Mechanisms of plastic deformation of WC–Co and Ti(C, N)–WC–Co

International Journal of Refractory Metals & Hard Materials 24 (2006) 135–144

www.elsevier.com/locate/ijrmhm

Mechanisms of plastic deformation of WC–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 deformation of the cutting edge of inserts made from WC–Co and Ti(C, N)–19%WC–Co upon radial turning was measured

as a function of the cutting speed. Ab initio calculations of some grain boundary geometries in WC–Co indicated that Co segregates

to WC/WC grain boundaries in submonolayer proportions and increases their strength and resistance to Co infiltration. This was

also confirmed with TEM-EDX. SEM studies showed that during plastic deformation the hard phase skeletons in both materials

were partly broken up and infiltrated by binder phase. TEM observations showed a considerable deformation of the WC grains

and some infiltrated grain boundaries were facetted along low energy planes. In the cermet, no deformation of the hard phase grains

could be seen. The plastic deformation at high temperatures observed by three-point bending corresponds to a number of relaxation

processes, detected by internal friction, related to grain boundary sliding. Both materials deform by grain boundary sliding, accom-

modated in the cermet by Co diffusion and in the WC–Co by deformation of the hard phase and diffusion of Co.

� 2005 Elsevier Ltd. All rights reserved.

Keywords: Cemented carbides; Characterisation; Plastic deformation; DFT; Internal friction

1. Introduction

The continuous development of cemented carbidesused for metal cutting inserts has led to materials that

have an outstanding resistance to different wear mecha-

nisms. Grain size, carbide mixture, binder fraction and

composition can be controlled to design materials for al-

most every possible working condition. The deposition

of different coatings and gradient sintering are further

examples of production methods that have made it pos-

sible to increase cutting speeds, and thereby productiv-ity. However, since the inserts are very resistant to

abrasive wear, the gradually aggravated working condi-

0263-4368/$ - see front matter � 2005 Elsevier Ltd. All rights reserved.

doi:10.1016/j.ijrmhm.2005.04.009

* Corresponding author. Tel.: +46 31 772 3325; fax: +46 31 772 3224.

E-mail address: [email protected] (G. Ostberg).

tions can cause temperatures as high as 1000 �C, andeven above [1,2], in the tool materials and, in combi-

nation with very high loads, the tools can deformplastically.

A number of studies on the behaviour of hardmetals

at high temperatures (800–1200 �C) have been presented

[3–12] but common for most investigations is that they

are performed under laboratory conditions with meth-

ods like three-point bending and tensile creep tests.

In this work, five different aspects to study material

properties have been combined:

• Deformation of cutting inserts with a turning opera-

tion under realistic conditions has been obtained. The

degree of plastic deformation has been studied as a

function of cutting speed with the purpose of describ-

ing the deformation mechanisms that are acting

mailto:[email protected]

0

50

100

150

200

250

300

350

400

200 250 300 350 400 450 500 550

Cutting speed (m/min)

Plas

tic d

efor

mat

ion

(µm

)

WC-CoCermet

Fig. 1. Results from plastic deformation tests. Every point on each

curve is an average of six tests.

136 G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 135–144

under service conditions. However, although the cut-

ting conditions are controlled there is a lack of

knowledge about the temperature with this method.

• Three-point bending tests have been performed on

the same materials to measure the macroscopic defor-

mation mechanisms under controlled conditions.• Internal friction (IF) measurements have been used to

probe the mobility of microstructural defects and to

assess physical parameters linked to the thermal acti-

vation of relaxation processes.

• The microstructure has been characterized with scan-

ning electron microscopy (SEM) and transmission

electron microscopy (TEM). SEM was used to get a

good overview of the microstructure whereas TEMprovided the means to study the microstructure with

very high resolution as well as the crystal structure.

By comparing the microstructure on different length

scales before and after deformation, it is possible to

determine what mechanisms have operated.

• Ab initio calculations have been performed on the

WC–Co system. By using first principle density func-

tional theory (DFT) calculations it has been possibleto predict properties like grain boundary strength and

tendency for elements to segregate to grain

boundaries.

2. Materials and experimental procedures

Two model materials were studied. One was a cermet

made from a 0.8 lm (FSSS) WC powder, two 1.5 lmTi(C, N) powders and a 1.3–1.7 lm Co powder, which

were mixed by milling. The other material was a cemen-

ted carbide made from a 5.0 lm WC powder and a

0.8 lm Co powder. Both materials were designed to

have the same binder phase volume fraction of

10.19 vol% at sintering temperature. The Co-magneticvalues were adjusted to target the same amount of

W-dissolution in the binder, which means unusually

low solution hardening by W in Co for the cermets.

