International Journal of Refractory Metals & Hard Materials 24 (2006) 145–154
www.elsevier.com/locate/ijrmhm
Effect of TaC on plastic deformation ofWC–Co and Ti(C,N)–WC–Co
Gustaf Ostberg a,*, Katharina Buss b, Mikael Christensen a, Susanne Norgren c,Hans-Olof Andren a, Daniele Mari b, Goran Wahnstrom a, Ingrid Reineck c
a Department of Applied Physics, Chalmers University of Technology, SE-412 96 Goteborg, Swedenb Ecole Polytechnique Federale de Lausanne, Institut de Physique de la Matiere Complexe, CH-1015 Lausanne, Switzerland
c R&D Materials and Processes, AB Sandvik Coromant, SE-126 80 Stockholm, Sweden
Received 11 November 2004; accepted 12 April 2005
Abstract
The plastic deformation resistance of metal cutting inserts made from a WC–Co cemented carbide, a Ti(C, N)–WC–Co cermet
and corresponding materials with additions of TaC has been studied. The cermets were produced with both high and low carbon
activity, resulting in a total of six materials. Ab initio calculations of some WC/WC grain boundary geometries suggest that both Co
and Ta segregate substitutionally to the boundary and improve the grain boundary strength when substituting carbon. However,
only Co segregation was found experimentally, probably due to (Ta, W)C formation. Plastic deformation tests were performed with
a turning operation under controlled conditions. For the WC–Co, the addition of Ta had a positive effect for lower cutting speeds
but at higher speeds the effect was negative. Three-point bending tests indicated a beneficial effect of Ta in WC–Co, which was also
confirmed by internal friction (IF) measurements. However, after thermal cycling, the effect of Ta could be smaller, or even negative.
The Ta cermet produced with low carbon activity exhibited a better plastic deformation resistance during cutting but no apparent
effects of Ta could be seen either in IF measurements or in three-point bending tests of the cermets. However, a correlation was
found between plastic deformation during turning and IF spectra. In the cermet materials, binder phase lamella formation promotes
grain boundary sliding at high temperatures.
� 2005 Elsevier Ltd. All rights reserved.
Keywords: Cemented carbides; Characterisation; Plastic deformation; DFT; Internal friction
1. Introduction
Cemented carbides are composite materials with a
hard carbide or carbonitride skeleton embedded in a
tough binder metal and are commonly used in metal cut-
ting since they have an outstanding ability to resist high
temperatures and loads.
In metal cutting industry, the largest economical sav-
ings are usually made by increasing the number of pro-
0263-4368/$ - see front matter � 2005 Elsevier Ltd. All rights reserved.
doi:10.1016/j.ijrmhm.2005.04.010
* Corresponding author. Tel.: +46 31 772 3325; fax: +46 31 772 3224.
E-mail address: [email protected] (G. Ostberg).
duced parts per time unit. This means that cutting
speeds have to be increased, resulting in high tool loadsand working temperatures, which can be as high as
1000 �C. Due to the excellent resistance to abrasive wear
of modern wear resistant coated tools, these severe cut-
ting conditions can cause the material to deform plasti-
cally before any significant wear by other mechanisms
has occurred. Thus, it is often the deformation of the
tool, rather than the abrasive wear at its surface, that
determines its lifetime.Another study of plastic deformation of cemented
carbide cutting tools [1] has shown that the hard phase
skeleton is broken up during deformation and films of
146 G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 145–154
binder phase form between the grains and facilitate
grain boundary sliding.
Over the years different measures have been taken to
improve the deformation resistance of cemented car-
bides and cermets. One example is to reduce the grain
size, and thereby increase the hardness, by adding tran-sition metal carbides, such as VC and Cr3C2 [2–4]. In
cermets, Mo2C has been commonly used due to its
favourable effect on the wetting between the ceramic
skeleton and the binder phase which improves the
mechanical properties of the material [5]. Different
studies of mechanical properties have also been per-
formed on materials where Mo2C was added [6–10].
Other cubic carbides like TiC and TaC have been usedto increase the hardness of WC–Co since they have a
higher hardness than WC [11]. The advantageous effect
of TaC on cermets has also been reported by Rolander
et al. [12] who found that tantalum increased the plas-
tic deformation (PD) resistance during metal cutting.
