve the toughness, it is
necessary to achieve the lowest levels of coarse IM
particles. The removal of excess amounts of the soluble
particles is possible by controlling all the stages of
processing. But, the limits on the reduction of the Fe and
Si impurities are set by the cost and the availability of
high-purity materials. Consequently, these impurities are
always present in commercial alloys. They react with Al
and alloying elements such as Mg and Cu to form a large
Materiali in tehnologije / Materials and technology 42 (2008) 5, 191–196 191
UDK 669.715:539.42 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 42(5)191(2008)
number of phases6. In addition, these alloys contain Mn
and Cr, which may also be present in the form of coarse
IM particles, since they combine with Fe, Si and Al.
Therefore, it is of interest to predict the variation in
the fracture properties as a function of the micro-
structural parameters, such as the volume fraction of
coarse IM particles, their size and their spatial
distribution. However, most of the available information
is concerned with the properties of wrought alloys.
Systematic and in-depth quantitative microstructural and
fractographic examinations of commercial AA 7000
alloys in the form of thick plates cut out from forgings
have not been widely conducted. It is the purpose of this
contribution to report on a microstructural and
fractographic investigation of the effect of compositional
variations on the attributes associated with coarse IM
particles and the fracture toughness of modern AA 7000
alloy forgings (the high-zinc variant) in the over-aged
condition as a function of test orientation. The failure
mechanisms are identified and the individual contri-
butions to the overall fracture are quantitatively assessed.
The data are then used to obtain a relationship between
the microstructural parameters and the plane-strain
fracture toughness, with the results being utilized for the
modeling of toughness.
2 EXPERIMENTAL
Three industrial alloys with Zn, Mg, and Cu levels
broadly in the range of the AA 7049 composition were
received in the form of hot-forged  50-mm-thick
pancake-type plates. The chemical composition of each
alloy is given in Table 1. The amounts of alloying
elements are very near to the nominal ones. The only
difference in composition between the alloys is the total
(Fe+Si) content, which increases gradually from alloy 1
through alloy 3. Figure 1 shows the cutting of the tests
specimens from the received plates.
All three alloys were solution treated at 460 °C for
1 h, water-quenched, and aged to a T73 temper. The
two-step T73 over-aged treatment consisted of the aging
of the specimens for 5 h at 100 °C and 5 h at 160 °C.
Light optical microscopy (LOM) and image analyses
were used to characterize the microstructure of the
as-heat-treated plates. Metallographic sections were
taken from the corresponding R-C and L-R planes. The
specimens were then prepared using standard metallo-
graphic techniques. A selective etching and energy-
dispersive X-ray spectroscopy (EDS) analysis on a
scanning electron microscope (SEM) were used to
identify the IM phases present. The volume fraction of
all the soluble and insoluble coarse IM particles, VV, their
size expressed by the average intercept length, L, and the
mean free path, , characterizing the space distribution
were assessed with the line-intercept method. The
measurements were carried out on 500 uniformly
sampled microstructural frames at a magnification of
1000 times. Plane-strain fracture-toughness tests were
performed in accordance with ASTM E399 on the
specimens of the corresponding R-C and L-R orien-
tations, i.e., on the single-edge-notched three-point
bending specimens (SEB) of R-C orientation and the
compact-tension (CT) specimens of L-R orientation. The
specimens were fatigue pre-cracked according to the
ASTM standard specifications. In all cases, three
specimens were tested. The KIC values for the R-C
orientation specimens were obtained from the J-integral
data. JIC was evaluated with the unloading compliance
technique, with a single specimen for each JIC result. The
J-integral and the crack growth, a, were calculated in
accordance with ASTM E1152 and ASTM E813. On a
broken specimen an SEM fractographic examination was
performed to explain the fracture mechanism. The
fracture surface morphology was investigated in the
central region of the plastic zone ahead of the fatigue
pre-crack. The area fraction of the microvoid-induced
transgranular fracture regions, AAt, the intergranular
fracture regions, AAi, and the coarse IM particles, AAp,
were estimated. The area measurements were performed
on SEM fractographs by tracing the areas on a digitizing
M. VRATNICA ET AL.: THE EFFECT OF COMPOSITIONAL VARIATIONS ON THE FRACTURE TOUGHNESS ...
192 Materiali in tehnologije / Materials and technology 42 (2008) 5, 191–196
Figure 1: (a) Schematic illustration of the specimen orientations used for the fracture-toughness tests (L-longitudinal direction, C-circumferential
or tangential direction, R-radial direction) and (b) locations of the metallographic planes for the microstructural analysis (b)
Slika 1: (a) Shema orientacije vzorcev, uporabljenih za preizkuse `ilavosti loma (L-vzol`na smer, C tangencialna smer, R – radialna smer) in (b)
mesto odvzema vorcev za mikrostrukturno analizo
tablet. These measurements provided the data to quantify
the contributions of the different fracture micromecha-
nisms to the plane-strain fracture-initiation process as a
function of the purity degree and the specimen orien-
tation.
3 RESULTS AND DISCUSSION
3.1 Analysis of the microstructural data
Typical microstructures of simple uniaxially forged
material after the full heat treatment are illustrated in
Figure 2.
All the forgings show a deformed dendrite cell
structure with coarse IM particles having an average size
of 1.27–2.43 µm. As expected, relatively coarse and
closely spaced precipitates, mostly situated on the grain
boundary surrounded by the precipitate-free zones
(PFZs), were also observed (Figure 2a). The TEM
characterization of the precipitation in AA 7000 alloys
by previous authors1,2 indicates that these particles are
variants of the -Mg(Cu,Al,Zn)2 phase. In the Al-rich
matrix there was a dense population of uniformly
distributed dispersoids1,2,4 and fine precipitates of the 
and ' phases1,2 that contributed to the precipitation
hardening. On the other hand, the coarse IM particles are
inhomogeneously distributed and aligned in the direction
of the prevailing deformation, as observed on the
metallographic plane of L-R orientation (Figure 2b).
The micrographs also illustrate the IM particles that are
irregularly shaped and of different types (Figures 2c and
d). The combined use of the metallographic and EDS
analyses indicates that these particles are of the
following types: (a) soluble Mg(Cu,Al,Zn)2, S-CuMgAl2,
and, most often observed, Mg2Si, (b) the Fe-bearing
phases Al7Cu2Fe, (Cu,Fe,Mn)Al3 and a very little of
(Cu,Fe,Mn)Al6, (c) the Cr-bearing phase (Cu,Fe,Mn,
Cr)Al7, and (d) another type of Si-containing phases
(Fe,Cr,Mn,Cu)3SiAl12.
The identified phases were found in all three alloys;
however, variations in the composition caused large
changes in their fraction and their morphological charac-
teristics. This observation was also supported by the
image analyses. As can be seen from the data presented
in Table 2, the alloy 3 had the highest percentage of
M. VRATNICA ET AL.: THE EFFECT OF COMPOSITIONAL VARIATIONS ON THE FRACTURE TOUGHNESS ...
Materiali in tehnologije / Materials and technology 42 (2008) 5, 191–196 193
Table 1: Chemical composition of the investigated alloys (in mass fractions, w/%)
Tabela 1: Kemi~na sestava raziskanih zlitin (v masnih dele`ih, w/%)
Alloy Elements
Zn Mg Cu Mn Cr Zr Ti V B Fe Si
1 7.45 2.47 1.53 0.25 0.17 0.15 0.015 0.003 0.003 0.12 0.11
2 7.30 2.26 1.55 0.29 0.18 0.13 0.015 0.007 0.003 0.16 0.09
3 7.65 2.26 1.55 0.25 0.18 0.11 0.017 0.005 0.003 0.26 0.11
Table 2: Results of the image analysis and the plane-strain fracture-toughness tests
Tabela 2: Razultati analize slike in dolo~itve `ilavosti loma
Alloy Plane KIC/(MPa·m1/2) AA*
IM phase characteristics
Type VV,  /% L/µm λ/µm
1
R-C 45.50 ND***
MgZn2+S 0.159 1.63 1024.8
Fe-rich** 0.227 2.00 883.5
Mg2Si 0.125 1.74 1394.1
L-R 43.16
P 0.152 MgZn2+S 0.147 1.56 1054.7
I 0.278 Fe-rich** 0.236 2.08 878.0
T 0.570 Mg2Si 0.144 1.94 1350.0
2
R-C 42.63 ND***
MgZn2+S 0.048 1.29 2685.6
Fe-rich** 0.357 1.99 557.0
Mg2Si 0.094 1.70 1806.7
L-R 40.96
P 0.299 MgZn2+S 0.095 1.27 1333.8
I 0.304 Fe-rich** 0.440 2.06 465.5
T 0.397 Mg2Si 0.134 1.92 1425.1
3
R-C 40.53 ND***
MgZn2+S 0.119 1.63 1363.4
Fe-rich** 0.590 2.43 409.8
Mg2Si 0.147 1.95 1318.7
L-R 37.67
P 0.378 MgZn2+S 0.046 1.38 3031.4
I 0.287 Fe-rich** 0.532 2.37 444.3
T 0.335 Mg2Si 0.146 2.09 1434.8
* Area fractions of microvoid-induced transgranular fracture regions (t), intergranular fracture region (i), and coarse constituent particles (p); **
Fe-rich phases = Al7Cu2Fe + (Cu,Fe,Mn)Al3 + (Cu,Fe,Mn)Al6 + (Cu,Fe,Mn,Cr)Al7 + (Fe,Cr,Mn,Cu)3SiAl12 ; *** ND = not determined
* dele` povr{ine mikrojami~astega transkristalnega preloma (t), interkristalen prelom (i) in velikih IM- zrn (p); ** Z Fe bogate faze = Al7Cu2Fe +
(Cu,Fe,Mn)Al3 + (Cu,Fe,Mn)Al6 + (Cu,Fe,Mn,Cr)Al7 + (Fe,Cr,Mn,Cu)3SiAl12 ; *** ND = ni dolo~eno
coarse particles, with an average of volume fraction of
0.79 %, while the alloys 1 and 2 had a similar volume
fraction, varying between about 0.52 % and 0.58 %. The
significant increase in the amount of coarse particles,
serving as the crack-initiation sites, is a direct
consequence of an increase in the total (Fe+Si) content.
For all the alloys, i.e., the volume fraction of phases
containing Al, Mg, Cu, and Zn (Mg(Cu,Al,Zn)2 and
S-CuMgAl2) is lower than   0.15 %. Note also that,
since the Si content is practically unchanged from one
alloy to the other, the volume fraction of Mg2Si particles
is constant. This implies that the Fe content plays an
important role in the formation of coarse particles. The
volume fraction of grey particles (Fe-containing phases)
increases almost linearly with the increase in the Fe
content. As a result, the coarse particles are distributed
over shorter distances. These features can provide planes
of easy crack growth, thereby reducing the deformation
capacity of the matrix.
3.2 Toughness behavior
Table 2 shows how reducing the total (Fe+Si)
content and thereby removing most of the coarse IM
particles improves KIC. As expected, the toughness was
the highest for the alloy 1, with the lowest (Fe+Si)
content of the mass fraction (w) 0.23 %. By increasing
the impurity level from 0.23 % to 0.25 % and in turn the
volume fraction of the Fe- and Si-containing particles
from the volume fraction () 0.380 % to 0.574 %, the
KIC value in the R-C orientation decreased by approxi-
mately by 6.5 %. Since in alloy 2 the Si content is w =
0.02 % and lower than that in the alloy 1, it is concluded
that large variations in the amount of coarse particles and
the fracture toughness can occur with a relatively small
change in the Fe content. The toughness decreased
further when going to a purity of w = 0.37 %. The drop
in the fracture toughness, due to the presence of
undesirable particles –  = 0.298 % greater than in the
microstructure of alloy 1 – is almost of 11 %.
The trend is similar for the L-R orientation. Although
the amount of damaged particles has the greatest
influence, a more probable explanation for the toughness
degradation found in this work may be based on the
synergistic effect of the coarse particles’ volume
fraction, their sizes and spacing. Namely, as the (Fe+Si)
content increases the particle size increases, while the
mean free path decreases, with a concomitant decrease in
toughness. Thus, although the average particle size vary
between 1.70 µm and 2.08 µm for the alloys 1 and 2, no
distinct variation of the Fe-containing particle size with
the orientation of the metallographic plane were
M. VRATNICA ET AL.: THE EFFECT OF COMPOSITIONAL VARIATIONS ON THE FRACTURE TOUGHNESS ...
194 Materiali in tehnologije / Materials and technology 42 (2008) 5, 191–196
Figure 2: Optical microstructures of the over-aged alloy 3 (a),(c),(d) and alloy 1 (b) observed in R-C (a),(c) and L-R planes (b),(d) etched in 10
% H3PO4 at 50 °C for 5 min (a) and Keller's reagent at 20 °C for 5 s (b), (c), (d). Type of phases: A = Mg(Cu,Al,Zn)2, B = S-CuMgAl2, C =
Mg2Si, D = (Cu,Fe,Mn)Al3 or (Cu,Fe,Mn)Al6, E = Al7Cu2Fe, F = (Cu,Fe,Mn,Cr)Al7, G = (Fe,Cr,Mn,Cu)3SiAl12.
Slika 2: Opti~na slika prestarane zlitine 3 (a), (b), (c) in (d) in zlitine 1 (b) v ploskvah R-C (a) in (c) ter L-R v ploskvah (b) in (d), jedkano v 10 %
HPO4 5 min (a) in s Keller reagentom 5 s pri 20 °C (b), (c) in (d). Vrste faz: A – Mg(Cu,Al,Zn)2, B – S-CuMgAl2, C – Mg2Si, D –
(Cu,Fe,Mn)Al3 ali(Cu, Fe,Mn) Al6, E – Al7Cu2Fe, F – Cu,Fe,Mn,Cr)Al7, G – Fe,Cr,Mn,Cu)3SiAl12
observed, whereas for the alloy 3 the particle size
seemed to vary from 1.95 µm to 2.43 µm, with only
minor differences between the R-C and L-R orientation
planes. It should be noted that the Fe-containing particles
were larger than those of the Mg2Si phase and hence,
more detrimental to KIC, although a contribution from the
other phases particles cannot be ruled out. Namely,
larger particles are cracked more often. Also, the
particles aligned in the loading direction are more prone
to this type of damage. Therefore, the distribution of
coarse particles is of particular importance. The metallo-
graphic examination showed that the spacing of the weak
paths is also affected (at least for the R-C orientation
plane). Decreasing the amount of particles leads to an
increase in the mean interparticle distances. The  value
increases by 60 % to 80 %, going from alloy 3 to alloy 1,
causing the absence of most particle stringers and
decreasing the extent of coarse voiding with the IM part-
icles. The comparison of the R-C and L-R orientations
indicates that for a given degree of purity, the fracture
toughness is different for a different direction of testing.
All three alloys showed a higher toughness for the
R-C-orientation specimens, as compared to that for the
L-R orientation. This toughness anisotropy is parti-
cularly notable for alloy 3, with the highest (Fe+Si)
content. The KIC value in the R-C direction is higher by
7 % than the short-transverse value. This is primarily
attributed to the anisotropic orientations of the coarse IM
particles. This is consistent with the previously reported
results1,5 and indicates that the alignment of these
particles seriously reduces KIC and that it is even more
efficient than the other possible paths of weakness, e.g.,
the grain boundaries with stringers of precipitates. Since
coarse particles are brittle and fracture or separate from
the matrix when the local strain exceeds a critical value,
they provide preferential crack paths ahead of a crack at
a high stress intensity. When such a stress intensity is
applied, the presence of these preferential crack paths
decreases the energy needed for the crack propagation.
Hence, the toughness degradation can be attributed to an
acceleration of the crack initiation and growth by the
damaged particles.
The relative contributions of the different fracture
modes to the plain-strain fracture initiation process (and,
therefore, to KIC) were determined from the fracture
surfaces of the failed toughness specimens. SEM obser-
vations of the fracture surfaces revealed the multi-
mechanism of the fracture process as a competition
between coarse voiding at the IM particles and trans-
granular/intergranular failure. The controlling mecha-
nism varies with the alloy purity, with the combination
of transgranular/intergranular shearing dominating in
alloy 1 (Figure 3a), and the coarse voiding becoming
progressively more important as the (Fe+Si) content
increases (Figure 3b).
The transgranular microvoid-induced fracture gene-
rated with the formation of voids around the matrix
precipitates was observed in all cases. The shear planes
were covered with fine dimples, commonly associated
with very fine particles. Many extremely fine dimples,
representing areas of extensive deformation that
preceded the ductile fracture, were observed on the walls
of very large dimples associated with the aggregates of
mostly fractured IM particles. Places called ridges,
where the crack had changed its direction of propa-
gation, could also be identified. A higher magnification
examination showed that they appear to be formed on
the grain boundaries. Flat intergranular areas were
characterized by a fine population of shallow dimples
and coarse grain-boundary precipitates on some of them.
The presence of the ridges suggests a competition
between intergranular and transgranular shear-fracture
mechanisms, although a greater extent of grain-boundary
failure has not been observed.
In spite of the general similarity of the fracture
surfaces for the alloy 1 through alloy 3, significant
differences were also observed. The crack-propagation
modes and the toughness levels are dependent on the
extent of the primary void growth prior to the outset of
M. VRATNICA ET AL.: THE EFFECT OF COMPOSITIONAL VARIATIONS ON THE FRACTURE TOUGHNESS ...
Materiali in tehnologije / Materials and technology 42 (2008) 5, 191–196 195
Figure 3: SEM micrographs showing the R-C fracture plane of alloy 1
(a) and the L-R fracture plane of alloy 3 (b).
Slika 3: SEM-posnetek preloma v R-C-ploskvi pri zlitini 1 (a) in
L-R-preloma zlitine 3 (b)
the secondary failure mode (i.e., intergranular and
transgranular shear crack growth) and the intervoids
linking. Since the primary voids initiate at coarse IM
particles, it may be expected that an increase in the
(Fe+Si) level, which in turn increases the amount of
these particles, leads to an increased number of fractured
particles, and therefore, a smaller extent of intergranular/
transgranular shear fracture. Figure 2b illustrates this
point very clearly, and it is also confirmed by the
quantitative data in Table 2. The area fraction, AAt, of
the microvoid-induced transgranular fracture decreases
systematically with an increase in the volume fraction of
the Fe- and Si-containing particles, while the area
fraction, AAp, of the coarse IM particles on the fracture
surface increases, leading to a decrease in the fracture
toughness.
4 CONCLUSIONS
Quantitative microstructural and fractographical data
showed that the coarse Fe- and S-containing particles
had a very detrimental effect on the fracture toughness of
over-aged Al-Zn-Mg-Cu forgings. The drop in KIC was
consistent with a change in the dominant failure
mechanism, from intergranular/transgranular shear to
extensive coarse void growth as the (Fe+Si) content
increases from w = 0.23 to 0.37 %. This effect was
attributed to the amount, size and spatial arrangement of
the coarse IM particles and demonstrates the orientation
dependence of the fracture toughness and the morpho-
logical anisotropy of the microstructure.