After milling and spray drying, the powders were

pressed as triangular cutting inserts, TNMG 160408-

PF. In addition, bars for the preparation of IF and

three-point bending specimens were manufactured from

the same powders. Vacuum sintering was performed incycles with a maximum temperature of 1410 �C for

WC–Co and 1480 �C for the cermet, both with an iso-

thermal hold for 1 h at sintering temperature. Finally,

the cutting inserts were PVD coated with a 5 lm Ti(C,

N)/TiN layer to minimize abrasive wear.

2.1. Turning tests

Plastic deformation of the materials was obtained by

a radial turning (facing) operation under controlled con-

ditions. A cylindrical workpiece made from SS 2541

steel (0.34 wt% C, 1–1.5 wt% Cr, 3 wt% Ni, Mo) was

cut with a depth of cut of 1 mm and a feed of 0.3 mm/

rev with a cutting time of 30 s. During each test the cut-

ting speed was kept constant and tests were performed

for cutting speeds from 250 to 550 m/min, to get an

increasing degree of deformation, until failure. Six in-serts were tested for each cutting speed and material

and their average plastic deformation, measured as

the depression of the rake face, is plotted in Fig. 1. No

values are plotted for 550 m/min since fracture was

obtained at this speed.

2.2. Scanning electron microscopy

All scanning electron microscopy was performed with

a Leo Ultra55 FEG-SEM equipped with an Oxford

INCA system for EDX analysis. Before the deformed

inserts were cut up, low magnification studies of the

shape of the inserts were performed by using secondary

electron (SE) mode. In order to increase the lateral res-

olution of the microstructural analysis, secondary elec-

tron (SE) mode was used. Although this technique isvery sensitive to topographic contrast, it can still give

atomic number contrast, and with a higher resolution

than with backscattered electrons. SEM-EDX was

merely used qualitatively, in order to confirm that the

Ti(C, N)/TiN coating was still intact after deformation.

To prepare SEM specimens, the inserts were cut,

using a Buehler Isomet high speed saw, perpendicular

to the cutting edge and through the deformed zone.The cross-section was then molded into a conductive

compound and ground on SiC paper to get a flat sur-

face. Finally, the surface was polished with 6, 3 and

1 lm diamond spray on paper.

2.3. Transmission electron microscopy

TEM analyses were performed in a Philips CM200FEG-TEM equipped with a Link Isis� EDX system.

G. Ostberg et al. / International Journal of Refractory Metals & Hard Materials 24 (2006) 135–144 137

For the most part, bright-field and dark-field imaging

was used to analyse different features but crystal struc-

tures and orientation relations between grains were

determined with selected area diffraction (SAD). TEM-

EDX was used to measure any segregation in grain

and phase boundaries.TEM specimens of undeformed materials were pre-

pared by cutting �200 lm slices from the inserts. B

3 mm discs were cut out, using a Gatan 601 ultrasonic

cutter, and then ground to a thickness of 100 lm. The

discs were dimple ground to 20 lm thickness and ion

polished to electron transparency in a Gatan precision

ion polishing system (PIPS), with 3–5 keV acceleration

voltage and incidence angles of 3–5�.The method used to prepare TEM specimens from

deformed materials involves focused ion beam (FIB)

milling and a procedure called in situ lift-out. In this

procedure a thin foil is picked out from the deformed re-

gion of the insert and attached to a copper grid before it

is thinned to electron transparency by ion polishing. A

more detailed description on this method can be found

elsewhere [13].

2.4. Three-point bending

Samples were spark cut to 35 · 7 · 3.5 mm3 in size

and the sides were planed to be parallel. A Schenk

RMC 100-1 mechanical testing machine was used for

testing at a constant strain rate of _e ¼ 1.5� 10�5 s�1 (de-

tailed description given in [14]). Carbide skeleton sam-ples were produced by etching the cobalt in a boiling

37% HCl + 5 g/l KClO3 water solution. Bending tests

were performed on complete and skeleton samples at

different temperatures between 750 and 1200 �C.