It was suggested that tantalum should influence the
interfacial energies, resulting in a stronger hard phase
skeleton.The purpose of this study, which can be considered as
a continuation of the work of Rolander et al., is to get a
fundamental understanding of how Ta affects the micro-
structure, and thereby the PD resistance, of cemented
carbides and cermets. Plastic deformation of cutting in-
serts by turning tests under realistic conditions are com-
pared to internal friction (IF) and three-point bending
measurements under controlled conditions. The micro-structure is characterised with scanning electron micros-
copy (SEM) and transmission electron microscopy
(TEM) and the chemistry is analysed with energy dis-
persive X-ray analysis (EDX). Furthermore, first princi-
ple ab initio calculations are made on the WC–Co
system to predict segregation tendencies and interfacial
energies.
2. Materials and experimental procedures
Six model alloys were studied in this work and their
compositions after sintering are listed in Table 1. The
cermets were produced in two versions; one with high
(HC) and one with a lower carbon (LC) activity. All
Table 1
Composition as determined by chemical analysis after sintering (at%) of the
Alloy Co Ta
WC–Co 9.71 0.00
WC–Co–Ta 9.79 1.06
Ti(C, N)–WC–Co–HC 9.13 0.00
Ta–Ti(C, N)–WC–Co–HC 8.80 2.10
Ti(C, N)–WC–Co–LC 9.27 0.00
Ta–Ti(C, N)–WC–Co–LC 9.23 2.02
materials were designed to have the same binder volume
fraction of 10.2% at sintering temperature and the W-
solution in the Co-binder was to be comparable for
the cemented carbides and the HC cermets, resulting
in the unusually high carbon additions for a cermet.
Ta was added in amounts above the solubility limit inorder to avoid coarse sluggish precipitation of Ta into
the cubic (Ta, W)C phase in the cemented carbide. In-
stead, the nuclei of the cubic carbide can be retained
in the liquid at sintering temperature in order to precip-
itate the cubic phase formed during cooling on an al-
ready present phase. Moreover, the cermets were made
with as similar Ti/W (�10) and N/(C + N) (�0.37) ratios
as possible. The alloys were produced by powder metal-lurgical methods including mixing by milling, spray-dry-
ing, pressing and sintering. The materials were pressed
as triangular cutting inserts for the deformation testing
and as bars for the preparation of IF and three-point
bending specimens. Sintering temperatures were
1410 �C for the cemented carbides and 1480 �C for the
cermets. After sintering, the cutting inserts were coated
with a 5 lm TiC/Ti(C, N) layer to minimize abrasivewear. Some basic properties of the sintered materials
are shown in Table 2.
2.1. Turning tests
Plastic deformation (PD) of the materials was mea-
sured as the depression of the rake face of the insert after
a radial turning (facing) operation under controlled con-ditions. A cylindrical workpiece made from SS 2541
(0.34 wt% C, 1–1.5 wt% Cr, 3 wt% Ni, Mo) steel was
cut with a depth of cut of 1 mm and a feed of 0.3 mm/
rev. The cutting speed was kept constant during each
test and the cutting time was 30 s. Tests were performed
at cutting speeds between 300 and 475 m/min and for
every cutting speed a new cutting insert was used. Two
of such test sequences were performed for all six materi-als. In Fig. 1, the average plastic deformation of the two
test sequences are plotted for each material.
2.2. Electron microscopy
All scanning electron microscopy was performed with
a Leo Ultra55 FEG-SEM equipped with an Oxford
six model alloys studied in this work
W Ti N C
45.49 0.00 0.00 44.80
44.25 0.00 0.00 44.90
3.92 42.08 16.41 28.46
3.85 40.04 14.71 30.49
3.83 42.33 16.72 27.86
3.93 40.13 16.34 28.35
Table 2
General properties of the six alloys studied
Alloy Density
(g/cm3)
Porosity Grain sizea
(lm)
HV3 Co-magnetic W in ss
(at%)bCoercivity
(lT/cm3)
WC–Co 14.97 A02B00C00 2.10 1493 5.10 3.12 13.12
WC–Co–Ta 14.76 A06B08C00 2.59 1405 5.18 2.85 13.31
Ti(C, N)–WC–Co–HC 6.34 A04B02C00 2.05 1594 11.63 4.34 13.96
Ta–Ti(C, N)–WC–Co–HC 6.64 A02B02C00 2.52 1473 11.47 3.50 12.03
Ti(C, N)–WC–Co–LC 6.3 A04B04C00 Not measured 1468 11.34 4.77 13.92
Ta–Ti(C, N)–WC–Co–LC 6.67 A04B02C00 Not measured 1478 9.46 6.62 13.82
a Measured by mean linear intercept method.b Calculated from the Co-magnetic measurements.