5 REFERENCES
1 N. U. Deshpande, A. M. Gokhale, D. K. Denzer, John Liu,
Metallurgical Materials Transactions A, 29A (1998) 4, 1191
2 B. Morere, J.-C. Ehrstrom, P. J. Gregson, I. Sinclar, Metallurgical
Materials Transactions A, 31A (2000) 10, 2503
3 N. Kamp, I. Sinclair, M. J. Starink, Metallurgical Materials Transac-
tions A, 33A (2002) 4, 1125
4 R. C. Dorward, D. J. Beerntseen, Metallurgical Materials Transac-
tions A, 26A (1995) 9, 2481
5 A. M. Gokhale, N. U. Deshpande, D. K. Denzer, J. Liu, Metallur-
gical Materials Transactions A, 29A (1998) 4, 1203
6 L. F. Mondolfo, Aluminium alloys, Structure and Properties, Butter-
worts, London 1976
M. VRATNICA ET AL.: THE EFFECT OF COMPOSITIONAL VARIATIONS ON THE FRACTURE TOUGHNESS ...
196 Materiali in tehnologije / Materials and technology 42 (2008) 5, 191–196
P. JUR^I, M. HUDAKOVA: WEAR MECHANISM OF DUPLEX-COATED P/M VANADIS 6 LEDEBURITIC STEEL
WEAR MECHANISM OF DUPLEX-COATED P/M
VANADIS 6 LEDEBURITIC STEEL
MEHANIZEM OBRABE LEDEBURITNEGA JEKLA P/M VANADIS
6 Z DUPLEKSNO PREVLEKO
Peter Jur~i,1 Mária Hudáková2
1ECOSOND Ltd.,Prague 5, Czech republic
2MtF STU Trnava, Slovak republic
p.jurci seznam.cz
Prejem rokopisa – received: 2007-11-21; sprejem za objavo – accepted for publication: 2008-02-05
The wear mechanism of duplex-coated P/M Vanadis 6 ledeburitic steel was investigated. Duplex layering by plasma nitriding
and PVD CrN coating increased the wear resistance substantially; however, the scattering of the results was too large. The
samples differed in the adhesion of the PVD overlay to the nitrided substrate and showed a very different wear resistance and
wear mechanism, too. The specimens with good adhesion showed a critical load needed for coating delamination of more than
130 N, low wear, a relatively little damaged PVD layer, and a rapid wear of counterparts. If the adhesion of the coating was
smaller, the wear rate increased rapidly and the worn surface showed symptoms of massive cracking and a total collapse of the
coating, a strong decrease of the friction coefficient, and of the wear rate of the counterparts.
Key words: ledeburitic P/M steel, duplex coating, wear rate, wear mehanism, coating composition
Raziskan je bil mehanizem obrabe jekla P/M Vanadis 6 z dupleksno prevleko. Dupleksna prevleka iz plazemskega nitriranja in
PVD CrNi-prekritja je pomembno pove~ala odpornost proti obrabi, vendar je bil pri tem prevelik raztros rezultatov. Vzorci so
se razlikovali v adheziji med PVD-plastjo in podlago, zato sta bila razli~na odpornost proti obrabi in mehanizem obrabe. Vzorci
z dobro adhezijo so imeli kriti~no silo za delaminacijo prekritja nad 130 N in relativno majhno obrabo, relativno malo
po{kodovano PVD-plast in hitro obrabo nasprotnega dela. Pri majhni adheziji prekritja, je hitrost obrabe hitro rastla, obrabljena
povr{ina pa je kazala simptome masivnega razpokanja in totalnega kolapsa prekritja, veliko zmanj{anje koeficienta obrabe in
hitrosti obrabe nasprotnega dela.
Klju~ne besede: ledeburitno P/M-jeklo, dupleksno prekritje, hitrost obrabe, mehanizem obrabe, sestava prekritja
1 INTRODUCTION
Ledeburitic steels are widely used in industrial
operations like metal cutting, but also in the cold
working and sheet-metal forming of large series of parts
for the automotive industry. The steels are usually used
in the as-heat-treated condition and with a hardness of
HRC 57–60. Nevertheless, in many cases the heat
treatment itself cannot ensure a sufficient service time
for the tools and they must be submitted to a surface
treatment that also gives them a better resistance to
environmental effects.
Thin layers prepared by various PVD (physical
vapour deposition) methods are mostly used for the
surface treatment of tool steels1-6. PVD layers can be
deposited over a large range of chemical compositions
and mechanical properties, at low temperatures and
without any influence on the core properties of the tools.
Their common disadvantage is the adhesion to the steel
substrate, since PVD layers differ considerably from
steels in terms of mechanical and physical properties.
The surface has to meet many criteria to resist
various mechanisms of wear. The first of these criteria,
and established recently as not the most important, is
hardness1. The ratio of the hardness to the Young’s
modulus was established as being the most important
criterion for sufficient wear resistance1. This ratio differs
significantly for steels and for PVD layers, and may
cause serious problems in the surface processing of
tools. For example, the well-known layers such as TiN,
TiAlN and TiB2 have a hardness of about HV 3000 (four
times more than tool steels) at a Young’s modulus of 400
GPa (2 times higher than steel)3-5. Also, for the
promising CrN layers, the Young’s modulus is higher
than that of steels (up to 240 GPa), whereas the hardness
can be varied according to the actual chemical
composition across a very wide range, from HV 1500 to
approximately HV 25006.
The heat treatment affects the Young’s modulus in a
very limited range. On the other hand, the hardness can
be increased very significantly with an appropriate heat
treatment. For the steel Vanadis 6, the hardness in the
as-delivered state is about HB 255 (HRC 23)8, and after
heat treatment a hardness of HRC 62 (HV 750) can be
easily achieved. An additional surface hardness increase
(and also the hardness/Young’s modulus ratio) can be
achieved by nitriding. As reported9,10 for the Vanadis 6
steel after plasma nitriding, a hardness of HV 1100 can
be achieved. On the other hand, the fracture toughness is
lowered very significantly by the nitriding11,12. The
minimizing of the undesirable effect of nitriding on the
fracture toughness is a significant goal of our other
Materiali in tehnologije / Materials and technology 42 (2008) 5, 197–202 197
UDK 620.178.1:669.14.018.25 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 42(5)197(2008)
investigations. In this paper we report on the results of
an investigation of the wear mechanism of CrN+
plasma-nitriding duplex-layers deposited on a Vanadis 6
substrate.
2 EXPERIMENTAL
Specimens of the steel Vanadis 6 (2.1 % C, 7 % Cr, 6
% V, Fe bal.) were heat treated to a hardness of HRC 60,
the recommended value for cold-work applications.
Plasma nitriding was performed in a RUBIG Micropuls
– Plasmatechnik® device at temperatures of 500 °C and
530 °C and processing times of 60 min and 120 min,
respectively.
The CrN layers of thickness 2 µm or 5 µm were
deposited with vacuum arc sputtering in a Balzers BAI
730M device, and the adhesion of the layers was
examined with a scratch test. The microstructural
examinations and the analysis of the worn surfaces were
performed with scanning electron microscopy. Wear
testing was carried out without lubrication using the
ring-on-plate configuration of the counterpart and the
tested specimen, at loads of 50 N and 150 N, respec-
tively. Hollow cylinders from the ball-bearing steel
100Cr6 heat treated to HRC 60 were used as the counter-
parts. The total movement (sliding in a combination with
slow rotation) distance was of 10 km, and the wear
weight loss was determined also after (1, 2.5 and 5) km
of sliding. Glow discharge optical emission spectroscopy
(GDEOS) was used to determine the elements’ depth
profiles in the near-surface region. The phase consti-
tution of the CrN layers was identified with X-ray
diffraction.
3 RESULTS
In a previous paper13, the best nitriding parameters
with respect to the adhesion of the PVD layers were
found to be: a temperature of 530 °C, a processing time
of 120 min and a reactive atmosphere with the
composition N2 : H2 = 1 : 3. Tribological tests showed a
large scatter in terms of the wear resistance. Therefore, it
was decided to examine other combinations of nitriding
parameters with the aim to reduce the dispersion of
results and consider also the possibility of improving of
the fracture toughness, although this was the main goal
of other investigations.
An example of the microstructure of a duplex-layer
(CrN + plasma nitriding) is shown in Figures 1 and 2.
On the surface, a CrN layer of thickness 5 µm is
deposited. The nitrided inter-layer has a thickness of
about 40 µm. It differs from the non-nitrided material
mainly in the topography caused by the etching. Also,
there is a continuous or quasi-continuous layer of nitride
found at the boundaries of the original grains. No
inhomogeneities were found in the surface regions and
the PVD layer was free of cracks and pores.
As shown in Table 1, all the newly used
combinations of plasma nitriding and CrN layering
produced a minimal adhesion of approximately 50 N.
This indicates that the pre-treatment ahead of the plasma
nitriding is a promising method for increasing the
adhesion of the PVD CrN coating. On the other hand, the
adhesion values again exhibited a large dispersion, even
for specimens processed under the same conditions.
The results of the measurements of the wear
resistance for a load of 50 N in Figure 3, demonstrate
the good wear behaviour of the duplex-treated material
in comparison to a material without surface processing,
only nitrided or only CrN layered9. When compared to
specimens pre-nitrided at 530 °C for 120 min and then
CrN layered14, the results of the tests were not
sufficiently reliable. However, it is clear that the
pre-nitriding at 500 °C for 120 min. is more efficient for
the wear behaviour of the duplex layer than the
pre-nitriding at 530 °C for 60 min. Also, in contrast to
previous tests,9,14 it was confirmed that a thinner CrN
layer had a better wear resistance. This wear behaviour is
to be expected, since internal stresses, often leading to
coating damage, increase with the coating thickness
under given testing conditions15.
P. JUR^I, M. HUDAKOVA: WEAR MECHANISM OF DUPLEX-COATED P/M VANADIS 6 LEDEBURITIC STEEL
198 Materiali in tehnologije / Materials and technology 42 (2008) 5, 197–202
Figures 1, 2: SEM micrographs of duplex layer. Above – overview,
below – detail
Slika 1, 2:. SEM-posnetka dupleksnega prekritja. Zgoraj splo{ni
videz, spodaj detajl
The measurements gave a lower weight loss for the
counterparts after sliding with the CrN layer with a
thickness of 2 µm (Figure 4). No major changes in
weight loss were found, which would indicate that no
delamination of the coatings occurred during the sliding
test. It is concluded that the cause of the different wear
was the differences in the friction coefficient (Figure 5).
A higher friction coefficient induced rather more wear
(Figure 3).
Table 1: Measurements of adhesion with a scratch test. Yellow lines –
results in9,14
Tabela 1: Dolo~anje adhezije s trgalnim preskusom. Rumene vrstice –
rezultati v9,14
Processing Lc3/N Lc4/N Lc5/N
CrN 2 µm 18 44 60
CrN 5 µm 26 56 65
Nitriding 530 °C/120
min + CrN 2 µm 44 70
not
identified
Nitriding 530 °C/120
min + CrN 5 µm 90 94 99
Nitriding 500 °C/120
min + CrN 2 µm
50, not
identified 79, 156 116, -
Nitriding 500 °C/120
min + CrN 5 µm 55, 148
not
identified 129, 158
Nitriding 530 °C/60 min
+ CrN 2 µm 46,135
not
identified,
100
108, 138
Nitriding 530 °C/60 min
+ CrN 5 µm 47, 52
not
identified 98, 134
The use of a higher load (150 N) led to increasing
differences in the wear behaviour between the specimens
processed by various combinations of plasma nitriding
and CrN layering (Figure 6), and also between speci-
mens processed in the same batch (Figure 8). However,
the general tendency of improved wear resistance after
plasma nitriding was retained. The assessment of other
P. JUR^I, M. HUDAKOVA: WEAR MECHANISM OF DUPLEX-COATED P/M VANADIS 6 LEDEBURITIC STEEL
Materiali in tehnologije / Materials and technology 42 (2008) 5, 197–202 199
Figure 3: Weight loss of specimens as a function of sliding distance
and surface treatment, load of 50 N
Slika 3: Izguba mase vzorcev v odvisnosti od dol`ine drsenja in povr-
{inske obdelave; obremenitev 50 N
Figure 6: Weight loss of specimens as a function of sliding distance
and surface treatment, load of 150 N
Slika 6: Izguba mase vzorcev v odvisnosti od dol`ine drsenja in povr-
{inske obdelave; obremenitev 150 N
Figure 4: Weight loss of counterparts as a function of sliding distance
and surface treatment, load of 50 N
Slika 4: Izguba mase nasprotnega dela v odvisnosti od dol`ine drsenja
in povr{inske obdelave; obremenitev 50 N
Figure 5: Friction coefficient as a function of sliding distance, load of
50 N
Slika 5: Koeficient trenja v odvisnosti od dol`ine drsenja; obreme-
nitev 50 N
effects on the wear resistance became more difficult due
to sudden changes in the sliding surface due to the
PVD-layer delamination. Therefore, the mean values
from Figure 6 do not represent the wear behaviour
accurately because of the large dispersion due to
delamination on at least one specimen from the three
tested.
Also, the figures of weight loss for the counterparts
exhibited a large scatter. The mean values in Figure 7
give only partial information about the process during
the wear testing. In the cases of an absence of
delamination, the weight loss of the specimen was
minimal, the counterpart underwent intensive wear
(Figure 8) and the friction coefficient increased slowly.
On the other hand, if the adhesion of the CrN layer was
poor, the wear rate on the specimen’s side changed
suddenly from low to very high, and the weight loss of
the counterpart became minimal. Simultaneously, the
fiction coefficient also decreased as a result of the
decreased surface contact between the specimen and the
counterpart (Figure 9). To understand these differences
in wear behaviour, selected specimens with good (71)
and worse (72) adhesion were subjected to a more
detailed investigation.
These observations correspond well with the
measurements of adhesion from the scratch test,
although the type of contact differs for those during the
scratch test (point contact of the diamond indenter) and
during the wear-resistance test (linear contact of a
hollow cylinder with approximately equal hardness)
differs considerably. As shown already, the adhesion of
50 N, which is the minimum acceptable value for various
industrial applications, was achieved for all the
P. JUR^I, M. HUDAKOVA: WEAR MECHANISM OF DUPLEX-COATED P/M VANADIS 6 LEDEBURITIC STEEL
200 Materiali in tehnologije / Materials and technology 42 (2008) 5, 197–202
Figure 10: Worn surface of the specimen with good adhesion of the
CrN layer
Slika 10: Obrabljena povr{ina vzorca z dobro adhezijo CrN-plasti
Figure 8: Wear of specimen and counterpart in the case of good
(yellow) and poor (red) adhesion
Slika 8: Obraba vzorca in nasprotnega dela za primer dobre (rumeno)
in slabe (rde~e) adhezije
Figure 9: Friction coefficient for the specimen with good adhesion
(yellow) and poor adhesion (red) of the CrN coating
Slika 9: Koeficient obrabe za vzorec z dobro (rumeno) in slabo
(rde~e) adhezijo CrN-plasti
Figure 7: Weight loss of counterparts as a function of sliding distance
and surface processing of the specimen, load 150 N
Slika 7: Izguba mase nasprotnega dela v odvisnosti od dol`ine drsenja
in povr{inske obdelave vzorca; obremenite 150 N
specimens. Large differences in the wear resistance
found for a load of 150 N can be attributed to the
scattering of the adhesion values (ranged from 50 N to
148 N) determined during the scratch test.
Figure 10 shows the surface of a specimen with
excellent PVD layer adhesion. The surface shows two
areas of different micro-morphology. The first area (part
A) is typical for the practically undamaged layer and
only the friction producing parallel tracks in the sliding
direction occurred during the wear test. On the second,
"B", area a slightly deformed and cracked coating
without clear marks of delamination is found. The
preservation of the coating on the surface in spite of the
high contact pressure confirms its excellent adhesion.
In the case of worse adhesion of the coating, the
situation is significantly different (Figure 11). The wear
process was probably equal to the case of good adhesion
(Figure 10) until the occurrence of cracking in the CrN
layer at various places. This was mostly, however, at the
boundary between the worn and the unaffected parts of
the specimen (A). From this moment on the friction
coefficient started to decrease (Figure 9). The
degradation of the layer continued with its crumbling,
predominantly in the vicinity of the primary cracks (B),
up to the total material removal in selected areas (C).
The removed fragments of the CrN coating accelerated
the weight loss on the specimen’s side due to the
removal itself, while on the counterpart’s side, the
weight loss was practically stopped due to minimizing of
the sliding contact (minimizing of the surface
undergoing the wear) and lowering of the friction
coefficient (Figure 9).
Specimens with good (71) and poor (72) wear
behaviour also differ in the chemistry of the PVD layers
and, as a consequence, in the phase constitution, too. The
CrN layer on specimen 71 contains less chromium than
that on specimen 72 (Figures 12 and 13), but it has
more nitrogen. The difference is 1.5–2.0 %, depending
on the depth. At the same time, the layer on specimen 71
consists mainly of the Cr2N compound, while in the
layer on specimen 72, the content of nitrogen in the
chromium solid solution is significantly increased
(Figure 14). The specimens also differ in terms of the
texture of the layer. The diffraction peaks of Cr2N(110) at
37° of the two-theta angle and Cr2N(300) at 67° are higher
for specimen 72, with a poor adhesion, than for
specimen 71, with good adhesion. The differences
indicate a different mechanism of growth for both layers.
At first sight this is surprising, since the layers were
prepared on specimens processed in the same way as in
the heat-treatment stage and in the nitriding stage.
Another possible explanation of the different wear
behaviour of the specimens, and also of those processed
under the same processing conditions, is a possible
exceeding of the load-carrying capacity of the substrate.
Nevertheless, the scratch test indicated a large scatter of
the critical load values and, therefore, the differences in
the layer formation displayed by X-ray diffraction
P. JUR^I, M. HUDAKOVA: WEAR MECHANISM OF DUPLEX-COATED P/M VANADIS 6 LEDEBURITIC STEEL
Materiali in tehnologije / Materials and technology 42 (2008) 5, 197–202 201
Figure 13: Depth profiles of elements throughout the near surface
region – specimen with poor adhesion
Slika 13: Porazdelitev elementov v CrN-plasti blizu povr{ine, vzorec
s slab{o adhezijo
Figure 12: Depth profiles of elements throughout the near-surface
region – specimen with good adhesion
Slika 12: Porazdelitev elementov v CrN-plasti blizu povr{ine, vzorec
z dobro adhezijo
Figure 11: Worn surface with poor adhesion of the CrN layer
Slika 11: Obrabljena povr{ina vzorca s slabo adhezijo CrN-plasti
measurements are more probably the primary expla-
nation for the different wear resistance of the specimens.
The cause of different layer growth, in a given batch,
too, can probably be attributed to the positional effect of
the specimen with respect to the chromium source
(target) during the sputtering. No other effects with a
possible influence on the layer growth are known.
4 CONCLUSIONS
1) For all the combinations of plasma nitriding and
CrN coating, the critical load measured with the scratch
test was at least 50 N, which is considered to be an
acceptable value for industrial applications. In some
cases the adhesion was much better and the critical load
exceeded 130 N. However, a large scatter of adhesion
values was found, also for specimens processed under
the same nitriding and/or coating conditions.
2) Wear testing at a load of 50 N did not show a
significant scatter of results. The wear of specimens
proceeded in a steady manner and slightly better results
were found for the thinner CrN layers.
3) The use of a load of 150 N emphasized the
difference in the wear behaviour of the specimens, also
for those prepared under the same processing conditions.
For the specimens with excellent adhesion, the wear rate
on the specimen’s side was low and on the counterpart’s
side it was much greater. The friction coefficient
increased slightly with the increased sliding distance. For
the specimens with poor adhesion, a sudden and
dramatic increase in the wear rate on the specimen’s side
occurred during the test. This increase is connected with
the decrease of the friction coefficient and the lowering
of the wear on the counterpart’s side.
4) In the case of specimens with excellent adhesion
the CrN coating remained on the surface, although in a
slightly damaged form. Coatings with small adherence
cracked and were progressively removed from the
contact surface.
5) Layers with excellent adhesion differ from those
with poor adhesion in chemistry and phase constitution.
In the first there was less chromium in the solid solution
and more Cr2N compound was found in the second.
Moreover, the different texture probably also indicates a
difference in the growth mechanism.
ACKNOWLEDGEMENTS
Authors wish to thank the Ministry of Education and
Youth of the Czech Republic and the Slovak Republic
for the financial support for the solution of the Project
Eureka E!3437 PROSURFMET.