2.5. Mechanical spectroscopy

Mechanical spectroscopy measures the internal fric-tion (IF), or material damping, by applying a cyclic

stress with a subresonant inverted forced torsion pendu-

lum on bars, 35 · 4 · 1 mm3 in size. The IF was mea-

sured either at constant frequency as a function of

temperature between 300 and 1500 K or at constant

temperature as a function of frequency between 1 mHz

and 10 Hz. A more thorough description of the theoret-

ical formalism has been given by Mari et al. [14]. Basi-cally, the IF is the ratio of dissipated energy over

maximum stored elastic energy per cycle, measured as

phase lag between the applied stress and the strain re-

sponse. The dynamic modulus of the material is ob-

tained simultaneously as the ratio of the stress and the

strain amplitude. The IF spectra can be decomposed

into a number of relaxation peaks, each of which ac-

counts for a distinct relaxation process.In the particular case of the standard anelastic solid

[15], the IF can be expressed as a function of the angular

frequency x, the relaxation time s, and the relaxation

strength D:

IF ¼ tanu ¼ DðxsÞa

1þ ðxsÞ2a. ð1Þ

Such a function is called a Debye peak if a = 1, which is

broadened if a < 1 (1/a being the broadening factor).

The peak shows a maximum at xs = 1 with an ampli-

tude of D/2. The IF spectra can be composed of several

relaxation peaks, each of which accounts for a distinct

relaxation process. For thermally activated mechanisms,

the relaxation time depends on temperature following anArrhenius law:

s ¼ s0eEactkT ; ð2Þ

where Eact is the activation energy of the mechanism, k

Boltzmann�s constant and s0 a constant corresponding

to the inverse of the attempt frequency.

3. Ab initio calculations—method

In the computer simulations, the ðWCð0001Þ=WCð1�210ÞÞ asymmetric tilt boundary was used as a

model system for WC/WC grain boundaries. Calcula-

tions were performed for eight different translation

states in this boundary. In four of them, the close

packed (0001) interface atomic plane was metal termi-

nated, and in the other four it was carbon terminated.The structures were modelled in a supercell slab geome-

try with periodic boundary conditions. For all struc-

tures, density functional theory (DFT) as implemented

in the Vienna ab initio simulation package (VASP)

[16–19] code is used. The exchange-correlation func-

tional is approximated with the Perdew–Wang 1991

version of the generalized gradient approximation

(GGA-PW91) [20]. The plane-wave pseudopotentialmethod with Vanderbilt ultrasoft pseudopotentials

[21,22] was used. The plane wave cutoff energy was set

to at least 24 Ry (327 eV) in all calculations. Atomic

structure relaxations were performed with a quasi-New-

ton algorithm, and the structures were optimized until

the total energies were converged to at most 1 meV.

4. Results

4.1. Turning tests

A plot of the degree of deformation as a function of

cutting speed can be seen for both materials in Fig. 1. At

low and high cutting speeds there does not seem to be

any significant difference in deformation between thetwo materials, but at speeds around 450 m/min the cer-

met is more deformed than the WC–Co, and this differ-

ence appears significant.


4.2. Characterisation

A first, low magnification, study was performed on

the deformed cutting inserts with SEM in SE mode.

As can be seen in Fig. 2, the flank and rake face have

been depressed and by using EDX spot analysis at somepoints around the cutting edge it could be determined

that the coating was still intact since a clear Ti signal

in the spectrum was detected.

The SEM images in Fig. 3 show cross-sections of the

undeformed and deformed materials. In the undeformed

materials, the typical WC–Co and cermet microstruc-

tures can be seen. The WC–Co has a continuous hard

phase skeleton consisting of prismatic WC grains andthe cermet has rounded hard phase grains with a core/

rim structure.

When comparing the micrographs for the unde-

formed and deformed materials, it is evident that the

hard phase skeletons in both deformed materials are

partly broken up and the binder phase has formed

lamellae between the grains by infiltrating up to 10%

of the grain boundaries (marked with arrows in themicrographs). Since the resulting compressive force for

the deformation is directed diagonally from the upper

left corner across the images, as indicated with black ar-

rows in the figure, it can be seen that the Co lamellae are

oriented roughly parallel to this direction.

When looking at the structure in the TEM, the infil-

trated grain boundaries are seen very clearly in both

materials. Fig. 4 shows some typical lamellae, appear-ing as sharp lines, or bands, between the grains and

with a width varying between 20 and 50 nm. In bound-

aries perpendicular to the lamellae, GBS seems to have

taken place. Although no signs of bulk dislocation

movements or slip can be seen in the hard phase

Fig. 2. SEM micrograph of a WC–Co insert deformed at 500 m/min.