Plastic deformation
0
0.2
0.4
0.6
0.8
1
1.2
300 350 400 450 475Cutting speed (m/min)
PD (m
m)
WC-Co Hi C cermet Lo C cermetWC-Co + Ta Hi C cermet + Ta Lo C cermet + Ta
Fig. 1. Results from plastic deformation tests. Every point in each
curve is an average of two measurements.
G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 145–154 147
INCA system for EDX analysis. TEM analyses were
performed in a Philips CM200 FEG-TEM equipped
with a Link Isis� EDX system.
SEM samples were prepared by cutting cross-
sections through the deformed cutting edge as is shown
in Fig. 2. Further details on the procedures for material
Fig. 2. SEM micrograph showing a top-view of the deformed cutting
edge of an insert. The highlighted area marks the material removed
when making the cross-section.
analysis and sample preparation are described elsewhere
[1].
2.3. Three-point bending
Three-point bending tests were performed at a con-
stant strain rate of _e ¼ 1.5� 10�5 s�1. Carbide skeleton
samples were produced by etching the cobalt. Bending
tests were performed on complete and skeleton samples,
with size 35 · 7 · 3.5 mm3, at different temperatures be-
tween 800 and 1200 �C. Details about the procedure aregiven elsewhere [1].
2.4. Mechanical spectroscopy
Mechanical spectroscopy is used to determine the IF
which is a measure of the dissipative processes, e.g., dif-
fusion and dislocation movements, acting at different
temperatures. The IF was measured both at constantfrequency, as a function of temperature and at constant
temperature as a function of frequency on bars,
35 · 4 · 1 mm3 in size. A more thorough description of
the method can be found elsewhere [1].
3. Ab initio calculations—method
Ab initio calculations were performed for ten differ-
ent translation states in the WC(0001)/WC(1�210)asymmetric tilt boundary, shown in Fig. 3. In five of
the translational states, the close packed (0001) plane
interface is metal terminated, and in the other five it is
carbon terminated.
In the calculations of the energetics, only the internal
energy given by the total energies obtained from densityfunctional theory (DFT) is taken into account. This is the
dominating contribution to the interface energetics, and
all other terms are neglected based on the assumption
of them being small [13] together with a tendency for can-
cellation of temperature-dependent terms for differences
in free energy between the considered structures [14].
Further details about the calculations are given else-
where [1,15].
Fig. 3. Relative position of interface atoms at all studied grain
boundary geometries containing segregated (Co, Ta) metal atoms.
Segregated atoms are white, C atoms black, and W atoms grey. The
atoms in two close packed (0001) layers, and one (1�210) layer are
displayed. Atoms in the same layer are connected.
148 G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 145–154
4. Results
4.1. Turning tests
From Fig. 1, it can be seen that the PD resistance of
WC–Co is unaffected or slightly improved by the Ta
addition at cutting speeds up to 450 m/min but at higher
speeds the Ta containing material undergoes larger
deformation.
For the cermets the behaviour is very different
between the high and low carbon materials. The PDresistance for the HC cermet is strongly impaired by
the addition of Ta whereas for the low carbon material
it is clearly improved over the whole range of measure-
ment.