5 REFERENCES
1 Oberle, T.: J. of Metals, (1951) 3, 438–439
2 Matthews, A., Leyland, A.: HTM 56 (2001) 1, 5
3 Kí`, A., Zetek, M.: In.: Proc. of the Conf. Vrstvy a povlaky 2004,
Ro`nov pod Radho{tm, 7.–8. 10. 2004, 12
4 [míd, M. et al.: In: Proc. of the Conf. Vrstvy a povlaky 2004,
Ro`nov pod Radho{tm, 7.–8. 10. 2004, 79
5 Podgornik, B. et al.: Wear 254 (2003), 1113
6 Sokovi~, M., Panjan, P., Kirn, R.: J. Mater. Proc. Techn., 157–158
(2004), 616
7 Salas, O. et al.: Surf. Coat. Techn. 172 (2003), 117
8 Vanadis 6 – SuperClean, Uddeholm Tooling AB, Sweden, 2001
9 Jur~i, P., Panjan, P.: Surface Processing of the PM Vanadis 6 Steel
with Plasma Nitriding and CrN PVD – Coating, In: Proceedings of
the European Powder Metallurgy Congress, Prague, Czech
Republic, 2005
10 Jur~i, P., Hnilica, F., Suchánek, J., Stola, P.: Structural Features of
the Cr-V Ledeburitic Steel Saturated by Nitrogen, Mater. Tehnol., 38
(2004) 1–2, 13
11 Hnilica, P., ^makal, J., Jur~i, P.: Alterations in Fracture Behaviour
of the Cr-V Ledeburitic Steel Vanadis 6 Caused by Plasma Nitriding,
In.: 11th Conference on Materials and Technology, book of abstracts,
Portoro`, Slovenia, October 1–3, 2003
12 Jur~i, P., Hnilica, F., ^makal, J., Pechmanová, J.: In: Proceedings of
the 8th National Conference Degradation of Structural Materials,
September 2.-4, 2003, Terchová, Slovakia, p. 137
13 Jur~i, P., Musilová, A., Stola, P., Hrubý, V.: In: Proceedings of the
Conference Carburizing and Nitriding, November 27–28, 2001,
Brno, Czech Republic, p. 81
14 Jur~i, P., Hudáková, M.: Surface Processing of the PM Vanadis 6
Steel with Plasma Nitriding and CrN PVD – Coating, In: Pro-
ceedings of the 16th Int. Conf. Metal 2007, May 2007, Hradec n.
Moravicí, Czech Republic, CD – ROM
15 Diao, D. F. et al.: The Maximum Tensile Stress on a Hard Coating
Under Sliding Friction. Tribology Intern., 27 (1994) 4, 267–272
P. JUR^I, M. HUDAKOVA: WEAR MECHANISM OF DUPLEX-COATED P/M VANADIS 6 LEDEBURITIC STEEL
202 Materiali in tehnologije / Materials and technology 42 (2008) 5, 197–202
Figure 14: X-ray diffraction patterns from PVD CrN layers
Slika 14: Difraktogram rentgenskih `arkov v PVD-plasti
D. KLOB^AR ET AL.: ANALIZA TOPLOTNIH RAZPOK NA ORODJIH ZA TLA^NO LITJE ALUMINIJA
ANALIZA TOPLOTNIH RAZPOK NA ORODJIH ZA
TLA^NO LITJE ALUMINIJA
ANALYSIS OF THERMAL CRACKS ON DIE CASTING DIES
Damjan Klob~ar1, Janez Tu{ek1, Matej Pleterski1, Ladislav Kosec2, Mitja Muhi~3
1Fakulteta za strojni{tvo, Univerza v Ljubljani, A{ker~eva 6, 1000 Ljubljana, Slovenija
2Naravoslovnotehni{ka fakulteta, Univerza v Ljubljani, A{ker~eva 12, 1000 Ljubljana, Slovenija
3TKC d.o.o., Trnovska 8, 1000 Ljubljana, Slovenija
Prejem rokopisa – received: 2007-09-21; sprejem za objavo – accepted for publication: 2008-02-05
Podalj{anje trajnostne dobe orodij je ekonomsko zelo pomembno, zato smo raziskali po{kodbe orodij za tla~no litje
aluminijevih zlitin. Analizirali smo vzroke in mehanizme nastanka razpok, ki nastanejo zaradi temperaturnega utrujanja med
tla~nim litjem. Izdelali smo napravo za potopni preskus, ki omogo~a simulacijo pogojev med tla~nim litjem teh zlitin in
kontrolirano temperaturno utrujanje jekla. Epruvete iz jekla H13 so bile razli~no toplotno obdelane. Na vogale nekaj epruvet je
bilo po postopku TIG navarjeno jeklo maraging, ki je bilo presku{ano v navarjenem stanju in optimalno staranem stanju. Po
dolo~enem {tevilu ciklov so bile povr{ine epruvet vizualno pregledane, vzorci pa so bili razrezani in pregledani na svetlobnem
in vrsti~nem elektronskem mikroskopu. Podro~ja okoli razpoke so bila analizirana z energijsko disperzijsko spektroskopijo
rentgenskih `arkov (EDS). V vogalu epruvete smo izmerili {tevilo in velikost razpok ter padec trdote. Rezultati so pokazali, da
je odpornost jekla proti temperaturnemu utrujanju odvisna od vrste jekla in njegove toplotne obdelave. Dobro odpornost proti
temperaturnemu utrujanju ima kakovostno toplotno obdelano jeklo H13. [irjenje povr{inskih razpok dodatno pospe{i oksidacija
povr{ine.
Klju~ne besede: tla~no litje, aluminijeve zlitine, termi~no utrujanje, potopni preskus, orodno jeklo za delo v vro~em, jeklo
maraging
The aim of this research was to analyse the aluminum alloy die casting die failures with the aim of prolonging in-service die
life. An extensive analysis of thermal fatigue cracks in aluminium alloy die casting was performed. An immersion test apparatus
was developed that enables the simulation of conditions during aluminum alloy die casting and the controlled thermal fatigue
testing of materials. Specimens of the AISI H13 tool steel were differently heat treated. On the edges of some specimens, the
maraging steel was gas tungsten arc (GTA) weld cladded. These samples were tested in as-welded and in optimally aged
condition. After completion of a particular number of cycles, the specimen edge surface was examined visually. The specimens
were then cut and examined with light and scanning electron microscopy (SEM). The energy-dispersive X-ray spectroscopy
(EDS) mapping of areas around the cracks was performed, also, and the number and length of thermal fatigue cracks were
assessed and Vickers hardness profiles were determined. The results confirmed the good thermal fatigue resistance of the
properly hardened H13 steel. The surface cracks growth is accelerated by surface oxidation.
Keywords: aluminium alloy, die casting, thermal fatigue cracking, immersion test, hot work tool steel, maraging steel
1 UVOD
Tla~no litje aluminijevih zlitin je velikoserijski
postopek izdelave izdelkov zahtevnih geometrijskih
oblik v ozkih dimenzijskih tolerancah. Med tla~nim
litjem prite~e v orodje talina s hitrostjo od 30 m/s do 100
m/s in s temperaturo okoli 700 °C. Polnilni tlaki so od 50
MPa do 80 MPa 1. Trajnost orodij zmanj{ujejo a) tem-
peraturni cikli, ki povzro~ajo temperaturne razpoke na
povr{ini orodja zaradi temperaturnega utrujanja jekla, b)
korozija in sprijemanje zaradi oksidacije aluminija na
povr{ino gravure 2,3, c) erozija zaradi toka taline, ~) pre-
lomi orodja zaradi temperaturnih {okov in d) segrevanje
jekla, ki povzro~a nestabilnost mikrostrukture in
degradacijo mehanskih lastnosti 4-7.
Orodja za tla~no litje aluminija se najpogosteje
po{kodujejo zaradi razpok, ki nastanejo zaradi tempe-
raturnega utrujanja jekla. Razpoke, ki se pojavijo v
orodju, povzro~ajo nesprejemljive odtise na povr{ini
izdelka.
Trajnost orodja lahko podalj{amo a) z zmanj{anjem
temperaturnih napetosti med obratovanjem orodja, b)
izbiro jekla in toplotno obdelavo, ki zagotavljata ve~jo
odpornost proti temperaturnemu utrujanju, c) s povr{in-
skimi prevlekami na orodjih in ~) reparaturnim
varjenjem, s katerim saniramo po{kodovano orodje.
Delovne napetosti v orodju zmanj{amo z bolj blagimi
parametri tla~nega litja, optimalnim notranjim segre-
vanjem orodja (temperiranjem), z optimalno geometrijo
izdelka/orodja s ~im manj ostrimi prehodi ter z minimal-
nim vnosom zaostalih napetosti med njegovo izdelavo.
Povr{inske prevleke uspe{no zmanj{ujejo korozijo,
spajanje in erozijo povr{ine gravure, vendar niso odporne
proti temperaturnemu utrujanju 8. Novo razvite ve~slojne
prevleke `e dosegajo izbolj{ano odpornost proti
temperaturnemu utrujanju 9,10.
Poleg klasi~nih CrMoV-jekel za delo v vro~em, imajo
dobro vzdr`ljivost tudi orodja, izdelana iz jekel maraging
z ni`jo vsebnostjo niklja (od 10 % do 14 %) 11,12. Njihova
prednost pred klasi~nimi orodnimi jekli za delo v vro~em
sta manj{i modul elasti~nosti in manj{i temperaturni
raztezek, kar oboje zmanj{uje temperaturne napetosti.
Ve~ja toplotna prevodnost jekel maraging pove~a
prevajanje toplote po materialu in s tem zmanj{uje
Materiali in tehnologije / Materials and technology 42 (2008) 5, 203–210 203
UDK 669.14.018.252:669.715:539.42 ISSN 1580-2949
Izvirni znanstveni ~lanek/Original scientific article MTAEC9, 42(5)203(2008)
temperaturne gradiente in napetosti, kar lahko prispeva k
dalj{i obratovalni dobi orodja.
Razvoj orodnih jekel za delo v vro~em gre v smeri
izbolj{anja stabilnosti mehanskih lastnosti v obratovanih
razmerah 12. K temu veliko prispevajo toplotne obdelave,
ki prepre~ujejo oz. zavirajo spremembe mikrostrukture z
oblikovanjem bolj stabilnih karbidnih izlo~kov
(M(CrMoV)C) med popu{~anjem. Naloga teh izlo~kov
je zaviranje gibanja dislokacij, kar pove~uje odpornost
jekla proti meh~anju 13-15.
Raziskave na podro~ju napovedovanja trajnosti
orodij 13,14,16,17 in opisa mehanizmov po{kodb 5,18,19 pote-
kajo v realnih razmerah na orodjih med obratovanjem ali
v laboratorijskih razmerah. Ugotavljanje odpornosti jekla
proti temperaturnemu utrujanju izvajamo s preizkusi
izotermnega utrujanja (pri konstantni temperaturi se
spreminja obremenitev), termomehanskega utrujanja in
temperaturnega utrujanja v najrazli~nej{ih izvedbah
14,20-22.
2 EXPERIMENTALNI DEL
Preizku{ali smo jeklo 1.2344 (kemi~ne sestave: 0,4
% C; 5,15 % Cr; 1,05 % Si; 0,4 % Mn; 1,35 % Mo in
1 % V) ter vare iz jekla maraging 1.6356 (kemi~ne
sestave: 0,02 % C; 1,35 % Mo; 18 % Ni; 12 % Co; 0,1 %
Al in 1,6 % Ti).
2.1 Potopni preizkus
Naredili smo napravo za potopni preizkus (PP)
temperaturnega utrujanja, ki omogo~a kontrolirano
temperaturno utrujanje jekla (slika 1a). Cikel je trajal 21
s in je zajemal trisekundno dr`anje epruvet v emulziji
(32 °C), ki prepre~i sprijemanje aluminija na epruvete,
{tirisekundno gibanje epruvet do posode z raztaljeno
aluminijevo zlitino, desetsekundno dr`anje epruvet v tej
zlitini s temperaturo   700 °C in {tirisekundno vra~anje
epruvet do posode z emulzijo. Epruvete so bile vseskozi
notranje hlajene z vodo s temperaturo 20 °C. To
povzro~a velike tla~ne napetosti na povr{ini epruvete, ko
je le-ta potopljena v talino aluminijeve zlitine, in natezne
napetosti, ko je epruveta potopljena v hladno emulzijo.
Najve~je napetosti se pojavijo v vogalu epruvete.
Velike tla~ne napetosti povzro~ijo v vogalu epruvete
plasti~no deformacijo, saj napetosti krepko presegajo
tiste pri te~enju jekla pri tej temperaturi 24. V hladni fazi
cikla, ko je epruveta pomo~ena v hladno emulzijo,
nastanejo natezne napetosti, ki povzro~ajo nukleacijo
razpok. Razpoke nato rastejo zaradi malocikli~nega
utrujanja.
2.2 Oblika in priprava epruvet
Preizkusili smo dve obliki epruvet: klasi~no za
potopni preskus (slika 1b) in epruveto z optimalno
oblikovanim vogalom (slika 1c). Slednjo smo razvili na
osnovi analize napetosti z metodo kon~nih elementov
(MKE) in omogo~a ve~je temperaturne gradiente oz.
napetosti med potopnim preskusom ter bolj zahtevne
pogoje presku{anja. Z navojem M16 je bila epruveta
pritrjena na pnevmatski manipulator.
Epruvete so bile izdelane iz jekla H13 in toplotno
obdelane, tj. kaljene in popu{~ene (tabela 1). Toplotna
obdelava 1 pomeni kakovostno toplotno obdelavo s
kaljenjem v vakuumu, toplotna obdelava 2 pa toplotno
obdelavo slab{e kakovosti, s kaljenjem na zraku. Na
nekaj epruvet, kaljenih v vakuumu, smo po postopku
TIG navarili jeklo maraging. Varjenje je potekalo v
za{~itni atmosferi plina argona s pretokom 10 L/min.
Varilni tok je bil 82 A, varilna napetost 11 V, ~as
varjenja je bil okoli 150 s, vnos energije pa okoli 1185
J/mm. Nekaj epruvet smo po varjenju starali 3 h pri
temperaturi 480 °C, druge pa so bile presku{ene v
navarjenem stanju. Epruvete so bile mehansko obdelane
(slika 1c).
D. KLOB^AR ET AL.: ANALIZA TOPLOTNIH RAZPOK NA ORODJIH ZA TLA^NO LITJE ALUMINIJA
204 Materiali in tehnologije / Materials and technology 42 (2008) 5, 203–210
Slika 1: a) Shematski prikaz naprave za termi~no utrujanje materiala, b) klasi~na epruveta za presku{anje in c) optimalna oblika epruvete
Figure 1: a) Scheme of the thermal fatigue test apparatus, b) shape of the classic thermal fatigue test specimen, and c) optimal test specimen
geometry
2.3 Presku{anje in vrednotenje podatkov
Z epruvetami smo naredili 20 000 preizkusov
temperaturnega utrujanja. Robove epruvet smo po vsakih
4 000 ciklih vizualno pregledali, ali so razpokani. Po 20
000 ciklih smo epruvete razrezali preko vogala vzdol`
njene dol`ine (slika 4 a) in naredili mikroobruse. Vzorce
smo jedkali v 4-odstotni raztopini nitala. Analizirali smo
mikrostrukturo, izmerili profil mikrotrdote po Vickersu
od vogala epruvete proti centru ter na svetlobnem
mikroskopu fotografirali in v programu Image Tool
analizirali dol`ino in {tevilo razpok vzdol` prerezanega
vogala epruvete. Izmerjene vrednosti smo statisti~no
obdelali. V vrsti~nem elektronskem mikroskopu (SEM)
Jeol-JSM 5610 smo z elektronsko mikroanalizo izmerili
kemi~no sestavo materiala na povr{ini in v okolici
razpoke.
2.4 Analiza z metodo kon~nih elementov
Z MKE in s programom ABAQUS smo simulirali
temperaturno utrujanje 23. Naredili smo sekven~no veza-
no termomehansko analizo in dolo~ili optimalno obliko
epruvete ter zahtevne pogoje potopnega preizkusa 22.
3 REZULTATI
3.1 FEM-analiza
Toplotne napetosti (slika 2 b) med potopnim pre-
skusom so posledica temperaturnih gradientov (slika 2
a) med povr{ino in notranjostjo epruvete. Najve~ji gra-
dienti temperature oz. napetosti se pojavijo na povr{ini v
za~etku pomakanja epruvete v vro~o talino in hladno
emulzijo. Kasneje se le-ti zmanj{ajo zaradi prenosa
toplote v notranjost epruvete. Posledice cikli~nega
obremenjevanja epruvet so bile vidne kot razpoke na
povr{ini.
3.2 Povr{inske razpoke
Majhno {tevilo dolgih razpok se je pojavilo v jeklu
H13, kaljenem na zraku (slika 3 a). V enakem jeklu,
kaljenem v vakuumu, se je pojavilo ve~je {tevilo kraj{ih
razpok (slika 3 b). V jeklu maraging v navarjenem sta-
nju (slika 3 c) in staranem stanju (slika 3 d) se je poja-
vilo veliko {tevilo kratkih in komaj opaznih razpok.
3.3 Razpoke v prerezu vogala
Razpoke v epruveti iz jekla H13, kaljeni na zraku, so
bile velike in redke (slika 4 b in 3 a). Njihova velikost je
bila posledica razoglji~enja povr{ine pri segrevanju pred
kaljenjem in ima manj{o trdoto (slika 5) in slab{e
mehanske lastnosti. ^ela (vrhovi) razpok so bila najbolj
pogosta na prehodu delno razoglji~enega v naoglji~eno
jeklo (slika 4 b). V vogalu epruvete iz jekla H13, ki je
bila toplotno obdelana v vakuumu, so bile razpoke bist-
veno kraj{e (slika 4 c). Podobno velja tudi za povr{inske
navare jekla maraging v navarjenem (slika 4 d) in stara-
nem stanju (slika 4 e).
3.4 Trdota jekla na povr{ini epruvete
Razen pri epruveti iz jekla H13, ki je bila kaljena v
vakuumu, se je pomemben padec trdote pojavil pri vseh
D. KLOB^AR ET AL.: ANALIZA TOPLOTNIH RAZPOK NA ORODJIH ZA TLA^NO LITJE ALUMINIJA
Materiali in tehnologije / Materials and technology 42 (2008) 5, 203–210 205
Slika 3: Mikrorazpoke na povr{ini vogala epruvet po 20 000 ciklih: a)
jeklo H13, kaljeno na zraku, b) jeklo H13, TO v vakuumu, c) jeklo
maraging v navarjenem stanju in d) navarjeno jeklo maraging starano
3 h na temperaturi 480 °C
Figure 3: Micro-cracks on the edge surface after 20 000 cycles of
immersion: H13 tool steel quenched in a) air and b) vacuum and
maraging steel in c) as-welded condition and d) welded and aged 3 h
at 480 °C
Slika 2: MKE-analiza a) temperatur in b) napetosti v epruveti med
potopnim preskusom. Parametri potopnega preskusa: TAL = 690 °C,
tAL = 10 s, TZRAK = 28 °C, tZRAK = 4 s, TEMULZIJA = 32 °C, tEMULZIJA
= 3 s, THLADILNA VODA = 20 °C, tCIKLA = 21 s
Figure 2: Finite element analysis of a) temperatures and b) stresses in
test specimen during the immersion test. Immersion test parameters:
TAL = 690 °C, tAL = 10 s, TAIR = 28 °C, tAIR = 4 s, TEMULSION = 32
°C, tEMULSION = 3 s, TCOOLING WATER = 20 °C, tCYCLE = 21 s
analiziranih epruvetah (slika 5). Najve~je zmanj{anje
trdote (∆HV   280) smo opazili v jeklu H13, kaljenem
na zraku (slika 4 b). Nekoliko manj{a padca trdote sta se
pojavila v navarih iz jekla maraging (∆HV   180) kot
posledica prestaranja.
3.5 [tevilo in velikost razpok
Najdalj{e razpoke so bile v jeklu H13, kaljenem na
zraku (slika 6). Njihova povpre~na dol`ina je bila 72,8
µm, najdalj{e pa so bile dolge 230 µm (slika 6 a).