4.2. Characterisation
In Fig. 4, the microstructures of the undeformedmaterials are compared. At a first glance, the morphol-
ogies of the cemented carbides with and without Ta look
quite similar. However, except for the presence of cubic
(W, Ta)C grains the WC grains (a-phase) are also
slightly larger and have a less faceted grain shape in
the Ta material. Some porosity and a less homogeneous
binder phase distribution can also be seen when Ta is
added.When comparing the morphologies of the HC cer-
mets, quite large differences can be seen between the
Ta and no-Ta material. The fraction of rim area is high-
er and the fraction of cores is lower in the hard phase
grains (c-phase) of the Ta material. Furthermore, the
grains are more rounded and the grain size is signifi-
cantly larger. To explain this, thermodynamic calcula-
tions of the effect of Ta on the melting point of thebinder phase were made. In Fig. 5, the melting temper-
ature is plotted as a function of carbon activity during
sintering and it can be seen that Ta lowers the melting
temperature overall and also makes the melting point
more sensitive to variations in carbon activity. Obvi-
ously, the liquid state sintering has been taken to a later
stage due to the lowered melting point of the binder
phase, which explains the morphology of the HC Ta cer-met. Thus, a direct comparison should not be done be-
tween the turning test results of the high carbon
cermets, since the two materials represent completely
different microstructures.
In the LC cermets, there are small differences between
the morphologies of the Ta and no-Ta materials (Fig.
4(e) and (f)). The most obvious difference is that the rims
of the hard phase grains in the Ta cermet have higherbrightness, indicating a higher solution of heavy
elements.
TEM-EDX linescans of grain and phase boundaries
in the Ta materials are shown in Fig. 6. The EDX probe
size is 8 nm at FWHM and, if considering beam broad-
ening effects in the specimen, the total probe size should
be slightly larger than 10 nm at FWHM. Hence, in Fig.
6(a) and (d) there are small overlaps of the spots. A clearindication of segregation of Co can be seen both in the
a/a boundaries and the c/c boundaries. However, nei-
ther in the grain boundaries, nor in the Co/WC phase
boundaries, any signs of Ta segregation or enrichment
could be seen. It should be noted that the presence of
Co in the hard phase grains is caused by background
noise and is only apparent. The Ta content in WC is also
just apparent due to the overlap between the Ta and Wpeaks in the EDX spectrum. Hence, what is seen as Ta
in the WC is just the tail of the W peak.
In Fig. 7 the six materials are shown after deforma-
tion by turning testing. All materials exhibit a partly
broken up hard phase skeleton where binder phase has
infiltrated some grain boundaries and formed lamellae
between the grains. This process has been described pre-
viously [1].In Fig. 8, the deformed region of the Ta cemented
carbide has been outlined. Apparently, the porosity
which can be seen in the undeformed bulk is absent.
A noticeable difference seen in the TEM between the
Ta and no-Ta HC cermet is that there appears to be a
larger misfit between the cores and the rims in the hard
phase grains (see Fig. 9). The deformed HC Ta cermet
also seems to have wider and more frequently occurringbinder lamellae than the no-Ta material. However, this
is most likely just a sign of the higher degree of deforma-
tion rather than a direct effect of the Ta addition.
4.3. Three-point bending
Stress–strain curves at different temperatures for the
cemented carbide and the cermets are shown in Fig.10(a) and (b), respectively, for comparison between the
no-Ta and Ta materials. For the WC–Co, the addition
Fig. 4. SEM micrographs of the undeformed materials. (a) WC–Co, (b) WC–Co with Ta, (c) high carbon cermet, (d) high carbon cermet with Ta,
(e) low carbon cermet and (f) low carbon cermet with Ta. The somewhat uneven contrast of the binder phase is due to topographic effects.
G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 145–154 149
of Ta leads to an increase of the flow stress at 900 and
1000 �C. As expected, the fracture strain decreases
whereas the flow stress increases. However, at 1200 �C,the flow stress of the material with Ta is equivalent tothat of the WC–Co without Ta.
The cermets show surprising results in three-point
bending. Despite the significant differences that they dis-
play in the turning test results, all four cermets perform
very similarly at all temperatures in three-point bending.
Apparently, three-point bending is not sensitive enough
for the mechanisms that lead to the deformation of the
cutting edge or the test conditions are too different toallow a comparison.
4.4. Mechanical spectroscopy
In general, the WC–TaC–Co shows a lower IF than
WC–Co and the IF is clearly reduced all along the tem-
perature spectrum (Fig. 11). However, at high tempera-
ture, the structure with tantalum seems more unstable.