Povpre~na gostota razpok v vogalu je bila okoli 4,7
razpoke na milimeter dol`ine vogala. Razpoke v drugih
epruvetah so bile bistveno kraj{e. Pri jeklu H13, ki je
bilo kaljeno v vakuumu, je bila povpre~na dol`ina 2,3
µm, pri jeklu maraging v navarjenem stanju 3,7 µm ter
pri staranem jeklu maraging 5,2 µm. Povpre~na gostota
D. KLOB^AR ET AL.: ANALIZA TOPLOTNIH RAZPOK NA ORODJIH ZA TLA^NO LITJE ALUMINIJA
206 Materiali in tehnologije / Materials and technology 42 (2008) 5, 203–210
Slika 5: Potek trdote od povr{ine vogala proti centru epruvete po
preizkusu utrujanja
Figure 5: Hardness profile from the edge to the centre of the speci-
men after the immersion test
Slika 4: Razpoke na vzdol`nem prerezu kvadrastega vzorca za preizku{anje temperaturne utrujenosti: a) priprava vzorca, b) jeklo H13, kaljeno
na zraku, c) H13, TO v vakuumu, d) jeklo maraging v navarjenem stanju in e) navarjeno jeklo maraging starano 3 h na temperaturi 480 °C
Figure 4: Edge cracks at along immersion test specimen. a) Sample preparation, b) H13 tool steel air quenched, c) H13 tool steel vacuum
quenched, d) as-welded maraging steel, and e) maraging steel welded and aged for 3 h at 480 °C
Slika 6: Velikost razpok a) in {tevilo razpok b) na dol`ini milimetra
robu
Figure 6: Edge crack a) length and b) density
D. KLOB^AR ET AL.: ANALIZA TOPLOTNIH RAZPOK NA ORODJIH ZA TLA^NO LITJE ALUMINIJA
Materiali in tehnologije / Materials and technology 42 (2008) 5, 203–210 207
Slika 7: Kemi~na sestava jekla v okolici oksidnega klina, EDS (jeklo H13)
Figure 7: Chemical composition of steel in the area of the oxide wedge, EDS (H13 tool steel)
razpok na dol`ino robu je bila pri jeklu H13, ki je bilo
kaljeno v vakuumu, 13,2 mm–1, pri jeklu maraging v
navarjenem stanju 36,5 mm–1, pri jeklu maraging v
staranem stanju pa 33,6 mm–1.
3.6 Kemi~na mikroanaliza jekla v okolici razpok
Z energijsko disperzijskim spektrometrom rentgen-
skih `arkov smo analizirali kemi~no sestavo materiala v
okolici razpok. Zna~ilna kemi~na sestava jekla H13 ob
razpoki po 20 000 ciklih PP je prikazana na sliki 7. Na
temenu razpoke (slika 7 b) je pove~ana koncentracija Al
in Si, kar ka`e, da je teme razpoke zapolnjeno z alu-
minijevo zlitino, analiza pa je pokazala tudi sledi ogljika
in kisika. V oksidni plasti ob razpoki je pove~ana
koncentracija kroma. Podobno je tudi v konici razpok
(slika 7 c). Kisik penetrira po razpoki do same konice in
povzro~a oksidacijo. V konici razpoke nismo zaznali
sledi aluminija iz taline zaradi ozke razpoke in ve~je
oddaljenosti od vrha.
Kemi~na sestava okolice razpoke iz jekla maraging,
staranega 3 h pri temperaturi 480 °C ter po 20 000 ciklih
PP je prikazana na sliki 8. Razpoka je imela na povr{ini
plast oksidov `eleza, ki je nastala zaradi difuzije `eleza
iz podpovr{inske plasti h kisiku. Podpovr{inka plast je
bila obogatena z nikljem in kobaltom, ki difundirata v
jeklo iz oksidirane plasti `eleza. V kratkih razpokah so
bile sledi silicija in aluminija.
4 DISKUSIJA
4.1 Toplotne razpoke
Nastajanje temperaturnih razpok lahko razdelimo v
tri korake; (1) nukleacija, (2) rast razpok in (3) pora-
jajo~a rast razpok do poru{itve oz. lu{~enja povr{ine 5.
Tvorba razpok je tesno povezana z lokalno plasti~no
deformacijo na povr{ini jekla, kar je zna~ilno za
malocikli~no utrujanje. Rast razpok pospe{i oksidacija
razpokane povr{ine z ve~anjem volumna te plasti,
D. KLOB^AR ET AL.: ANALIZA TOPLOTNIH RAZPOK NA ORODJIH ZA TLA^NO LITJE ALUMINIJA
208 Materiali in tehnologije / Materials and technology 42 (2008) 5, 203–210
Slika 8: Kemi~na sestava jekla v okolici oksidnega klina, EDS (jeklo maraging)
Figure 8: Chemical composition of steel in the area oft the oxide wedge, EDS (maraging steel)
nadaljnjo rast pa polnjenje razpok s talino, oksidacija in
popu{~anje (prestaranje) jekla na povr{ini.
4.2 Tvorba in rast razpok
Med vro~o fazo cikla litja se na povr{ini orodja
pojavijo velike tla~ne napetosti (slika 2 b), ki jih pove~a
tlak polnjenja gravure. Te napetosti sicer zavirajo
nukleacijo in rast razpok, zato bi bile za`elene, vendar
povzro~ajo lokalno plasti~no deformacijo. Le-ta je
posebej izrazita na mestih, kjer so prisotni zarezni
u~inki, ostri prehodi ali velike spremembe mase orodja.
V hladni fazi cikla se na povr{ini orodja pojavijo velike
natezne napetosti, ki so posledica plasti~ne deformacije.
Le-te pri temperaturi obratovanja prese`ejo kriti~no
vrednost nosilnosti jekla. Pri cikli~nem ponavljanju
obremenitev nastane malocikli~no utrujanje, ki povzro~a
nastanek in rast razpok.
Rast razpok pospe{i oksidacija povr{ine jekla.
@elezovi atomi iz podpovr{inskega sloja difundirajo na
povr{ino h kisiku, prazna mesta pa zapolnijo legirni
elementi (sliki 7 in 8). Na povr{ini orodja in razpoke se
pojavi plast, sestavljena prete`no iz oksidov `eleza, tam
pa so tudi oksidi aluminija in silicija, ki sta rezultat
reakcije med kisikom in aluminijevo zlitino. Na meji
med oksidno plastjo in jeklom (H13) se na zunanji strani
nalagajo kromovi oksidi oz. oksidi kobalta in niklja pri
jeklu maraging. Slabe strani oksidne plasti so nizka
temperaturna razteznost, ve~ji specifi~ni volumen in
krhkost. Oksidi in aluminijeve zlitine v razpokah pove-
~ajo natezne napetosti v klinu razpoke (med hladno fazo
cikla), ki omogo~ajo rast razpok (slika 7 b) 5,19.
Natezne napetosti povzro~ajo lokalne razpoke v
oksidni plasti in v polnilu razpoke (krhkost in razlika v
temperaturni razteznosti) ter odpiranje razpok. Te v
nadaljevanju delujejo kot kanali za dovajanje alumi-
nijeve zlitine v razpoko in omogo~ajo penetracijo kisika
do vrha razpoke, kjer povzro~ajo oksidacijo.
4.3 Rast razpok zaradi popu{~anja (prestaranja) jekla
Orodja za tla~no litje aluminija, izdelana iz CrMoV
orodnih jekel za delo v vro~em, delujejo nad tempe-
raturami sekundarnega utrjevanja 14. Pri takih pogojih se
skoraj v vseh jeklih pojavi nagnjenost k popu{~anju
jekla. Opazimo lahko tri stopnje meh~anja: (1) hitro
meh~anje, kontrolirano z redukcijo dislokacij, (2) navi-
dezno stacionarno meh~anje, ki ga vodi razvoj karbidov
in (3) drasti~no zmanj{anje napetosti zaradi pojava
velikih razpok 14. Ne`eleno popu{~anje jekla zmanj{uje
napetost te~enja jekla in povzro~a hitro napredovanje
razpok do globine popu{~ene plasti. Popu{~anje je
posebej o~itno pri jeklu H13, ki je imelo zaradi kaljena
na zraku razoglji~eno povr{ino (slika 5). Slab{e mehan-
ske lastnosti te povr{ine so povzro~ile rast razpok, kar je
sprostilo napetosti v njihovi okolici ter s tem prepre~ilo
nastanek novih razpok. Rast razpok prepre~uje ve~ja
trdota (sliki 5 in 6 a), zato se je v jeklu H13, ki je bilo
TO v vakuumu, pojavila ve~ja gostota kraj{ih razpok.
V navaru iz jekla maraging je trdota padla zaradi
prestaranja med temperaturnim utrujanjem (slika 5).
Kraj{e razpoke so se pojavile v nestaranem varu, ker
pred preizku{anjem ni bil staran. Kljub temu pa toplotne
obdelave staranja pri sanaciji orodij ne smemo izpustiti,
saj s tem tvegamo pojav drugih po{kodb zaradi pre-
majhne trdote delovne povr{ine.
Gostota temperaturnih razpok je v korelaciji z nape-
tostjo te~enja jekla (slika 6 b), ki zna{a za jeklo H13
1230 MPa, za jeklo maraging pa 1688 MPa 24,25. V
materialih z vi{jo napetostjo te~enja se pojavi ve~ja
gostota kraj{ih razpok. Ker je trdota jekla v korelaciji z
napetostjo te~enja bi pri~akovali v jeklu maraging kraj{e
razpoke oz. bolj{o odpornost proti temperaturnemu
utrujanju, ~e bi prepre~ili prestaranje jekla maraging med
delom. Proti prestaranju so bolj odporna jekla maraging
z manj Ni (od 10 % do 14 %).
Tabela 1: Toplotna obdelava epruvet iz jekla H13
Table 1: Heat treatment of H13 test specimens.
TO 1 – vakuum – kakovostna TO TO 2 – brez za{~itne atmosfere – TOslabe kakovosti
KALJENJE 1005 °C/35 min KALJENJE 1000 °C/30 min
POPU[^ANJE 1 600 °C/150 min POPU[^ANJE 1 530 °C/150 min
POPU[^ANJE 2 570 °C/240 min
kaljenje = quenching, popu{~anje = tempering, vakuum = vacuum,
brez za{~itne atmosfere = without protective atmosphere
5 SKLEPI
V orodjih za tla~no litje ne`eleznih kovin je toplotno
utrujanje glavni krivec za po{kodbe. V vro~i fazi cikla se
na povr{ini orodja pojavijo velike tla~ne napetosti, ki jih
dodatno zvi{uje tlak polnjenja livne votline s talino. Na
mestih hitrih in ostrih prehodov se pojavi plasti~na
deformacija jekla. V hladni fazi cikla se na teh mestih
pojavijo velike natezne napetosti, ki povzro~ijo nukle-
acijo razpok in rast s pove~evanjem {tevila ciklov.
Oksidacija povr{ine orodja in povr{ine razpok ter
polnjenje razpok z aluminijevo talino pove~ajo natezne
napetosti v klinu med hladno fazo in omogo~ajo rast
razpok. V jeklih z ve~jo napetostjo te~enja in ve~jo
trdoto se pojavi ve~je {tevilo kraj{ih razpok, medtem ko
se pri materialih z manj{o napetostjo te~enja in manj{o
trdoto pojavi manj{e {tevilo dalj{ih razpok. Trajnost
orodij za tla~no litje aluminijevih zlitin lahko podalj-
{amo z uporabo trdnej{ih, trdih in `ilavih materialov, ki
so bolj odporni proti popu{~anju (staranje) med obrato-
vanjem.
6 LITERATURA
1 A. Srivastava, V. Joshi, R. Shivpuri, Computer modeling and
prediction of thermal fatigue cracking in die-casting tooling, Wear
256 (2004), 38–43
2 V. Joshi, K. Kulkarni, R. Shivpuri, R. S. Bhattacharya, S. J. Dikshit,
D. Bhat, Dissolution and soldering behavior of nitrided hot working
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Materiali in tehnologije / Materials and technology 42 (2008) 5, 203–210 209
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relation to interface mass loss and tribology in die casting dies. Wear
256 (2004), 1232–1235
4 A. Persson, S. Hogmark, J. Bergstrom, Simulation and evaluation of
thermal fatigue cracking of hot work tool steels. International Journal
of Fatigue 26 (2004), 1095–1107
5 A. Persson, S. Hogmark, J. Bergstrom, Failure modes in field-tested
brass die casting dies. Journal of Materials Processing Technology
148 (2004), 108–118
6 A. Persson, S. Hogmark, J. Bergstrom, Thermal fatigue cracking of
surface engineered hot work tool steels. Surface and Coatings
Technology 191 (2005), 216–227
7 Young, W.: Why Die Casting Dies Fail, Paper No. G-T79-092,
(1979), 1-7. St.Louis, Missouri, NADCA. 10th SDCE International
die casting exposition & congress
8 Y. Wang, A study of PVD coatings and die materials for extended
die-casting die life. Surf. and Coat. Techn. 94–95 (1997), 60–63
9 K. Kulkarni in sod., Thermal cracking behavior of multi-layer
LAFAD coatings on nitrided die steels in liquid aluminum
processing. Surf. and Coat. Techn. 149 (2002), 171–178
10 P. Panjan in sod., Improvement of die-casting tools with duplex
treatment. Surface and Coatings Technology 180–181 (2004),
561–565
11 Marlok, Longer die life Better quality, www.metsopowdermet.com
(2005). Metso Powdermet.
12 Fuchs, K. D.: Hot-work tool steels with improved properties for die
casting applications, Rosso, M., Actis Grande, M, in Ugues, D. II,
(2006). 17–26. Torino, Proceedings of 7th International Tooling
Conference
13 Ahmer, Z. in sod., Cyclic behaviour simulation of
X38CrMoV5-47HRC (AISI H11) tempered martensitic hot work
tool steel, Rosso, M., Actis Grande, M, in Ugues, D. II, (2006).
513–520. Torino, Proceedings of 7th International Tooling
Conference
14 Bergström, J., F. Rézai-Aria.: High temperature fatigue of tool steels,
Rosso, M., Actis Grande, M, in Ugues, D. II, (2006), 545–554.
Torino, Proceedings of 7th International Tooling Conference
15 Michaud, P. in sod., Influence of chemical composition on the
precipitation of secondary carbides in modified AISI H11 hot-work
tool steels, Rosso, M., Actis Grande, M, Ugues, D. (2006), 733–740.
Torino, Proceedings of 7th International Tooling Conference
16 Daffos, C., P. Lamesle, F. Rézai-Aria.: Fatigue-oxidation interaction
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thermomechanical loadings, Rosso, M., Actis Grande, M, in Ugues,
D. II, (2006), 471–478. Torino, Proceedings of 7th International
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effects on thermal fatigue damage accumulation and life of
die-casting steel X38CrMoV5 (AISI H11), Rosso, M., Actis Grande,
M, in Ugues, D. II, (2006), 461–470. Torino, Proceedings of 7th
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neering Failure Analysis 8 (2001), 355–359
19 F. Kosel, L. Kosec, Heat checking of hot work tools. Mechanical
Engineering Journal 29, (1983), E1-E8
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processing of Die Materials to Extend Die Life, Die Casting
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21 Klob~ar, D., J. Tusek, B. Taljat, G. Scavino.: Influence of thermal
fatique on materials for die-casting tooling, Rosso, M., Actis Grande,
M., Ugues, D. (2-5-2006). 479–485. Torino, Proceedings of 7th
International Tooling Conference
22 D. Klob~ar, J. Tu{ek, B. Taljat, Thermal fatigue of materials for
die-casting tooling. MSEA (2007)
23 Abaqus 6.4, (2003), Hibbit, Karlsson & Sorensen, Inc.
24 Materialne lastnosti jekla H13, www.sz-metal.si, (2003), 6-9-2003
25 QRO 90 Supreme, Hot work tool steel, Tool steel facts,
www.uddeholm.com, Uddeholm. 15.9.2005
D. KLOB^AR ET AL.: ANALIZA TOPLOTNIH RAZPOK NA ORODJIH ZA TLA^NO LITJE ALUMINIJA
210 Materiali in tehnologije / Materials and technology 42 (2008) 5, 203–210
M. PLETERSKI ET AL.: LASERSKO REPARATURNO VARJENJE TERMORAZPOK NA ORODJIH ...
LASERSKO REPARATURNO VARJENJE
TERMORAZPOK NA ORODJIH ZA TLA^NO LITJE
ALUMINIJA
LASER REPAIR WELDING OF THERMAL CRACKS ON
ALUMINIUM DIE CASTING DIES
Matej Pleterski1, Janez Tu{ek1, Ladislav Kosec3, Damjan Klob~ar1, Mitja Muhi~2,
Tadej Muhi~2
1Univerza v Ljubljani, Fakulteta za strojni{tvo, A{ker~eva 6, 1000 Ljubljana
2TKC, d. o. o., Trnovska 8, 1000 Ljubljana
3Univerza v Ljubljani, Naravoslovnotehni{ka fakulteta, A{ker~eva 12, 1000 Ljubljana
Prejem rokopisa – received: 2007-09-20; sprejem za objavo – accepted for publication: 2008-07-02
Sanacija po{kodovanih in izrabljenih orodij je ukrep, s katerim lahko v veliki meri pove~amo produktivnost in zmanj{amo
proizvodne stro{ke. Zato se v zadnjem ~asu vedno ve~ uporablja laserska tehnologija, katere prednost je predvsem v ozkem
lokalnem delovanju toplotne energije, majhni toplotno vplivani coni in zanemarljivih obrobnih zajedah. Orodja za tla~no litje so
med delovanjem izpostavljena kompleksnim termomehanskim obremenitvam in s tem visokim napetostim na povr{ini, ki
povzro~ijo utrujanje materiala in nastanek razpok. V ~lanku je predstavljena tehnologija sanacije razpok (`lebljenje, varjenje) z
bliskovnim Nd:YAG-laserjem, vklju~no z analizo mikrotrdot ob sami razpoki, `lebu in varu. Rezultati ka`ejo na to, da je z
laserjem mogo~e relativno hitro in enostavno odstraniti utrujeno obmo~je ob razpoki, ga navariti ter tako orodje usposobiti za
obratovanje.
Klju~ne besede: tla~no litje, orodja, razpoke, poprava, lasersko navarjanje, aluminij
Repair welding of damaged and worn-out tools is a measure, which can increase productivity and lower production costs to a
great extent. For this purpose, laser technology has recently been used with major benefits such as localized heating effect,
narrow heat effected zone and negligible undercuts. In the die-casting process, dies are submitted to a complex thermo
mechanical stressing and high stresses on the surfaces of dies are induced. That may lead to thermal fatigue cracking. In this
paper, the technology of thermal cracks repair (grooving, welding) with pulsed Nd:YAG laser is described. A micro-hardness
analysis of areas surrounding cracks, grooves and welds has been performed, also. The test results suggest that it is relatively
fast and easy to eliminate the fatigue area surrounding the cracks, and with proper welding to reestablish the tool operability.
Key words: die casting, dies, cracks, repairing, laser welding, aluminium
1 UVOD
Lasersko varjenje je ena izmed prvih industrijskih
aplikacij laserja. Ta alternativna tehnologija daje nove
mo`nosti obdelave, dodelave in sanacije, kjer je varjenje
z drugimi varilskimi metodami ote`eno ali celo onemo-
go~eno. V to vrsto laserske obdelave spada tudi lasersko
reparaturno zvarjanje in navarjanje, ki je relativno nova
tehnologija in se vedno bolj uveljavlja v livarski,
orodjarski ter drugi strojni industriji za sanacijo in
vzdr`evanje orodij. Orodja za tla~no litje so izdelana iz
kakovostnih jekel za delo v vro~em, ki so dobro odporna
proti toplotnim deformacijam, {okom ter termi~nim
razpokam na povr{ini. Ena izmed najpomembnej{ih
karakteristik orodja je trdota. Pri orodjih za tla~no litje
barvnih kovin mora biti le-ta okoli HRc = 45. Za
pove~anje trajnostne dobe jih lahko {e nitriramo in s tem
pove~amo njihovo povr{insko odpornost proti obrabi.
Orodja so med uporabo izpostavljena kompleksnim
termomehanskim obremenitvam in s tem visokim
napetostim na povr{ini, ki povzro~ijo utrujanje materiala
ter nastanek razpok na povr{ini, ne glede na vrsto
materiala, toplotno obdelavo in vrsto oplemenitenja
povr{ine orodja. Na sliki 1 je prikazana tipi~na termi~na
razpoka, ki se mo~no izra`a tudi na odlitku. Tak{ni
odlitki so nesprejemljivi, zato je treba orodja med
obratovalno dobo sanirati. Zato so se navadno
uporabljali procesi, kot so bru{enje, rezkanje in oblo~no
varjenje. Z laserjem je mogo~e razpoko iz`lebiti, pa tudi
Materiali in tehnologije / Materials and technology 42 (2008) 5, 211–214 211
UDK 621.791:669.715 ISSN 1580-2949
Izvirni znanstveni ~lanek/Original scientific article MTAEC9, 42(5)211(2008)
Slika 1: Odtisi temen razpok zaradi termi~nega razpokanja na ulitku
Figure 1: Imprints of cracks on a cast, as a result of thermal cracking
zavariti. Sun in sod. 1,2 so napravili {tudiji o trajnodina-
mi~ni in natezni trdnosti ter `ilavosti po pretaljevanju
razpok, ki so nastale pri trajni dinami~ni obremenitvi pri
povi{anih temperaturah. Treba je poudariti, da je {lo pri
obeh njihovih raziskavah zgolj za pretaljevanje razpok
brez dodajanja materiala. Vedani in sod. 3 so opravili
{tudijo s poudarkom na razvoju mikrostuktur in meta-
lur{kih problemih pri reparaturnem varjenju povr{insko
obdelanih orodnih jekel. Nekateri avtorji 4,5 so razisko-
vali lasersko reparaturno varjenje na razli~nih podro~jih.
Brown in sod. domnevajo, da je lasersko reparaturno
varjenje razpok na ladijskih plo~evinah bolj u~inkovito
kot varjenje z oblo~nimi postopki zaradi kraj{ih ~asov in
ni`jih stro{kov.
2 EKSPERIMENTALNI DEL
Eksperimenti so bili izvedeni na obrabljenem
stranskem jedru orodja za tla~no litje koluta za navijanje
varnostnega pasu v avtu, vzdol` katerega je potekala
razpoka. Orodje iz materiala 1.2343 (X38CrMoV5-1),
kaljeno na trdoto 45 ± 2 HRc (450 ± 30 HV) je prika-
zano na sliki 2. Orodje je bilo na za~etku odrezano na
rezu 1. Ta kos se je rabil za analizo trdote ob razpoki.
Razpoko na preostalem delu orodja se je nato po`lebilo z
bliskovnim Nd:YAG-laserjem mo~i 120 W. Po `leblje-
nju je bilo orodje odrezano na rezu 2. Drugi kos se je
rabil za analizo trdote ob `lebu. Preostanek orodja je bil
nato zavarjen z dodajnim materialom Uddeholm G3
premera 0,5 mm in odrezan na rezu 3. Uporabljeni
parametri za `lebljenje in varjenje so prikazani v tabeli
1. Pri iz`lebljenju z laserskim `arkom je treba izbrati
parametre z visoko gostoto energije v gori{~u laserskega
`arka. Material, tj. jeklo, je bilo treba upariti in ga
odstraniti iz razpoke. Za `lebljenje je zahtevana gostota
energije ve~ kot 1010 W/m2. To se dose`e z nastavitvijo
visoke mo~i na laserskem izvoru in izbiro le~e s kratko
gori{~no razdaljo. Parametri za varjenje morajo biti
prilagojeni predvsem velikosti iz`lebljenega utora in
premeru uporabljene `ice. V obeh primerih je bil kot
za{~itni plin uporabljen Ar ~istote 99,9996 %. Iz
odrezanih delov orodja so bili napravljeni makroobrusi,
na katerih so bile nato opravljene meritve mikrotrdot po
Vickersu pri obremenitvi 100 g oz. sili 0,981 N. Vsi
vzorci so bili jedkani z nitalom (4 %). Meritve so bile
opravljene na treh globinah orodja: (0,3, 0,6 in 1) mm.
Zaradi bolj{e preglednosti so v nadaljevanju predstav-
ljene zgolj meritve trdot na globini 1 mm.
Tabela 1: Prikaz nastavitvenih parametrov pri laserskem `lebljenju in
varjenju razpok
Table 1: Parameters for laser grooving and welding
Parameter @lebljenje Varjenje
Frekvenca las. bliskov y /Hz 6 7,4
^as trajanja las. bliska tp /ms 6,2 6,2
Povpre~na mo~ las. bliska Ppp /W 900 400
Povp. energija las. bliska Ep /J 15,3 11,9
Gori{~na razdalja le~e fl /mm 160 200
3 REZULTATI IN DISKUSIJA
Na sliki 3 je prikazan makroobrus razpoke. Z merit-
vami trdote, prikazanimi na sliki 4 a, smo ugotovili, da
je trdota okoli razpoke vi{ja, kar je razvidno tudi iz
diagrama na sliki 4 b, kjer je trdota ob razpoki v
povpre~ju za HV = 200 vi{ja od normalne trdote orodja.
To sicer ni zna~ilno za podro~ja ob termi~nih razpokah,
tako da lahko sklepamo, da je bilo to orodje pred na{o
analizo verjetno `e sanirano. Navadno trdota materiala
ob neposredni bli`ini termi~ne razpoke rahlo pade, kar je
posledica kemijske nehomogenosti jekla v obmo~ju ob
povr{ini razpoke. @elezovi atomi iz sloja pod povr{ino
razpoke difundirajo na povr{ino proti kisiku, na prazna
mesta pa difundirajo legirni elementi (Cr, Mo, V). Na
povr{ini razpoke se tako pojavi plast, ki je v veliki meri
sestavljena iz `elezovih oksidov in oksidov aluminija in
silicija, kar je posledica reakcije aluminijeve zlitine in
kisika. Na meji te oksidne plasti in osnovnega materiala
pa se kopi~ijo predvsem oksidi kroma. Slaba stran
oksidne plasti je nizka temperaturna razteznost, ve~ji
volumen in krhkost. Oksidi in aluminijeve zlitine v
razpokah pove~ajo natezne napetosti v klinu razpoke, kar
povzro~a njihovo rast 6. Razpoke v nadaljevanju delujejo
kot kanali za dovajanje aluminijeve zlitine v razpoko in
M. PLETERSKI ET AL.: LASERSKO REPARATURNO VARJENJE TERMORAZPOK NA ORODJIH ...
212 Materiali in tehnologije / Materials and technology 42 (2008) 5, 211–214
Slika 3: Makroskopska slika termi~ne razpoke
Figure 3: Macro section of a thermal crack
Slika 2: Stransko jedro z ozna~enimi podro~ji razreza
Figure 2: Test specimen with marked cuts
omogo~ajo penetracijo kisika do konice razpoke, kjer
povzro~i oksidacijo. Pri nadaljnjem obratovanju takega
orodja se razpoka {e hitreje {iri, postaja vedno globlja in
{ir{a. Vse skupaj vodi do nesprejemljivih ulitkov, orodje
pa lahko pri nadaljnjem obratovanju v takem stanju po~i
po celotni globini. Tako po{kodovana orodja se navadno
sanirajo (~e je sanacija sploh {e stro{kovno upravi~ena) z
rezkanjem razpoke in varjenjem po postopku TIG s
predgrevanjem, zajede pa se nato popravijo z laserjem.
Pri tak{ni sanaciji je treba veliko mehanske in toplotne
obdelave.
Na sliki 5 je prikazan makroobrus iz`lebljene raz-
poke, na katerem je lepo vidno toplotno vplivano
podro~je in tisto, ki je bilo med `lebljenjem pretaljeno in
ponovno zakaljeno (svetel rob tik ob `lebu). Ta, pri-
bli`no 0,2 mm {irok pas se izra`a tudi v diagramu
meritev trdot, ki je prikazan na sliki 6. Trdote so v tem
pasu tudi do HV = 300 ve~je od normalne trdote orodja
(izstopajo~a meritev, ki sega nad HV = 850, je posledica
kemi~ne nehomogenosti jekla v tem obmo~ju). Iz
diagrama je tudi razvidno, da trdota z oddaljenostjo od
zvarnega `lebu naglo pada in po 0,2 mm do 0,3 mm
dose`e normalno trdoto orodja. Razlog za tako visoke
trdote je predvsem v relativno veliki gostoti vnesene
energije, kar je pogoj za uspe{no `lebljenje obmo~ja ob
razpoki. Ker se orodja v primeru laserskega reparatur-
nega varjenja varijo v hladnem, se segreti material in
talina, ki se ne odstrani iz `leba, znova zakalita. Zaradi
velikega temperaturnega gradienta med tem obmo~jem,
segretim nad temperaturo tali{~a jekla, in osnovnim
materialom (sobna temperatura) se to obmo~je izredno
hitro ohladi in s tem mo~no zakali. Mikrostruktura, ki
nastane v takem primeru, je prikazana na SEM-posnetku
na sliki 7.
Na sliki 8 je prikazan makroobrus zavarjene razpoke.
Meritve trdot na varu so prikazane v diagramu slike 9.
Eden izmed ciljev pri reparaturnem varjenju je, da se
dose`e trdota zvara ~im bolj enaka trdoti osnovnega
materiala. Na to pa lahko vplivamo predvsem z izbiro
M. PLETERSKI ET AL.: LASERSKO REPARATURNO VARJENJE TERMORAZPOK NA ORODJIH ...
Materiali in tehnologije / Materials and technology 42 (2008) 5, 211–214 213
Slika 6: Profil trdote ob zvarnem `lebu
Slika 6: Hardness profile in the direction perpendicular to the groove
Slika 4: Merilna mesta (a) in profil trdote pravokotno na razpoko (b)
Figure 4: Location of measurements (a) and the hardness profile in
direction perpendicular to the crack (b)
Slika 7: Mikrostruktura jekla ob zvarnem `lebu (Sl. 5): A-osnovni
material; B-TVP; C-meja nataljevanja; D-obmo~je celi~ne mikro-
strukture, E-obmo~je aksialnih dendritov; SEM (SEI)
Figure 7: Image of microstructure at the groove: A-base material;
B-HAZ; C-solidification line; D-cellular dendrite area; E-axial
dendrite area; SEM (SEI)
Slika 5: Makroskopska slika iz`lebljene razpoke
Figure 5: Macro section of the groove
primernega dodajnega materiala in varilnimi parametri.
Zaradi majhne koli~ine taline in velikega gradienta
temperature uporabljamo pri laserskem reparaturnem
varjenju dodajne materiale, ki imajo zni`an dele` C (0,2
%) in zvi{ano vsebnost Cr (1,5 %) in Mo (3 %) glede na
osnovni material. Prav tako moramo biti pazljivi pri
nastavitvi parametrov, da vnesena energija ni prevelika,
vendar {e tolik{na, da se dodajni material pretali in v
celoti zlije z osnovnim. Na sliki 10 je prikazan
SEM-posnetek obmo~ja na robu zvara, ki je z me{ano
celi~no-dendritno mikrostrukturo podobno tistemu ob
zvarnem `lebu, zato so tudi trdote v tem prehodnem
podro~ju visoke. Mikrostruktura v sredini vara pa je
prikazana na sliki 11. Gre za martenzitno mikrostrukturo
z redkej{imi neizrazitimi dendriti in trdoto rahlo vi{jo od
osnovnega materiala.
4 SKLEPI
Rezultati raziskave ka`ejo, da je z lasersko tehnolo-
gijo mogo~e relativno hitro in enostavno sanirati razpo-
kano orodje. Z ustreznimi parametri laserskega `arka je
mogo~e razpoko iz`lebiti ter tako odstraniti utrujeni
material ob razpoki. ^eprav se ob tako pripravljenem
`lebu pojavi ozko pretaljeno podro~je, ki se mo~no
zakali, je mogo~e ob nastavitvi primernih laserskih
parametrov in izbiri pravega dodajnega materiala orodje
zvariti tako, da je trdota zvara blizu trdoti osnovnega
materiala.
5 LITERATURA
1 Y. Sun, S. Hanaki, M. Yamashita, H. Uchida, H. Tsujii: Vacuum 73
(2004), 655–660
2 Y. Sun, H. Sunada, N. Tsujii: ISIJ International, 41 (2001) 9,
1006–1009
3 M. Vedani: Journal of Materials Science (2004), 241–249
4 E. Capello, D. Colombo, B. Previtali: Journal of Materials Pro-
cessing Technology 164–165 (2005), 990–1000
5 P. M. Brown, G. Shannon, W. Deans, J. Berd: Weld World (1999),
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6 F. Kosel, L. Kosec: Mechanical Engineering Journal 29, (1983),
E1-E8
M. PLETERSKI ET AL.: LASERSKO REPARATURNO VARJENJE TERMORAZPOK NA ORODJIH ...
214 Materiali in tehnologije / Materials and technology 42 (2008) 5, 211–214
Slika 11: Mikrostruktura v varu; fina dendritna mikrostruktura; SEM
(SEI)
Figure 11: Microstructure in the weld; Fine dendritic microstructure;
SEM (SEI)
Slika 9: Profil trdote preko zvara
Figure 9: Hardness profile across the weld
Slika 10: Mikrostruktura jekla na robu zvara: A-osnovni material;
B-TVP; C-obmo~je celi~no-dendritne mikrostrukture; D-pasovi
toplotnega vpliva zaradi ve~varkovnega polnjenja `leba; SEM (SEI)
Figure 10: Microstructure at the weld boundary: A-base material;
B-HAZ; C- dendrite area; D-layers of heat effect of multi-pass groove
filling; SEM (SEI)
Slika 8: Makroskopska slika navarjenega zvarnega `leba
Figure 8: Macro section of the welded groove
U. ESME ET AL.: USE OF ARTIFICIAL NEURAL NETWORKS IN BALL BURNISHING PROCESS ...
USE OF ARTIFICIAL NEURAL NETWORKS IN BALL
BURNISHING PROCESS FOR THE PREDICTION OF
SURFACE ROUGHNESS OF AA 7075 ALUMINUM
ALLOY
UPORABA UMETNIH NEVRONSKIH MRE@ ZA NAPOVED
HRAPAVOSTI POVR[INE PRI KROGELNEM GLAJENJU
ALUMINIJEVE ZLITINE AA 7075
Ugur Esme1, Aysun Sagbas2, Funda Kahraman3, M. Kemal Kulekci4
Mersin University, Tarsus Technical Education Faculty, Department of Mechanical Education, 33400, Tarsus-Mersin/Turkey
uesme2003 hotmail.com
Prejem rokopisa – received: 2008-06-25; sprejem za objavo – accepted for publication: 2008-07-09
Burnishing is a plastic deformation process, and it has become more popular as a finishing process. Thus, it is especially crucial
to select the burnishing parameters to reduce the surface roughness. In the present study, a surface roughness prediction model
using artificial neural network (ANN) is developed to investigate the effects of burnishing conditions during machining of AA
7075 aluminum material. The ANN model of surface roughness parameters (Ra) is developed considering the conditions as
burnishing force, number of tool passes, feed rate and burnishing speed. The experimental results were trained in an ANN
program and the results were compared with experimental values. It is observed that the experimental results coincided with
ANN results.
Keywords: Ball burnishing, surface roughness, modeling, artificial neural network
Glajenje je proces plasti~ne deformacije in je postalo zelo raz{irjeno kot kon~na obdelava. Za zmanj{anje hrapavosti povr{ine je
zelo je pomembna izbira parametrov glajenja.V tej raziskavi je bil z uporabo nevronske mre`e (ANN) razvit model glajenja pri
obdelavi aluminijeve zlitine AA 7075. Model parametra hrapavosti povr{ine (Ra) je bil razvit z upo{tevanjem pogojev: polirna
sila, {tevilo prehodov orodja, hitrost podajanja in hitrost poliranja. Eksperimentalni podatki so uporabljeni za ANN-program,
rezultati modela pa primerjani z eksperimentalnimi. Rezultati ANN se dobro ujemajo z eksperimentalnimi.
Klju~ne besede: krogelno glajenje, hrapavost povr{ine, modeliranje, umetna nevronska mre`a
1 INTRODUCTION
The surface quality is an important parameter to
evaluate the productivity of machine tools as well as
machined components. Hence, achieving the desired
surface quality is of great importance for the functional
behavior of mechanical parts 1. Surface roughness is
used as the critical quality indicator for the machined
surfaces and since, it affects several properties such as
wear resistance, fatigue strength, coefficient of friction,
lubrication, wear rate and corrosion resistance of the
machined parts 2. In today’s manufacturing industry,
special attention is given to dimensional accuracy and
surface finish. Thus, measuring and characterizing the
surface finish can be considered as a predictor for the
machining performance.
Burnishing is considered as a cold-working finishing
process differing from other cold-working surface
treatment processes such as shot peening and sand
blasting, etc. in that it produces a good surface finish and
also induces residual compressive stresses at the metallic
surface layers 3,4. Accordingly, the burnishing is
distinguished from chip-forming finishing processes
such as grinding, honing, lapping and super-finishing
which induce residual tensile stresses at the machined
surface layers 5. Also, burnishing is economically
desirable, because it is a simple and cheap process,
requiring less time and skill to obtain a high-quality
surface finish 6. The burnishing process can be achieved
by applying a highly polished and hard ball or roller onto
a metallic surface under pressure. As indicated in Figure
1, pressure causes the peaks of the metallic surface to
spread out permanently and fill the valleys 4, when the
applied burnishing pressure exceeds the yield strength of
the metallic material.
The surface of the metallic material will be smoothed
out and because of the plastic deformation the surface is
Materiali in tehnologije / Materials and technology 42 (2008) 5, 215–219 215
UDK 669.715:620.17.112 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 42(5)215(2008)
Figure 1: Schematic representation of ball burnishing process 7
Slika 1: Shematska predstavitev procesa krogelnega poliranja 7
work hardened and the material is left with a residual
stress distribution compressive on the surface 4. The
changes in surface characteristics due to burnishing will
cause improvements in surface hardness, wear resi-
stance, fatigue resistance, yield and tensile strength and
corrosion resistance, as claimed by many authors 8–11.
The aim of the present work was to investigate the
effect of burnishing parameters such as burnishing force
(F/N), number of tool passes (N), feed rate (f/(mm/min))
and burnishing speed (v/(r/min)) on the surface
roughness (Ra/µm) of AA 7075 aluminum with the use of
ANN.
2 MATERIAL AND EXPERIMENTAL
PROCEDURE
2.1 Material
In this study, high strength precipitation hardening
7XXX series wrought aluminum alloy AA 7075 was
used. The strength and good mechanical properties make
the AA 7075 aluminum alloy appropriatefor use in
aerospace industry. The chemical composition and
mechanical properties of the workpiece material is given
in Table 1.
Table 1: Chemical and mechanical properties of workpiece material
Tabela 1: Sestava in mehanske zna~ilnosti obdelovanca
Chemical
composi-
tion w/%
Al Cu Mg Cr Zn
90.0 1.60 2.50 0.23 5.60
Mechani-
cal pro-
perties
Tensile
strength
(MPa)
Yield
strength
(MPa)
Shear
strength
(MPa)
Fatique
strength
(MPa)
Hard-
ness
(HB 500)
220 95 150 160 60
The three part workpiece material shown in Figure 2,
was prepared with the dimensions of 30 mm diameter
and 60 mm in length with each segment with 20 mm in
length.
2.2 Machines and Equipments
A 18 mm diameter steel ball was used for burnishing.
The detailed drawing is shown in Figure 3. When the
ball or roller is pressed against the surface of the metallic
specimen, a pre-calibrated spring was compressed used
mainly to reduce the possible sticking of the tool onto
the surface.
The experiments were performed on a FANUC
GT-250B CNC machining center. The burnishing tool
was mounted on the CNC turret as shown in Figure 4.
Dry turning and burnishing were used in all the
experimental work and alcohol was used to clean the
specimens before burnishing. The cleaning of the ball
was carried out continuously in order to prevent any hard
particles from entering the contact surface between the
tool and the specimen, such hard particles usually
leaving deep scratches that may damage the burnished
surface of the specimen. The Phynix TR-100 model
surface roughness tester was used to measure the surface
roughness of the burnished samples. Cut off length was
chosen as 0.3 for each roughness measurement. Six
measurements of surface roughness were taken from the
samples and average of the roughness was used in
modeling.