Whereas the IF of WC–Co without Ta is very stable,
i.e., heating and cooling are very similar, the IF of
WC–TaC–Co increases at high temperature and a hys-
teresis is formed above 1125 �C. Moreover, it can be ob-served that the IF continues to increase at high
temperature upon repeated thermal cycling and above
1180 �C the IF of WC–TaC–Co becomes even higher
than that of WC–Co. In fact, the increase of IF at high
temperature in the Ta containing sample is due to the
presence of a peak named PW4, which cannot be seen
in the temperature spectra at 1 Hz [1]. As can be seen
in Fig. 12 the maximum of this peak is located at1125 �C at 10�2 Hz and PW4 is enhanced by the presence
of Ta.
The general temperature spectrum of the cermets
resembles very much the one formerly presented [16].
However, at the maximum temperature of the tempera-
ture scans, the IF increases during a 2 h annealing. This
increase of IF is different for the four cermets. A cycle
composed of a heating and cooling scan (Fig. 13) leads
Fig. 5. Thermodynamical calculations showing the effect of Ta content
on the melting temperature of the binder phase as a function of carbon
concentration.
150 G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 145–154
to the formation of a hysteresis, the area of which is very
different depending on the type of cermet.
WC/WC boundary in WC-TaC-Co
0
1
2
3
4
5
6
7
8
9
10
0 10 20 30 40 50 60 70 80 90 100 110
Distance (nm)
Distance (nm)
Com
posi
tion
(at%
)
Atomic% Co KAtomic% Ta L
Com
posi
tion
(at%
)
Com
posi
tion
(at%
)
(Ti,Ta,W)(C,N)/Co boundary in HC Ta cermet
0
10
20
30
40
50
60
70
80
90
100
0 10 20 30 40 50 60 70 80
Atomic% Co KAtomic% Ta LAtomic% W LAtomic% Ti K
a b
dc
Fig. 6. TEM-EDX linescans of (a) a WC/WC boundary in WC–TaC–Co,
boundary in the HC Ta cermet and (d) a (Ti, W, Ta)(C, N)/Ti(C, N) bounda
grains is an artifact due to background noise. In (d) Ti is used as balance.
4.5. Ab initio calculations
The heat of segregation to WC/WC grain boundary
has been calculated for Co and Ta. In each case, the seg-
regant is replacing a tungsten (carbon) atom in the
(1�210) interface plane in tungsten (carbon) terminatedgrain boundaries (Fig. 3). The results for segregation
to all studied grain boundary geometries as well as seg-
regation to free WC surfaces are given in Table 3. Posi-
tive values indicate situations where segregation is
energetically favourable.
It can be seen that segregation to free WC surfaces is
favourable for Co, but not for Ta. The propensity for
Co to adsorb on the WC surface can be related to thegood wetting of Co on WC.
For segregation to WC/WC grain boundaries, it is
found that Ta will not segregate to any tungsten termi-
nated boundaries, while it might be favourable for Co to
segregate. Segregation to carbon terminated boundaries
is highly favourable for both Co and Ta, although the
tendency for segregation is larger for Co.
The effect of segregated intergranular atoms on thegrain boundary strength has also been investigated.
The strength is taken to be the ideal work of separation,
Wsep, calculated as the difference between the sum of
Distance (nm)
Distance (nm)
WC/Co boundary in WC-TaC-Co
0
10
20
30
40
50
60
70
80
90
100
0 20 40 60 80 100 120 140
Com
posi
tion
(at%
) Atomic% Co KAtomic% Ta LAtomic% W L
(Ti,Ta,W)(C,N)/Ti(C,N) boundary in HC Ta cermet
0
1
2
3
4
5
6
7
8
9
10
0 20 40 60 80 100 120 140
Atomic% Co KAtomic% Ta LAtomic% W L
(b) a Co/WC boundary in WC–TaC–Co, (c) a Co/(Ti, Ta, W)(C, N)
ry in the HC Ta cermet. Note: The presence of Co and Ta in the WC
Fig. 7. SEM micrographs of the plastically deformed materials. (a) WC–Co deformed at 475 m/min, (b) WC–Co with Ta deformed at 475 m/min, (c)
high carbon cermet deformed at 500 m/min, (d) high carbon cermet with Ta deformed at 450 m/min, (e) low carbon cermet deformed at 400 m/min
and (f) low carbon cermet with Ta deformed at 450 m/min. Some examples of binder phase lamellae have been marked by white arrows. The direction
of the total macroscopic force is marked with black arrows. Note that the somewhat uneven contrast of the binder phase is due to topographic
contrast.