U. ESME ET AL.: USE OF ARTIFICIAL NEURAL NETWORKS IN BALL BURNISHING PROCESS ...
216 Materiali in tehnologije / Materials and technology 42 (2008) 5, 215–219
Figure 4: Ball burnishing experimental set up
Slika 4: Eksperimentalna priprava za krogelno glajenje
Figure 2: Dimensions of workpiece material
Slika 2: Mere obdelovanca
Figure 3: Detailed drawing of the ball burnishing tool: (1) casing; (2)
adapter cover; (3) spring
Slika 3: Na~rt orodja za krogelno glajenja. (1) ohi{je, (2) prilago-
ditveni pokrov, (3) vzmet
3 MODELING WITH ARTIFICIAL NEURAL
NETWORK (ANN)
Computers are an integral part of day to day activities
in engineering design and engineers have utilized
various applications to assist them improve their design
12. ANN mimics some basic aspects of the brain
functions 13–15. It is based on the neural structure of the
human brain, which processes information by means of
interaction between many neurons 13,16. In the past few
years there has been a constant increase in interest of
neural network modeling in different fields of materials
science. The basic unit in the ANN is the neuron. The
neurons are connected to each other with weight factor.
A network is usually trained using a large number of
input with corresponding output data 17.
The ANN architecture used modeling of surface
roughness is illustrated in Figure 5. It consists of many
simple processing neurons organized in a sequence of
layers: input, intermediate (hidden) and output layers.
The simulation problem consists of finding a satisfactory
relationship between a set of neurons representing the
input data and associated known output. The selection of
the input parameters is a very important aspect of neural
network modeling 17. All relevant input parameters must
be represented as the input data of the neural network. In
this study burnishing force, number of passes, feed and
burnishing speed were used as inputs while surface
roughness was used as an output.
The ANN model used is 4 : 5 : 5 : 1 multilayer
architecture as shown in Figure 5. Yj (j = 1, 2,..,5) and Yi
(i = 1, 2,..,5) are the output of the hidden neurons.
3.1 The Training of the Network
Generally, there are three different learning
strategies. First, the trainer may tell the network what it
should learn (Supervised Learning), second, the trainer
may indicate whether or not the output is correct without
telling what the network should learn (Reinforcement
Learning) and finally, the network learns without any
intervention of the trainer (Unsupervised Learning). The
learning set consists of the inputs and the outputs used in
training the network. The required outputs take place in
this set in the case of supervised learning, while in other
cases, they are not found in it 17,18. In the present study,
the supervised learning approach was used. The com-
puter program has been developed under MATLAB 19
and as given in Table 2, a database of 30 experimental
results was used to train the ANN model.
Table 2: Experimental results and training set of ANN modeling
Tabela 2: Eksperimentalni rezultati in u~ni podatki za ANN-mode-
liranje
Exp.no
Burnishing
force
F/(9,86 N)
Number of
passes
N
Feed rate
f/(mm/min)
Burnishing
speed
v/(r/min)
Measured
surface
roughness
Ra/µm
1 9 2 0.62 200 0.30
2 10 3 0.80 400 0.37
3 11 2 0.60 500 0.37
4 12 3 0.45 800 0.47
5 13 2 0.45 1000 0.44
6 14 4 0.45 600 0.65
7 15 4 0.45 600 0.71
8 16 2 0.27 200 0.60
9 17 3 0.62 600 0.69
10 18 4 0.45 600 0.89
11 19 3 0.27 400 0.85
12 20 2 0.27 500 0.78
13 21 3 0.45 600 0.91
14 22 4 0.27 1000 1.12
15 23 2 0.62 700 0.75
16 24 3 0.45 600 1.06
17 25 2 0.27 200 1.02
18 9 4 0.27 200 0.38
19 10 2 0.62 300 0.33
20 12 4 0.45 400 0.54
21 16 3 0.80 500 0.63
22 13 3 0.60 600 0.51
23 15 2 0.27 700 0.55
24 16 3 0.62 800 0.64
25 17 4 0.45 900 0.82
26 20 3 0.62 1000 0.81
27 14 2 0.45 400 0.49
28 16 4 0.80 600 0.76
29 11 3 0.27 800 0.42
30 10 2 0.45 800 0.33
3.2 Testing Stage
In order to understand whether an ANN is making
good predictions, test data that has never been presented
to the network are used and the results are checked at
this stage. The statistical methods of root mean square
error (RMSE), the coefficient of multiple determination
(R2) values have been used for making comparisons
17,20–23. These values are determined by the following
equations:
RMSE
n
a pj j
j
= −
⎛
⎝
⎜
⎞
⎠
⎟∑1
2
1 2/
(1)
( )
( )
R
a p
p
j j
j
j
j
2
2
2
1= −
−⎛
⎝
⎜
⎜
⎜⎜
⎞
⎠
⎟
⎟
⎟⎟
∑
∑
(2)
U. ESME ET AL.: USE OF ARTIFICIAL NEURAL NETWORKS IN BALL BURNISHING PROCESS ...
Materiali in tehnologije / Materials and technology 42 (2008) 5, 215–219 217
Figure 5: The constructed ANN model
Slika 5: Razviti ANN-model
where; p is the predicted value, a the actual value and n
the number of samples.
4 RESULTS AND DISCUSSION
The comparisons of experimental results with the
ANN predictions have been depicted in terms of
percentage error for validation set of experiments. From
Table 3 it is evident that for our set of data the neural
network predicts the surface roughness nearer to the
experimental values. In the prediction of surface
roughness values the average errors for ANN is found to
be as 3.30 %.
The average RMSE was found to be as 0.0048. The
value of the multiple coefficient of R2 between
experimental results and ANN prediction is obtained as
0.998. This value showed that ANN model fits well with
U. ESME ET AL.: USE OF ARTIFICIAL NEURAL NETWORKS IN BALL BURNISHING PROCESS ...
218 Materiali in tehnologije / Materials and technology 42 (2008) 5, 215–219
Table 3: Validation set used for ANN analysis
Tabela 3: Podatki za preverjanje ANN-analize
Exp.no
Burnishing
force
F/(9,81 N)
Number of
passes
N
Feed rate
f/(mm/min)
Burnishing
speed
v/(r/min)
Measured
surface
roughness
Ra/µm
ANN
Predicted
surface
roughness
Rap/µm
Error
%
RMSE R2
1 10 2 0.62 200 0.34 0.36 -6.75 0.0051 0.9960
2 10 3 0.80 600 0.36 0.38 -4.23 0.0034 0.9984
3 11 4 0.27 200 0.50 0.52 -4.42 0.0049 0.9982
4 12 3 0.45 400 0.47 0.45 5.22 0.0055 0.9970
5 13 3 0.45 1000 0.49 0.48 1.91 0.0021 0.9996
6 15 3 0.10 600 0.64 0.63 2.22 0.0032 0.9995
7 17 4 0.27 600 0.84 0.87 -3.46 0.0065 0.9989
8 18 4 0.27 800 0.89 0.91 -2.12 0.0042 0.9996
9 21 2 0.62 800 0.72 0.74 -3.25 0.0052 0.9990
10 22 3 0.45 600 0.96 0.96 0.50 0.0011 1.0000
11 23 2 0.27 300 0.92 0.91 1.36 0.0028 0.9998
12 24 4 0.62 200 1.09 1.04 4.59 0.0112 0.9977
13 25 2 0.80 1000 0.87 0.89 -2.29 0.0045 0.9995
14 10 3 0.62 900 0.37 0.39 -4.64 0.0038 0.9980
15 11 4 0.45 800 0.47 0.48 -1.72 0.0018 0.9997
16 12 2 0.80 300 0.39 0.38 1.57 0.0014 0.9997
17 13 3 0.80 600 0.46 0.45 2.94 0.0030 0.9991
18 14 4 0.60 700 0.64 0.63 1.88 0.0027 0.9996
19 9 2 0.27 800 0.37 0.38 -2.38 0.0020 0.9995
20 25 4 0.27 1000 1.16 1.06 8.62 0.0224 0.9911
Average error: 3.30%
Average RMSE: 0.0048
Average R2: 0.998
Figure 7: Learning behavior of ANN model
Slika 7: U~no vedenje ANN-modela
Figure 6: Actual average surface roughness against ANN prediction
Slika 6: Dejanska hrapavost proti ANN-napovedi
the experimental results. Figure 6 illustrates the ANN
predictions against the experimental results.
The training of the neural network was performed
with an allowable error of 0.01 (sum of squared error
over the output neurons). The learning behavior of this
particular network is shown in Figure 7.
5 CONCLUSION
In this study, for the modeling of the effects of ball
burnishing parameters (burnishing force, number of
passes, feed rate and burnishing speed) on the surface
roughness of the AA 7075 aluminum alloy depending on
various processing parameters, an ANN-based approach
has been suggested and successfully implemented. As
Figure 6 indicates for each average surface roughness
value the predictions of the ANN are very close to the
experimental results. It may be concluded that the ANN
may be used as a good alternative for the analysis of the
effects of burnishing parameters on the average surface
roughness. In the field of surface roughness, ANNs are
good alternative to conventional empirical modeling.
The advantages of the ANN compared to classical
methods are speed, simplicity and capacity to learn from
the experimental results and also none need for a wider
experimental study. Because of this fact that, engi-
neering effort may be reduced in the areas where ANN
modeling is preferred.
In this study the focus was to predict the average
surface roughness in ball burnishing process. The results
from ANN model will allow to improve determination of
the average surface roughness value and help to deter-
mine in a short time the behavior of the experimental
results.
6 REFERENCES
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(2006), 156–159
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(2001), 1541–1552
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U. ESME ET AL.: USE OF ARTIFICIAL NEURAL NETWORKS IN BALL BURNISHING PROCESS ...
Materiali in tehnologije / Materials and technology 42 (2008) 5, 215–219 219

M. AMBRO@I^ ET AL.: MERJENJE OBRABNE OBSTOJNOSTI STRUKTURNE KERAMIKE Al2O3
MERJENJE OBRABNE OBSTOJNOSTI STRUKTURNE
KERAMIKE Al2O3
WEAR-RESISTANCE MEASUREMENT OF STRUCTURAL Al2O3
CERAMICS
1Milan Ambro`i~, 2Stojana Veskovi~ Bukudur, 1Toma` Kosma~,
1Kristoffer Krnel, 1Darko Eterovi~, 1Natalija Petkovi~ Habe, 1Irena Pribo{i~
1Odsek za in`enirsko keramiko (K6), Institut "Jo`ef Stefan", Jamova 39, Ljubljana, Slovenija
2Hidria AET, d. o. o., Poljubinj 89, Tolmin, Slovenija
Prejem rokopisa – received: 2008-02-15; sprejem za objavo – accepted for publication: 2008-03-31
Obstojnost keramike proti obrabi med vsakdanjo uporabo je poleg drugih mehanskih lastnosti – trdote, togosti, tla~ne in natezne
trdnosti ter `ilavosti – velikega pomena. Pri obrabni obstojnosti so pomembne mnoge lastnosti, kot so trdota, `ilavost, koeficient
trenja med sti~nimi povr{inami, morebitne kemi~ne reakcije zaradi mo~no povi{ane temperature in nastanek »triboplasti«, tako
da je zelo te`ko najti enoli~no povezavo med danimi veli~inami. Zato je treba obrabno obstojnost posebej kvantitativno
opredeliti in jo izmeriti v ponovljivih razmerah. V ~lanku opisujemo preprost in poceni na~in merjenja obrabne obstojnosti
strukturne korundne keramike z obi~ajno napravo za bru{enje in poliranje kerami~nih vzorcev.
Klju~ne besede: strukturna keramika, obrabna obstojnost, faktor obrabe
Wear resistance of ceramics in the every-day use is of great importance, besides other mechanical properties, e.g., hardness,
stiffness, compressive and tensile strength and toughness. Several characteristics contribute to wear resistance, such as hardness,
toughness, coefficient of friction between the surfaces in contact, possible chemical reactions due to highly increased
temperature, and the formation of tribofilm, making the search for the unique connection between these quantities very difficult.
Therefore, it is necessary to define the wear resistance quantitatively and measure it in repeatable circumstances. A simple and
cost-efficient method for wear-resistance measurement of the structural ceramic material alumina with the help of the ordinary
grinding/polishing machine for ceramic samples is presented.
Key words: structural ceramics, wear-resistance, wear factor
1 UVOD
Pomembni podro~ji uporabe strukturne keramike sta
in`enirska in proteti~na keramika. Strukturna keramika
mora imeti zadovoljive mehanske lastnosti: veliko
trdnost in trdoto ter ~im bolj{o `ilavost. Kjer pa je njena
povr{ina izpostavljena vsakodnevnemu trenju z drugimi
sti~nimi povr{inami, npr. v sklepnih kerami~nih prote-
zah, mora imeti tudi dobro obstojnost proti obrabi.
Zato je treba obrabno obstojnost kvantitativno
opredeliti in jo meriti v ~im bolj ponovljivih razmerah.
Osnova sodobnih meritev na komercialnih napravah je
preprosta: z dolo~eno silo je treba pritisniti povr{ino
merjenca ob referen~no povr{ino (referen~na povr{ina je
lahko iz enakega ali pa druga~nega materiala kot presku-
{ani vzorec), spraviti drsni povr{ini v relativno gibanje in
nazadnje izmeriti u~inke obrabnega preskusa, npr.
izmeriti zmanj{anje prostornine merjenca. Veliko obrab-
nih preskusov je bilo narejenih na dveh pomembnih
in`enirskih in biomedicinskih kerami~nih materialih:
aluminijevem in cirkonijevem oksidu (Al2O3 in ZrO2) in
njunih kompozitov.1-11 Aluminijev oksid (Al2O3) – ko-
rundna keramika – ima veliko trdoto in zato tudi veliko
obrabno obstojnost. Danes se veliko uporablja merilo, da
je material dobro obrabno obstojen v vsakdanji uporabi,
~e ima faktor obrabe po ena~bi (1) manj{i od vrednosti
10–6 mm3/J.8,10
Za vedenje trdih materialov pri trenju in obrabi se je
uveljavilo tudi ime "tribolo{ke lastnosti materialov". Ime
izhaja iz tega, ker pri stiku in relativnem gibanju raz-
meroma gladkih povr{in dveh trdih materialov nastane
med njima {e tretja plast, na kratko triboplast, in sicer iz
aglomeriranih odkru{enih del~kov z ene ali obeh kera-
mi~nih povr{in.6 Kemijsko in fazno sestavo triboplasti ter
njeno hrapavost, ki seveda mo~no vpliva na koeficient
trenja, lahko preu~imo z vrsti~nim in presevnim
elektronskim mikroskopom (SEM in TEM), rentgensko
spektroskopijo, mikroskopom na atomsko silo (AFM),
ramansko spektroskopijo, opti~nim mikroskopom –
profilometrom itd.1-10 Ve~ poskusov ka`e na to, da se po
za~etnem spreminjanju koeficient trenja med sti~nima
povr{inama ustali pri neki vrednosti, kar ka`e na to, da
ga zares dolo~a nastala triboplast.6,7,11
Krell in Klaffke sta primerjala obrabo povr{in
vzorcev Al2O3, tetragonalnega ZrO2 (z molskim dele`em
Y2O3 3 %) in kompozita Al2O3/TiC (TiC je titanov
karbid) pri drgnenju ob krogle iz Al2O3 pri razli~nih
pogojih (npr. razli~ni vla`nosti zraka).3 Pri tem je imela
keramika Al2O3 razli~ne povpre~ne velikosti zrn, od 0,4
µm do 3 µm, odvisno od majhne koli~ine dodatkov, npr.
MgO, ki imajo bistven vpliv na potek sintranja. Za Al2O3
sta ugotovila, da se z zmanj{anjem kristalnih zrn
zmanj{a faktor obrabe. Nadalje, vla`en zrak ali teko~a
voda med sti~nima povr{inama zmanj{ata koeficient
Materiali in tehnologije / Materials and technology 42 (2008) 5, 221–225 221
UDK 620.178.1:666.3/.7 ISSN 1580-2949
Strokovni ~lanek/Professional article MTAEC9, 42(5)221(2008)
trenja, kar v glavnem vodi do zmanj{anja faktorja
obrabe. Faktorja obrabe Al2O3 in kompozita Al2O3/TiC
sta podobna, obraba ZrO2 pa je ve~ja.
Liu in sodelavci so ugotovili, da dodatek 6 % (pro-
storninski dele`) diopsida, MgCa(SiO3)2 precej izbolj{a
tribolo{ke lastnosti keramike Al2O3.10 Med drugim so
potrdili splo{no znano dejstvo, da pri ve~jih normalnih
silah med drsnimi povr{inami nastane bolj izrazito
puljenje celih zrn, s ~imer se mo~no pove~a obraba.
Njihova najpomembnej{a ugotovitev pa je verjetno ta, da
se faktor obrabe k = 10–6 mm3/J za ~isti Al2O3 pri danih
pogojih preskusa zmanj{a na desetino vrednosti pri
12-odstotnem dele`u diopsida.
^eprav obstajajo komercialne naprave za merjenje
obrabne obstojnosti materialov na osnovi ena~be (1), pa
si lahko pomagamo tudi z improviziranimi preskusi na
obi~ajni napravi za bru{enje in poliranje kerami~nih
vzorcev. Na tak na~in smo na Odseku za in`enirsko
keramiko Instituta "Jo`ef Stefan" izmerili obrabno ob-
stojnost keramike Al2O3, pripravljene na razli~ne na~ine
in z razli~nimi primesmi.
2 OPREDELITEV OBRABNE OBSOJNOSTI
KERAMIKE
Obraba materiala pri drgnenju ob enak ali druga~en
material je tem ve~ja, ~im ve~ja je sila trenja med
povr{inama, ta sila pa nara{~a s pravokotno silo, ki ti{~i
podlagi eno ob drugo. Obraba nara{~a tudi s ~asom
drgnjenja oziroma z relativnim premikom med sti~nima
povr{inama. Zato je smiselno opredeliti faktor obrabe k
takole:
k
V
F s
= ∆
n
(1)
kjer je ∆V zmanj{anje volumna materiala oz. pre-
skusnega vzorca zaradi obrabe, Fn je pravokotna
(normalna) sila med sti~nima povr{inama, s pa je
razdalja, ki jo pri trenju naredi ena povr{ina relativno
glede na drugo. Primerna enota za faktor obrabe je npr.
mm3/J. Ena~ba (1) je v skladu s pri~akovanjem, da je
zmanj{anje volumna materiala sorazmerno z delom
zunanje sile, saj je za trganje medatomskih vezi v
materialu potrebna energija, {tevilo potrganih vezi pa je
sorazmerno tako z volumnom odnesenega materiala kot
z dovedeno notranjo energijo. Zavedati pa se moramo,
da ta sorazmernost velja le do dolo~ene mere, saj so
prisotni razli~ni mehanizmi odna{anja materiala s
povr{ine. Po drugi strani je sila trenja Ftr med sti~nima
pov{inama bolj vplivna od normalne sile Fn, saj je delo
sile trenja, ki je vzporedna z relativnim gibanjem obeh
povr{in, razli~no od ni~, medtem ko je delo normalne
sile ni~. Sila trenja je sicer sorazmerna z normalno silo,
vendar pa je odvisna tudi od koeficienta trenja ktr: Ftr =
ktrFn. Tako je lahko pri isti normalni sili Fn in pri sicer
enaki keramiki hitrost obrabe zelo razli~na, ~e imamo
razli~ne koeficiente trenja med sti~nima povr{inama,
npr. zaradi oblivanja povr{in z vodo. Vendar, ~e je pri
neki aplikaciji koeficient trenja vedno enak, je normalna
sila sorazmerna s silo trenja, tako da je ena~ba (1)
smiselna.
3 EKSPERIMENTALNO DELO
Zaradi sodelovanja z industrijo (AET – Hidria,
Tolmin) nas je zanimala obrabna obstojnost aluminije-
vega oksida Al2O3. Obrabne preskuse lahko naredimo
kar na napravi za poliranje kerami~nih vzorcev, ~e
poskrbimo za ponovljive pogoje presku{anja. Postopek
obrabnih preskusov na okroglih plo{~icah (tabletkah) z
uporabo diamantne paste z delci povpre~ne velikosti 45
M. AMBRO@I^ ET AL.: MERJENJE OBRABNE OBSTOJNOSTI STRUKTURNE KERAMIKE Al2O3
222 Materiali in tehnologije / Materials and technology 42 (2008) 5, 221–225
Slika 1: a) Struersova polirna naprava z nastavkom Pedepin 2; b) Pove~an izrez naprave z nosilcem vzorcev in vpetim preskusnim vzorcem
Figure 1: a) Struers polishing machine with the part Pedepin 2; b) Enlarged section of the machine with the sample holder and attached testing
sample
µm je bil podrobneje opisan v reviji Vakuumist.12 Tu
opi{emo obrabne preskuse, kjer smo vzorce brusili na
napravi Pedepin 2 danskega podjetja Struers (slika 1a).
Iz podjetja AET smo prejeli vzorce v obliki votlih valjev
z merami v tabeli 1. Vzorci so gredi vodnih ~rpalk iz
serijske proizvodnje. Od celotne dol`ine vzorcev smo za
posamezni preskus odrezali kose, dolge okrog 20 mm
(slika 1b). Keramik je bilo {est vrst, glede na vrsto
uporabljenih kerami~nih prahov in dodanih primesi. Na
primer oznaka AET (tip 6) pomeni keramiko iz me{anice
ve~ vrst za~etnih prahov, ki jih v AET navadno upo-
rabljajo za serijsko proizvodnjo bele korundne keramike,
medtem ko ima rjava keramika (tip 5) prime{an man-
ganov in titanov oksid. Zna~ilna temperatura sintranja
teh vzorcev je bila 1640 °C, le za rjavo keramiko je bila
ta temperatura 1300 °C. Gostote keramik so bile izmer-
jene z Arhimedovo metodo, ~eprav bi jih lahko preprosto
izra~unali iz mase in dimenzij vzorcev.
@e pri `aganju preskusnih dvocentimetrskih kosov z
diamantno `ago pri ponovljivih pogojih (enaka obte`ba,
temeljito o~i{~enje lista `age pred `aganjem itd.) smo
dobili prvo dokaj zanesljivo informacijo o obrabni
obstojnosti razli~nih keramik. ^asi `aganja, ki so se
gibali okrog vrednosti 5 min ali ve~, so se namre~ za
razli~ne skupine zelo razlikovali, medtem ko so si bili za
vzorce iz iste keramike med seboj podobni. Dalj{i ~as
`aganja nakazuje ve~jo obrabno obstojnost keramike
(tabela 2). Sam obrabni preskus na napravi Pedepin je
potekal takole: Vzorec smo pritrdili na natan~no ozna-
~eno mesto na nosilcu in ga najprej na kratko zbrusili z
bolj grobo brusno plo{~ico z diamantnimi zrni, tako da
smo poskrbeli za planparalelnost obdelovalne povr{ine z
drugo brusno plo{~ico v nadaljevanju preskusa. Vzorec
smo stehtali, nato pa ga obrusili s finej{o plo{~ico trikrat
po 1 min (za preskus ponovljivosti). Po vsaki minuti
bru{enja s finej{o plo{~ico smo vzorec ponovno stehtali
in izra~unali izgubo mase/volumna. Pred vsakim tehta-
njem smo vzorec o~istili v ultrazvo~ni kopeli z aceto-
nom. Zelo pomembno za ponovljivost rezultatov je bilo
tudi ~i{~enje brusnih plo{~ic pred vsakim preskusom.
4 REZULTATI IN DISKUSIJA
Ker smo preskuse izvajali pri nespremenljivih po-
gojih (stalna normalna sila Fn ≈ 50 N, frekvenca vrtenja
brusne plo{~ice 917 min–1, frekvenca vrtenja nosilca
vzorcev 300 min–1 itd.), nismo vsaki~ ra~unali faktorja
obrabe po ena~bi (1), temve~ nam je zadostovala
primerjava zmanj{anja njihove prostornine. Izidi meritev
so prikazani v tabeli 2. Naprava Pedepin nima vgraje-
nega merilnika normalne sile, zato je zgoraj omenjena
vrednost 50 N le groba ocena, ki smo jo dobili z uporabo
tehtnice. Na osnovi podatkov iz tabele (<V> ≈ 100 mm3),
ocenjene sile (Fn ≈ 50 N) in relativne drsne poti (s ≈ 65
m, izra~unali smo jo na osnovi frekvenc vrtenja ({tevila
obratov na minuto) polirne plo{~ice in nosilca vzorca,
ekscentri~nosti osi obeh vrtenj in ~asa preskusa) smo
ocenili faktor obrabe: k ≈ 0,03 mm3/J. Ta vrednost je
zelo velika v primerjavi s tistimi, navedenimi pri obi~aj-
nih preskusnih pogojih v literaturi, vendar pa je bru{enje
z brusno plo{~ico veliko bolj agresivno kot npr. med-
sebojno drgnjenje gladkih kerami~nih povr{in. @aganih
povr{in vzorcev nismo spolirali pred preskusom z
bru{enjem in pri~akovati je bilo, da bi bila obraba na
za~etku spoliranih povr{in veliko manj{a. Dodatni
preskusi so pokazali, da zmanj{anje volumna vzorcev ni
premo sorazmerno s ~asom, temve~ se hitrost obrabe
zaradi nastanka triboplasti hitro zmanj{uje. Za primer-
javo: faktor obrabe za laboratorijske Al2O3-tabletke, na
katerih smo naredili predhodne obrabne preskuse z
drgnjenjem ob fino polirno plo{~o, na katero smo poli-
vali redko vodno suspenzijo diamantnih delcev pov-
pre~nega premera 45 µm,12 je bil 6 · 10–4 mm3/J, vrstni
red razli~nih keramik po obrabni obstojnosti pa je bil
podoben kot pri preskusu z brusno plo{~ico. Ena
najpomembnej{ih veli~in pri obrabi materiala, koeficient
trenja, namre~ zelo variira pri razli~nih preskusnih
pogojih. Kljub tem razlikam pa je preskus z diamantno
brusno plo{~ico poceni, preprost in hiter, daje osnovno
informacijo o obrabni obstojnosti in tudi obdelovalnosti
razli~nih keramik, obe lastnosti (oziroma kompromis
med njima) pa sta pomembni za industrijsko tehnologijo.
^asa `aganja pri vzorcih skupin 3 in 6 sta samo
informativna, saj imajo ti vzorci bistveno druga~no
plo{~ino prereza (tabela 1), ~as `aganja pa je seveda
odvisen tudi od te veli~ine. Po drugi strani smo
ugotovili, da razli~ni prerezi vzorcev skoraj ni~ ne
vplivajo na zmanj{anje volumna med preskusom; to smo
preverili s primerjavo zmanj{anja volumna za isto
keramiko (isti votli valj na njegovem koncu, ki ima
druga~en premer votline kot v sredini valja). Iz tabele 2
je razvidno, da imajo keramike skupin 2, 3 in 6 pribli`no
M. AMBRO@I^ ET AL.: MERJENJE OBRABNE OBSTOJNOSTI STRUKTURNE KERAMIKE Al2O3
Materiali in tehnologije / Materials and technology 42 (2008) 5, 221–225 223
Tabela 1: Vrsta in mere vzorcev: ρ = gostota, M = masa, Dzun in Dnot = zunanji in notranji premer, L = dol`ina
Table 1: Type and dimensions of the samples: ρ = density, M = mass, Dzun and Dnot = outer and inner radius, respectively, L = length
Skupina Tip  /(g/cm3) M/g Dzun/mm Dnot/mm L/mm
1 MR52-23 3,76 20,66 9,95 2,30 77,07
2 MR32-23 3,76 20,78 9,95 2,30 77,13
3 KMS96 3,77 21,45 10,21 2,10 77,17
4 Alcoa 3,92 21,52 9,95 2,30 77,22
5 Rjava 3,92 24,08 9,95 2,77 89,13
6 AET 3,78 22,60 10,53 2,30 76,8
enako obrabno obstojnost, sledi keramika 1, manj
obstojni pa sta keramiki 4 in 5. ^asi `aganja potrjujejo
rezultate za zmanj{anje volumna med preskusi: ve~ji
volumenski obrabi ustreza dalj{i ~as `aganja vzorcev.
Tabela 2 Rezultati `aganja in obrabnega preskusa: <t> in δt sta
povpre~ni ~as in standardna deviacija ~asa `aganja vzorcev z
diamantno `ago (3 ali 4 meritve), <V> in δV sta povpre~no zmanj{anje
volumna vzorcev med enominutnim preskusom in ustrezna standardna
deviacija (3 meritve)
Table 2 Results of sawing and wear test: <t> and δt are average time
of sawing of samples with the diamond saw and its standard deviation,
respectively (3 or 4 measurements), <V> and δV are average volume
loss of samples after 1-minute test and the corresponding standard
deviation, respectively (3 measurements)
Skupina <t>/(min:s) t /s <V>/mm3 V/mm3
1 6:44 7 106 6
2 7:28 7 79 2
3 6:26 (inf) – 78 3
4 5:11 12 147 10
5 4:08 4 158 1
6 8:09 (inf) – 83 1
Povr{ine vzorcev po `aganju in preskusnem bru{enju
smo si ogledali z vrsti~nim elektronskim mikroskopom
(SEM). Slika 2 prikazuje zna~ilno povr{ino po `aganju
in bru{enju. Slike za vse vrste keramike so si podobne. S
slike lahko razberemo, da je glavni mehanizem pri
`aganju in bru{enju abrazija kristalnih zrn z diamantnimi
delci, puljenje celih kristalnih zrn s povr{ine pa je manj
pomembno.
Izmerili smo tudi trdoto razli~nih keramik z
Vickersovim odtisom z diamantno piramido pri dveh
obte`itvah, 2 kg in 5 kg. Pri tej metodi izra~unamo trdoto
z velikostjo odtisa (pod opti~nim mikroskopom vidimo
kvadrat, slika 3, sicer pa ima odtis obliko pravilne
4-strane piramide). Zna~ilne vrednosti trdote so okrog 15
GPa, med razli~nimi keramikami so sicer vidne razlike v
trdoti, a precej manj{e kot razlike v obrabnem preskusu
in tudi ni o~itne korelacije med trdoto in obrabno obstoj-
nostjo. Na slikah z opti~nega mikroskopa smo ocenili
poroznost keramik: le-ta je velikostnega reda 10 % in ni
opaziti bistvenih razlik za razli~ne keramike. Enako velja
za povpre~no velikost in velikostno porazdelitev por:
pore so v povpre~ju velike nekaj mikrometrov in so pri
rjavi keramiki nekaj manj{e kot drugod. Povpre~no
velikost kristalnih zrn smo izra~unali z analizo SEM
(Scanning Electron Microscope = vrsti~ni elektronski
mikroskop)-slik poliranih in termi~no jedkanih povr{in
kerami~nih vzorcev, njena vrednost pa je okrog 3 µm.
Razli~ne keramike torej nimajo o~itnih razlik v
mikrostrukturi in je verjetno za njihovo razli~no obrabno
obstojnost odlo~ilna morebitna razlika v strukturi snovi
med kristalnimi zrni.
5 SKLEP
Tudi z improviziranimi preprostimi preskusi na
standardni napravi za bru{enje in poliranje vzorcev se da
dokaj zanesljivo primerjati obrabno obstojnost razli~nih
kerami~nih materialov.
6 LITERATURA
1 Birkby I., Harrison P., Stevens R., The effect of surface trans-
formation on the wear behaviour of zirconia TZP ceramics, J. Eur.
Ceram. Soc. 5 (1989), 37–45
M. AMBRO@I^ ET AL.: MERJENJE OBRABNE OBSTOJNOSTI STRUKTURNE KERAMIKE Al2O3
224 Materiali in tehnologije / Materials and technology 42 (2008) 5, 221–225
Slika 3: Fotografija slike opti~nega mikroskopa povr{ine keramike
skupine 3 z odtisom Vickersove piramide. Na sliki je dobro opazna
tudi poroznost keramike.
Figure 3: Optical-microscope photograph of the surface of ceramics
from group 3 with the Vickers-pyramid indentation. The porosity of
the ceramic material is also evident.
Slika 2: SEM-fotografija `agane (a) in bru{ene (b) povr{ine keramike
skupine 3
Figure 2: SEM-photograph of the sawed (a) and ground (b) surface of
the ceramics from group 3
2 Fischer T., Anderson M. P., Jahanmir S., Influence of fracture tough-
ness on the wear resistance of yttria-doped zirconium oxide, J. Am.
Ceram. Soc. 72 (1989) 2, 252–257
3 Krell and D. Klaffke, Effects of grain size and humidity on fretting
wear in fine-grained alumina, Al2O3/TiC, and zirconia, J. Am.
Ceram. Soc. 79 (1996) 5, 1139–1146
4 Gee M. G., Jennett N. M., High resolution characterisation of tribo-
chemical films on alumina, Wear 193 (1996) 2, 133–145
5 Morita Y., Nakata K., Ikeuchi K., Wear properties of zirconia/
alumina combination for joint prostheses, Wear 254 (2003) 1–2,
147–153
6 Kalin M., Hockey B., Jahanmir S., Wear of hydroxiapatite sliding
against glass-infiltrated alumina, J. Mater. Res. 18 (2003) 1, 27–36
7 Basu B., Vleugels J., Van der Biest O., Microstructure-toughness-
wear relationship of tetragonal zirconia ceramics, J. Eur. Ceram. Soc.
24 (2004) 7, 2031–2040
8 Kerkwijk B., Garcya M., Van Zyl W. E. et al., Friction behaviour of
solid oxide lubricants as second phase in α–Al2O3 and stabilised
ZrO2 composites, Wear 256 (2004), 182–189
9 Novak S., Kalin M., Lukas P., Anne G., Vleugels J., Van der Biest
O., The effect of residual stresses in functionally graded alumina-
ZTA composites on their wear and friction behaviour, J. Eur. Ceram.
Soc. 27 (2007) 1, 151–156
10 Liu C., Zhang J., Sun J., Zhang X., Tribological properties of
pressureless sintered alumina matrix ceramic materials improved by
diopside, J. Eur. Ceram. Soc., in print
11 Singha Roy R., Guchhait H., Chanda A., Basu D., Mitra M. K., Im-
proved sliding wear-resistance of alumina with sub-micron grain
size: A comparison with coarser grained material, J. Eur. Ceram.
Soc., 27 (2007), 4737–4743
12 Ambro`i~ M., Obrabna obstojnost keramike, Vakuumist 27 (2007) 3,
10-15
M. AMBRO@I^ ET AL.: MERJENJE OBRABNE OBSTOJNOSTI STRUKTURNE KERAMIKE Al2O3
Materiali in tehnologije / Materials and technology 42 (2008) 5, 221–225 225

DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES
DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA –
DOCTOR'S, MASTER'S AND DIPLOMA DEGREES
DOKTORSKA DELA – DOCTOR'S DEGREES
Na Naravoslovnotehni{ki fakulteti Univerze v Ljubljani je dne 7.7.2008 pred
komisijo v sestavi: red. prof. dr. Radomir Turk kot predsednik in ~lani: red. prof.
dr. Ladislav Kosec, red. prof. dr. Anton Smolej in izr. prof. dr. Toma` Rodi~ Ale{
Nagode, univ. dipl. in`. materialov in metalurgije zagovarjal doktorsko disertacijo
z naslovom:
Analiza lezenja jekla 9Cr-1Mo-0,2V z upo{tevanjem napetostno odvisne
aktivacijske energije
An analysis of the creep behaviour of 9Cr-1Mo-0.2V steel based on the
stress-dependent activation energy
Doktorska disertacija je izdelana pod mentorstvom red. prof. dr. Ladislava Kosca.
ANALIZA LEZENJA JEKLA 9Cr-1Mo-0,2V Z
UPO[TEVANJEM NAPETOSTNO ODVISNE
AKTIVACIJSKE ENERGIJE
UDK: 539.3:669.14.018
POVZETEK
Vedenja jekla 9Cr-1Mo-0,2V pri kratkotrajnih preiz-
kusih lezenja ni mogo~e opisati z navadnim poten~nim
zakonom Arrheniusovega tipa, saj je izmerjen napetostni
eksponent n, definiran kot (  ln  εmin/  ln )T, zelo odvisen
od temperature, izmerjena navidezna aktivacijska ener-
gija lezenja Qc, definirana kot (  ln  εmin/ (–1/RT)) , pa od
napetosti.
Problema opisa vedenja tega jekla pri lezenju smo se
najprej lotili z uporabo koncepta napetostnega praga
(threshold stress), to je z uvedbo mejne napetosti, pod
katero je lezenje zanemarljivo. Izpeljali smo ena~bo, ki
sicer dobro opi{e vedenje jekla pri lezenju, vendar pa se
je izkazalo, da je mejna napetost  th mo~no odvisna od
temperature in/ali uporabljene napetosti ter celo spre-
minja predznak. Nizka vrednost izra~unane navidezne
aktivacijske energije lezenja, ki je dosegla le 185 kJ
mol–1, kar je celo manj od aktivacijske energije za
samodifuzijo v Fe-, pa je v tem primeru irelevantna za
pojasnjevanje delujo~ih mehanizmov lezenja.
Za opis vedenja jekla 9Cr-1Mo-0,2V pri lezenju smo
zato teoreti~no razvili nov model poten~nega zakona
lezenja z napetostno odvisno energijsko pregrado. Glede
na standardni model temelji na{ na hipotezi, po kateri
delujo~a napetost neposredno vpliva tudi na aktivacijsko
Materiali in tehnologije / Materials and technology 42 (2008) 5, 227–233 227
DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES
AN ANALYSIS OF THE CREEP BEHAVIOUR OF
9Cr-1Mo-0.2V STEEL BASED ON THE
STRESS-DEPENDENT ACTIVATION ENERGY
UDC: 539.3:669.14.018
ABSTRACT
The creep behaviour of 9Cr-1Mo-0.2V steel during
short-term creep tests cannot be accurately described by
a simple Arrhenius-type power-law model because the
apparent stress exponent n, defined as (  ln  εmin/  ln )T,
shows a strong temperature dependence, whereas the
apparent activation energy of the creep Qc, defined as
(   ln  εmin/ (–1/RT)) , shows a strong stress dependence.
The problem of describing the creep behaviour of the
examined steel was first dealt with by applying the
threshold-stress concept, i.e., the introduction of the
threshold stress  th, below which the creep deformation
is assumed to be negligible. An equation that satisfac-
torily describes the creep behaviour of the examined
steel was derived. However, it was observed that the
threshold stress  th strongly depends on the temperature
and/or the applied stress, and that sometimes it can even
change its sign. The low value of the calculated
activation energy of the creep, which was only 185 kJ
mol–1, is considerably smaller than the activation energy
for self-diffusion in Fe-, and thus it is irrelevant for an
explanation of the acting creep mechanism.
For this reason, a power-law, stress-dependent
energy-barrier model was theoretically developed in
order to describe the creep behaviour of 9Cr-1Mo-0.2V
energijo oziroma na energijsko pregrado, ki jo je treba
pri aktiviranju mehanizmov premagati, in ne samo na
potencialno energijo za~etnega in kon~nega stanja
mikrostrukturnih obmo~ij, v katerih se lezenje odvija.
Nov model poten~nega zakona smo izbolj{ali {e z
upo{tevanjem linearne odvisnosti aktivacijskega
volumna od temperature ter z normiranjem napetosti z
uvedbo stri`nega modula. Prvi~ smo ta model v {e ne
povsem kon~ni obliki uspe{no uporabili za opis lezenja
jekla 9Cr-1Mo-0,2V na osnovi literaturnih podatkov,
nato pa {e za opis vedenja tega jekla pri kratkotrajnem
lezenju (ang. short term creep test) v tem delu, pri ~emer
smo primerjali razli~ne vrste preizku{ancev oz.
obremenitve.