Fig. 8. SEM micrograph of the deformed WC–TaC–Co. The image
shows an overview of a cutting edge cross-section. The porosity (black
spots) seen in the bulk region is absent in the deformed region.
G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 145–154 151
energies of the free cleavage surfaces and the interface
energy of the intact boundary.
The results for the change in Wsep due to the presence
of segregants are given in Table 4. It is found that both
elements probably have a detrimental effect on the
boundary strength when present in tungsten terminatedboundaries. In contrast, both Co and Ta have a very
large strengthening effect in carbon terminated bound-
aries, where the effect is somewhat larger for Ta.
5. Discussion
5.1. WC–Co
In line with the results from the cutting tests, the
three-point bending measurements show an improved
resistance to plastic deformation of the WC–Co
Fig. 9. TEM micrographs of typical hard phase grains in the undeformed high carbon cermet without Ta (a) and with Ta (b). The material with Ta
apparently has more misfit dislocations in the interface between the core and the rim (marked with arrows).
Fig. 10. The effect of Ta on three-point bending deformation of (a) the
WC–Co and (b) the cermets at different high temperatures.
Fig. 11. IF heating–cooling cycle between RT and 1480 K of WC–Co
and WC–TaC–Co (at 1 Hz). Heating and cooling are indicated by
arrows.
Fig. 12. Frequency scans of WC–Co and WC–TaC–Co showing the
increase of PW4 upon the Ta addition.
152 G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 145–154
with Ta at intermediate temperatures (up to 1000 �C).However, this positive effect has disappeared at
1200 �C. Also in IF measurements, the tantalum seems
to have different effects on WC–Co at low and high
Fig. 13. IF heating–cooling cycle between RT and 1207 �C of all four
cermet grades (at 1 Hz). Heating is the lower and cooling the upper
curve of each hysteresis. During isothermal measurements at 1207 �Cthe IF continues to increase.
Table 3
Heat of segregation to free WC surfaces for Co and Ta
Surface Esegr (eV/atom)
Co Ta
WC(1210)-SubW 0.42 �0.09
WC(1210)-SubC 0.04 �1.07
Grain boundary
W-term (a) �0.06 �0.05
W-term (b) 0.26 �0.53
W-term (c) �0.22 �0.28
W-term (d) 0.60 �0.79
W-term (e) �0.52 �0.43
C-term (a) �0.52 �1.79
C-term (b) 2.38 2.36
C-term (c) 1.31 0.69
C-term (d) 2.03 1.52
C-term (e) 1.35 0.17
Table 4
Change in Wsep due to the presence of segregants
Grain boundary DWsep (J/m2)
Co Ta
W-term (a) �0.53 0.05
W-term (b) �0.18 (�0.48)
W-term (c) (�0.71) (�0.20)
W-term (d) 0.19 (�0.77)
W-term (e) �0.27 (�0.37)
C-term (a) (�0.62) (�0.79)
C-term (b) 2.58 3.79
C-term (c) 1.40 1.94
C-term (d) 2.20 2.86
C-term (e) 1.44 1.37
Figures in parenthesis mean that segregation does not occur for any
value of the carbon potential.
G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 145–154 153
temperatures. It has been shown that the PW4 peak is re-
lated to the infiltration of WC grain boundaries by the
cobalt, which thereby enhances grain boundary sliding
[1]. PW4, and possibly boundary infiltration, are en-
hanced by the presence of Ta.
Although predicted by the ab initio calculations, the
Ta does not segregate to the grain boundaries which
may be explained by that the driving force for formingcubic (Ta, W)C grains exceeds that of segregation. Thus,
the resistance to grain boundary infiltration, as calcu-
lated by the ab initio simulations, is governed by just
the Co segregation and the explanation of the positive
effect of Ta has to be found somewhere else than in
the a/a grain boundaries.