Jeklo 9Cr-1Mo-0,2V smo toplotno obdelali po
standardni toplotni obdelavi, nato pa opravili enoosne
stati~no natezne preizkuse lezenja pri konstantni
obremenitvi oz. konstantni napetosti ter merili lezenje {e
z upogibanjem tankih diskov s centralno delujo~o
obremenitvijo (ang. small-punch test), ki v preizku{ancu
povzro~i ve~osno rotacijsko simetri~no ter izrazito
heterogeno napetostno in deformacijsko stanje. Pri
enoosnih nateznih preizkusih lezenja so bile temperature
lezenja 625 °C, 650 °C oz. 675 °C, napetosti pa od 120
MPa do 240 MPa, medtem ko so bile temperature
preizku{anja pri merjenju lezenja z upogibanjem tankega
diska 650 °C, 675 °C oz. 690 °C, obremenitve pa od 350
N do 550 N. Preizkusi merjenja lezenja z upogibanjem
tankega diska so bili opravljeni na diskih razli~nih
debelin, in sicer 0,44 mm, 0,47 mm oz. 0,50 mm.
Ugotovili smo zelo dobro ujemanje med eksperimen-
talnimi podatki in izra~unanimi vrednostmi po modelu
poten~nega zakona z napetostno odvisno energijsko
pregrado. Navidezna aktivacijska energija za lezenje
postane v na{em modelu le {e nekoliko odvisna od
napetosti. Pri enoosnih stati~no nateznih preizkusih
lezenja se s pove~anjem napetosti od 120 MPa do 240
MPa napetostno odvisna navidezna aktivacijska energija
lezenja Qc, izra~unana iz minimalne hitrosti lezenja  εmin
zni`uje od pribli`no 645 kJ mol–1 do 600 kJ mol–1,
izra~unana iz ~asa do poru{itve tr pa od pribli`no 580 kJ
mol–1 do 548 kJ mol–1. Napetostni eksponent n pa je v
obeh primerih enak, n = 4,5.
Pri merjenju lezenja z upogibanjem tankega diska je
bila eksperimentalno potrjena veljavnost Monkman-
Grantove odvisnosti med minimalno hitrostjo upogibanja
diskov  δmin in ~asom do poru{itve tr,SP. Zato smo lahko v
izbolj{anem modelu poten~nega zakona z napetostno
odvisno energijsko pregrado minimalno hitrost upogi-
banja diska  δmin ustrezno zamenjali s ~asom do poru{itve
tr,SP, saj sta obe koli~ini skoraj obratno sorazmerni.
Ugotovili smo, da je optimalna debelina diska odvisna
od geometrije naprave za merjenje lezenja in je v na{em
primeru med 0,48 mm in 0,50 mm. Izra~unana navidezna
aktivacijska energija lezenja Qc je nekoliko odvisna od
obremenitve in je za debelino diska 0,50 mm pri
obremenitvi 550 N enaka 543 kJ mol–1, pri obremenitvi
228 Materiali in tehnologije / Materials and technology 42 (2008) 5, 227–233
DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES
steel. In contrast to the standard model, this model was
based on the hypothesis that the application of stress also
affects the activation energy, i.e., the energy barrier that
needs to be overcome when a local region transforms,
and not just the potential energy of the initial and final
local states. The new power-law model was additionally
improved by applying the linearity between the
activation volume and temperature as well as by
normalising the applied stress with the shear modulus.
This model was then used in a not-completely-final
version for the first time to describe the already
published creep data of 9Cr-1Mo-0.2V steel and then
also for a description of the short-creep data in the
present work, where different types of test specimens
and loadings were compared.
The 9Cr-1Mo-0.2V steel underwent a standard heat
treatment. After this uniaxial static-tensile creep tests at
a constant load and a constant stress were conducted as
well as small-punch creep tests, which led to a
multiaxial, rotary-symmetric, heterogeneous stress and
deformation state in the sample. While the uniaxial
tensile creep tests were performed at temperatures of 625
°C, 650 °C and 675 °C and initial stresses ranging from
120 MPa to 240 MPa, the small-punch tests were
performed at temperatures of 650 °C, 675 °C and 690 °C
and at loads ranging from 350 N to 550 N. In the case of
the small punch the creep properties were measured on
disks with different thickness, i.e., 0.44 mm, 0.47 mm
and 0.50 mm.
A very good correlation between the power-law,
stress-dependent energy-barrier model and the experi-
mental data was obtained. However, the apparent
activation energy shows less stress dependence than
when using the Arrhenius-type power-law. During
uniaxial static-tensile creep tests the stress-dependent
activation energy Qc calculated from the minimum creep
rate  εmin decreased from approximately 645 kJ mol–1 to
600 kJ mol–1 as the stress increased from 120 MPa to 240
MPa, while the stress-dependent activation energy
calculated from the time-to-rupture tr decreased from
approximately 580 kJ mol–1 to 548 kJ mol–1. The stress
exponent n was the same in both cases, i.e., n = 4.5.
The validity of the Monkman-Grant relationship
between the minimum deflection rate  δmin and the
time-to-rupture tr,SP for the small-punch creep test was
experimentally confirmed. Thus, the minimum deflec-
tion rate  δmin in the improved power-law, stress-depen-
dent energy barrier model was replaced by the
time-to-rupture tr,SP, since both parameters are almost
inversely proportional. It was found that the optimum
disk thickness is dependent on the geometry of the
small-punch apparatus for measuring the creep
properties. In our case the optimum disk thickness was
found to be between 0.48 mm and 0.50 mm. The
calculated apparent activation energy of the creep Qc was
slightly load dependent. The value of the apparent
activation energy for a disk thickness of 0.50 mm was
350 N pa 533 kJ mol–1. Vrednost obremenitvenega
eksponenta nSP pa je pri vseh debelinah diska 4,5, kar je
enako kot pri standardnih enoosnih nateznih preizkusih
lezenja.
DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES
Materiali in tehnologije / Materials and technology 42 (2008) 5, 227–233 229
543 kJ mol–1 at a load of 550 N, and 533 kJ mol–1 at a
load of 350 N. The value of the stress exponent nSP was
4.5 for all the disk thicknesses and was the same as
during the conventional uniaxial tensile creep tests.
Na Fakulteti za elektrotehniko Univerze v Ljubljani je dne, 2.7.2008, pred
komisijo v sestavi prof. dr. Stanislav Kova~i~ in ~lani prof. dr. Drago Matko, prof.
dr. Borut Zupan~i~ in izr. prof. dr. Toma` Kolenko Franci Vode, univ. dipl. in`.
zagovarjal doktorsko disertacijo z naslovom:
Razvoj sistema vodenja kontinuiranih pe~i za ogrevanje vlo`ka po predpi-
sanih krivuljah ogrevanja
Development of continuous furnace control system for stock reheating
according to the prescribed reheating curves
Doktorska disertacija je izdelana pod mentorstvom red. prof. dr. Draga Matka.
RAZVOJ SISTEMA VODENJA KONTINUIRANIH
PE^I ZA OGREVANJE VLO@KA PO PREDPISA-
NIH KRIVULJAH OGREVANJA
UDK: 621.771:519.68
POVZETEK
V disertaciji je predstavljen splo{en koncept
kontroliranega ogrevanja vlo`ka v kontinuirnih pe~eh v
jeklarski industriji, predstavljen na primeru potisne pe~i
v podjetju ACRONI, d. o. o. Vodenje ogrevanja slabov
je izvedeno z regulacijskim sledenjem referen~ni krivulji
ogrevanja (RKO) vsakega posameznega slaba, pri ~emer
regulacija temelji na matemati~nem modelu ogrevanja
slabov, ki deluje v realnem ~asu. RKO je podana kot
funkcija ~asa. ^asovno podan potek RKO omogo~a
elegantno in natan~no obravnavo zastojev. Sistem
vodenja kot odziv na zastoj v proizvodnji premakne
~asovno definirano RKO za ~as trajanja zastoja. S tem je
omogo~eno zaprtozan~no vodenje ogrevanja slabov tudi
med zastoji, in sicer {e vedno po RKO. Od u~inkovite in
natan~ne obravnave zastojev je najbolj odvisna
u~inkovitost ter energetski prihranki na pe~i pri uporabi
razvitega sistema. Zastoje zato sistemati~no razdelimo v
pet razli~nih vrst. Najprej jih lo~imo na tiste, katerih
trajanja ne poznamo, in na tiste, katerih trajanje
poznamo. Zadnje delimo na predvidene, nepredvidene,
na zastoje zaradi razli~nih ~asov ogrevanja razli~nih
jekel oz. debelin jekel ter na zastoje zaradi sinhronizacije
z valjarskimi ogrodji. Slednji omogo~ajo upo{tevanje
kapacitet valjarskih prog pri na~inu ogrevanja v pe~i. Za
vsako vrsto zastoja dolo~imo ~asovno premaknitev RKO
in podalj{anje intervala pomika slabov. RKO so lo~eno
definirane za razli~ne za~etne temperature s korakom 50
°C. Za potrebe prehoda iz ro~nega v avtomatsko vodenje
so za~etne RKO posnetek stacionarnega poteka
ogrevanja, kakr{nega predpisujejo tehnolo{ka navodila o
ogrevanju. S tem se izognemo spremembi na~ina
ogrevanja ob uvajanju sistema v proizvodnjo. RKO je
mo`no kadarkoli optimizirati in spremembe vnesti v
tabelo RKO v podatkovni zbirki. Na koncu dela je
DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES
230 Materiali in tehnologije / Materials and technology 42 (2008) 5, 227–233
DEVELOPMENT OF CONTINUOUS FURNACE
CONTROL SYSTEM FOR STOCK REHEATING
ACCORDING TO THE PRESCRIBED REHEATING
CURVES
UDC: 621.771:519.68
ABSTRACT
In the Ph.D. thesis is presented general concept of
controlled stock reheating in continuous furnaces for
steel industry, presented in the pusher-type furnace in
ACRONI, d. o. o. Controlled stock reheating is achieved
by the use of tracing reference reheating curves (RRC),
where the control of stock temperature is based on
real-time mathematical model of slab reheating process.
RRC is defined as a function of time. Such a definition
of RRC enables an exact and elegant consideration of
delays. A response of furnace control system to the delay
in production line is a time shift of RRC for the value of
delay duration. This enables closed-loop control of the
slab reheating process during normal operation as well as
during delays. Efficiency and precision of furnace
control during delays is the main factor for overall
efficiency and energy savings of furnace operation.
Delays are therefore systematically divided into five
delay types. Firstly, delays are divided into those, for
which the duration is not known and for those with
known duration. The last are further divided on
scheduled, unscheduled, delays due to different reheating
times of different steel grades and delays due to
synchronization with rolling mills. The last delay type
enables consideration of rolling mills capacities at the
reheating process in the furnace. For each delay type is
determined a time shift of RRC and a prolongation of
drop out interval. RRC are separately defined for
different initial temperatures with a step of 50 °C. To
introduce the presented furnace control system (FCS),
initial RRC are imitation of stationary reheating process,
which is defined by technological guidelines of the
reheating in furnace. Using such a RRC for introducing
FCS the material is reheated in the same way as by
primerjava rezultatov ro~nega vodenja z rezultati
simulacije avtomatskega vodenja. Avtomatsko vodenje
omogo~a za faktor 3 tesnej{e sledenje kon~ni tem-
peraturi slabov in za faktor 2 tesnej{e sledenje celotni
temperaturni krivulji v primerjavi z ro~nim vodenjem.
DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES
Materiali in tehnologije / Materials and technology 42 (2008) 5, 227–233 231
manual control. RRC can be optimized any time and the
changes can be updated in the table of RRC in database.
A comparison of manual and simulation of automatic
control is presented at the end. Automatic control by
FCS enables three times tighter tracing of end material
temperature and two times tighter tracing of whole RRC
in comparison to manual control.
Na Naravoslovnotehni{ki fakulteti Univerze v Ljubljani je dne 20. 2. 2008 pred
komisijo v sestavi: red. prof. dr. Radomir Turk kot predsednik in ~lani: izr. prof.
dr. Anton Smolej, doc. dr. Jelena Vojvodi~ Tuma, red. prof. dr. Franc Vodopivec,
red. prof. dr. Ivan An`el in izr. prof. dr. Bo{tjan Markoli, Gorazd Kosec, univ.
dipl. in`. materialov in metalurgije zagovarjal doktorsko disertacijo z naslovom:
Krhki prelom v coni toplotnega vpliva zvarov jekla Niomol 490 K
Brittle fracture in the heat affected zone of welds of steel Niomol 490 K
Doktorska disertacija je izdelana pod mentorstvom izr. prof. dr. Antona Smoleja
in somentorstvom doc. dr. Jelene Vojvodi~ Tuma.
KRHKI PRELOM V CONI TOPLOTNEGA
VPLIVA ZVAROV JEKLA NIOMOL 490 K
UDK: 621.791.05:669.14.018.298
POVZETEK
Cilj na{ega dela je bil ugotoviti, zakaj je prehodna
temperatura `ilavosti po Charpyju zvarov tanj{ih plo{~
(15 mm) opazno vi{ja, kot pri zvarih debelej{ih (25 mm)
plo{~ jekla Niomol 490 K, ~eprav je bilo varjenje
opravljeno z isto tehnologijo in z enako elektrodo. Da bi
na{el razlago za to razliko sem pripravil z razli~nimi
toplotnimi obdelavami jeklo razli~nih mikrostruktur, ki
so se razlikovale tudi po velikosti kristalnih zrn.
Segrevanje pri nanosu naslednjega varka sem simuliral s
kratkim segrevanjem v dvofazno temperaturno podro~je,
kjer sta obstojna avstenit in ferit. Preizkusi so obsegali
merjenje `ilavosti po Charpyju v razponu od popolnoma
krhkega do popolnoma duktilnega preloma, meritve
trdot, natezne preizkuse pri temperaturi krhkega loma ter
preiskave mikrostruktur in prelomnih povr{in. Na
podlagi eksperimentalnih rezultatov in njihove analize
sklepam, da ima jeklo z mikrostrukturo iz spodnjega
bainita mnogo ve~jo `ilavost, kot jeklo z mikrostrukturo
martenzita. Po ogrevanju v dvofazno podro~je (750 °C)
se mo~no zmanj{a `ilavost pri temperaturi 0 °C in povi{a
prehodna temperatura `ilavosti v primeru, ko je bila
za~etna mikrostruktura iz bainita, pri jeklu z za~etno
mikrostrukturo iz martenzita pa so spremembe mnogo
manj{e in druga~ne. Vpliv pogrevanja je druga~en in
mnogo manj{i pri drugih lastnostih, razlaga pa je v
na~inu in v hitrosti obremenjevanja oz. meritve. Vpliv
ogrevanja na `ilavost in prehodno temperaturo je mogo~e
razlo`iti na podlagi dveh podmen, tako da pri
temperaturi ogrevanja ni neposredne premene martenzita
v avstenit, ampak nastane sekundarni martenzit v pro-
cesu martenzit   ferit + cementit (popu{~eni martenzit)
  avstenit   sekundarni martenzit, ~e pa je izhodni
mikrostrukturi iz ferita in cementita (bainit), pa je
proces, ko nastane sekundarni martenzit v naslednjih
stopnjah: ferit + cementit (bainit)   avstenit   sekun-
darni martenzit, ter da je porazdelitev ferita in sekun-
DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES
232 Materiali in tehnologije / Materials and technology 42 (2008) 5, 227–233
BRITTLE FRACTURE IN THE HEAT AFFECTED
ZONE OF WELDS OF STEEL NIOMOL 490 K
UDC: 621.791.05:669.14.018.298
SUMMARY
The aim of this investigation was to determine, why
for the the low carbon Nb microalloyed steel Niomol 490
K and for identical welding procedure the Charpy
transition temperature was higher for welds of thinner
(15 mm) than for welds of thicker (25 mm) plates. With
thermal treatment specimens of the steel were prepared
with a different microstructure and grain size and the
reheating at the next pass deposition was simulated with
short time annealing in the two phase austenite + ferrite
temperature of 750 °C. For all types of investigated
microstructure the investigations consisted of: the
determination of the dependence Charpy notch toughnes
versus temperature from below the lower shelf threshold
up to to 60 °C, the determination of the tensile strength
by brittle fracturing and hardness; SEM examination of
microstructure and of fracture surfaces and the deter-
mination of the space orientation of clevage facets of the
brittle fracture.
It was found that notch toughness was much higher
with lower bainite, also for the steel quenched from 1250
°C in lead bath at 400 °C, than for martensite. After
reheating, the transition temperature was increased and
the 20 °C notch toughness was decreased very much for
bainite, while, both vere virtually unchanged for
martensite. The effect of reheat is much lower for other
properties.
The different effect of reheating on notch toughness
can be explained assuming that:
– no direct tranformation martensite austenite takes
place and secondary martensite is formed with the
process: martensite   ferrite + cementite  
austenite   secondary martensite, while for bainite
the proces is bainite (ferrite + cementite)   austenite
  secondary martensite
– and the distribution of secondary martensite and
ferrite is identical or very similar to the distribution
darnega martenzita enaka, ali zelo podobna porazdelitvi
ferita in cementita pred premeno v avstenit. Morfologija
mikrostrukture, predvsem porazdelitev ferita in sekun-
darnega martenzita po ogrevanju najmo~neje vplivata na
`ilavost in prehodno temperaturo `ilavosti in to tem bolj,
~im ve~ja so bila kristalna zrna avstenita iz katere je
nastala primarna mikrostruktura. Mikrostruktura ne
vpliva na orientacijo oz. prostorsko lego cepilnih
ploskev, ki so  110 pri izhodni primarni mikrostrukturi
iz martenzita ali bainita in pri mikrostrukturah, ki so
nastale z ogrevanjem le-teh.
Lokalna krhka podro~ja (LBZ local brittle zone)
zmanj{ajo `ilavost in povi{ajo prehodno temperaturo
`ilavosti v pasu najve~jih kristalnih zrn avstenita v coni
toplotnega vpliva (CTV). Dokazal sem, da je premena
avstenita v bainit (primarni) pri nanosu predhodnega
varka mnogo bolj neugodna za `ilavosti po nanosu
naslednjega varka, kot ~e se avstenit v CTV premeni v
primarni martenzit. Zato je po primarni premeni avste-
nita v bainit mnogo verjetneje, da bo po nanosu nasled-
njega varka nastala LBZ, ki bo bistveno bolj neugodna
kot v primeru primarne mikrostrukture iz bainita, kot po
premeni v martezit, ~eprav je izhodna trdota ve~ja pri
martenzitu. Pri jeklu Niomol 490 K in njemu podobnih
jeklih, se je pri varjenju izogibati premene avstenita v
bainit v pasu najve~jih kristalnih zrn avstenita v CTV,
posebej {e, ~e lahko pri gradnji ve~varkovnega vara
nastanejo LBZ.
Notranje napetosti, ki so povezane z nastankom
sekundarnega martenzita (LBZ) povzro~e plastifikacijo
matice in so najve~je pri temperaturi okolice. Prispevajo
k nukleaciji in propagaciji razpoke pri udarnem preiz-
kusu, pa tudi k lokalni poru{itvi delcev sekundarnega
martenzita (LBZ), s ~imer {e olaj{ajo nukleacijo razpoke
pri udarnem preizkusu.
DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES
Materiali in tehnologije / Materials and technology 42 (2008) 5, 227–233 233
od ferrite and cementite before the transformation to
austenite at reheating.
The distribution of ferrite and austenite before the
reheating affects the notch toughness and notch
transition temperature the strongest, the coarse were the
austenite grain size before the formation of secondary
martensite. The type of the initial and reheat
microstructure does not affect the cleavage lattice plane
 110.
Local brittle zones (LBZ), which may form in the
heat affected zone of welds, decrease the notch tougness
and increase the transition temperature the most in the
layer of coarsest austenite grain size heat affected zone.
Thus, for the steel Niomol 490 K, the transformation
austenite to lower bainite is more harmfull for after
reheat notch toughness than the transformation austenite
to primary martensite, as the microstructure of bainite is
more propensive to the formation of LBZ than that of
martensite, although the hardness is higher for
martensite. It is consluded that by welding of the steel
Niomol 490 K and similar structural steels, it is
recommended to avoid the trasformation of coarse
austenite to lower bainite in the coarse grained part of
the heat affected zone.