The initial, slightly positive, effect of Ta may be due to
the higher hardness of the cubic carbide grains whichmakes them less likely to deform and, thus, the deforma-
tion of the skeleton is hindered at first. However, at high-
er cutting speeds (temperatures) the effect of Ta is instead
negative, indicating that the c/a grain boundaries slide,
or get infiltrated by Co, easier than the a/a boundaries.
In the light of the above discussion, it can be expected
that slightly higher TaC contents will increase the
amount of the cubic carbides. Though this may improvethe hardness it will also increase the brittleness which re-
sults in earlier fracture of the inserts. If the TaC content
is much higher there will be formation of grains with a
core/rim structure and a direct comparison with the
materials studied here should not be made.
5.2. Cermets
The presence of Ta combined with the high carbon
content (largest IF hysteresis) seems to lead to an in-
creased sensitivity to cobalt infiltration of the grain
boundaries, which is seen in the SEM by the wider
and more frequent occurring lamellae.
None of the cermets show any significant differences
in bending, which is surprising in view of the differences
displayed by the turning tests. However, a correlationcan be found between the area of the hysteresis in the
IF temperature spectra (Fig. 13) and the wear of the
tools. From the IF point of view, the hysteresis can be
associated with a changing structure. The IF even
changes under isothermal conditions at 1207 �C (Fig.
11). Just as for the WC–Co, the IF increase at high tem-
perature is correlated to the presence of a peak PT4, due
to the Co infiltration of grain boundaries. The isother-mal increase of IF leading to the hysteresis is then re-
lated to a progressive infiltration of grain boundaries.
When the hysteresis is large, the high temperature infil-
trated state tends to maintain upon cooling.
It should be noted that at 1200 �C considerable solid
state sintering occurs. Microstructural coarsening by
dissolution/reprecipitation is therefore expected. Possi-
bly this is also an additional mechanism of plastic defor-mation, although the time available during a turning test
is much shorter.
154 G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 145–154
The plastic deformation measured after turning at
different cutting speeds, where the cutting speed can
be, even if only qualitatively, related to the temperature
at the cutting edge, seems correlated with IF spectra ob-
tained as a function of temperature. In contrast to the
three-point bending, where the experiments are carriedout after a time of temperature stabilization of about
45 min and which shows similar results for all cermet
grades, the turning application involves changing tem-
perature conditions. The differences observed for the
four cermets could then be related to these conditions
in correspondence with the IF results.
A positive influence of the Ta on the deformation
behaviour can then only be achieved when the carboncontent is carefully adjusted. With too high carbon con-
tent, the Ta does not have a positive effect.
The compositions of the grades studied here were
chosen to be comparable as regards solid solution of
W, grain size etc and have not been optimised as regards
mechanical properties. Thus, an increase in the TaC may
lead to an increase of the plastic deformation resistance
if the sintering temperature is adjusted accordingly, tocompensate for the lowered melting point of the binder
phase. However, for much higher contents of TaC, hea-
vy cores will start to form which may increase toughness
and tool life [17], but not necessarily resistance to plastic
deformation.
6. Conclusions
• Ab initio calculations predict that Ta and Co segre-
gate substitutionally to WC/WC grain boundaries
and thereby increase the work of separation for the
boundaries.
• Co segregation was found by TEM-EDX at the grain
boundaries but no signs of Ta segregations could be
seen, neither in the WC–TaC–Co nor in the cermets,but in the WC–TaC–Co cubic (W, Ta)C grains are
formed.
• In the WC–Co, Ta has different effects at different
temperatures. The initial slight positive effect is
explained by the strengthening effect of harder
(Ta, W)C grains on the carbide skeleton. The nega-
tive effect at higher temperatures is explained by slid-
ing or infiltration of c/a grain boundaries.• In the cermet materials, the presence of Ta in the bin-
der phase makes the its melting point very sensitive to
the C activity during sintering. Hence, the same sin-
tering cycle gives different microstructure for the
HC and LC cermets.
• No difference in deformation resistance is seen by
three-point bending for the four cermets, whereas
the deformation resistance of the inserts during turn-
ing show substantial differences. IF heating curves for
each of the cermets exhibit a hysteresis which can be
attributed to binder phase infiltration of grain bound-
aries. The area of the hysteresis can be correlated to
the deformation achieved by the turning tests.
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