The white line indicates the original shape. The plastic deformation

(PD) is measured as the depression of the rake face.

grains, movements of dislocations related to the GBS

most likely have taken place in the grain boundary

regions.

The dark spots which can be seen in the hard phase

grains are due to radiation damage caused by ion beam

thinning of the specimen.In some of the infiltrated boundaries in the WC–Co,

a faceting of one of the grains can also be seen (Fig. 5).

Furthermore, some WC grains in the deformed WC–Co

exhibit signs of deformation (see Fig. 6). The grains have

a high amount of slip lines and steps have formed at the

surfaces.

Regarding the binder phase, no obvious differences

between the WC–Co and the cermet or undeformedand deformed materials could be seen in the TEM. A

rather high amount of stacking faults, which are bor-

dered by partial dislocations, is present in the binder

of both materials before, as well as after, deformation.

A TEM-EDX linescan of the undeformed WC–Co,

shown in Fig. 7, indicate that Co segregates to the grain

boundaries.

4.3. Internal friction and three-point bending

In Fig. 8, three-point bending measurements per-

formed at different temperatures are plotted for both

materials. The WC–Co deforms mainly elastically at

750 �C but at 800 �C the first sign of plastic deformation

can be seen. At 900 �C the material shows a ductile

behaviour and above 1000 �C it is highly plastic. Forthe cermet the deformation is elastic up to at least

1000 �C. At 1100 �C some ductility can be noticed and

at 1200 �C the material exhibits a highly plastic

behaviour.

The results from three-point bending of the complete

structures and the corresponding skeletons are com-

pared for the two materials in Fig. 9.

Considering that the skeleton samples lack the contri-bution of the cobalt phase to the overall deformation

resistance and that the true sample cross-section is smal-

ler because of the voids present instead of the cobalt, the

flow stress of the skeletons is expected to lie below that

of the complete hardmetal. This is true for low temper-

atures, but a critical temperature, where the behavior

changes, can be defined. Above that temperature, the

skeletons appear stronger than the complete hardmetals.Up to 750 �C the complete WC–Co is stronger than

the WC skeleton, which exhibits brittle fracture, but at

900 �C the complete material becomes weaker although

the skeleton is still quite brittle. Thus, the transition tem-

perature should be around 800 �C for the WC–Co and

near 1100 �C for the cermet.

The two skeletons also exhibit a considerable differ-

ence in plasticity. At 1000 �C there is a significant plasticdeformation of the WC skeleton whereas the cermet

skeleton is still brittle at 1100 �C.

Fig. 3. SEM images of the undeformed WC–Co (a) and cermet (c) with corresponding deformed materials, deformed at 475 m/min (b) and 500

m/min (d), respectively. Examples of formation of binder phase lamellae are marked with arrows. The somewhat uneven contrast of the binder phase

is due to topographic contrast.

Fig. 4. TEMmicrographs of binder lamellae, marked with arrows, in WC–Co deformed at 500 m/min (a) and the cermet deformed at 400 m/min (b).


Typical spectra of IF and modulus for both materials

as a function of temperature (at 1 Hz) have been pub-

lished previously [23]. Several IF peaks have been attrib-

uted to the movement of dislocations in the cobalt

phase, to the movement of dislocations in the carbonit-

ride phase and to grain boundary sliding. In this paper,

Fig. 5. TEM micrograph of one faceted and one straight binder

lamella between WC grains in WC–Co deformed at 500 m/min. The

planes parallel to the grain boundaries and the facets are indicated by

the (SAD) diffraction patterns.

WC/WC boundary in WC-Co

0

0.5

1

1.5

2

2.5

3

3.5

0 5 10 15 20

Atomic% Co K

Fig. 7. TEM-EDX linescan of a WC/WC boundary in the WC–Co.

Note that the background noise contributes to about 2 at% of Co.


we focus on one of these peaks appearing at high tem-perature, which has not yet been clearly interpreted. A

high temperature peak is found in both WC–Co and

the cermet. The shift of these peaks in the frequency

Fig. 6. TEM micrographs of plastically deformed WC grains in WC–Co de

family of planes and steps have formed at the grain surface (arrows). The g

spectrum due to thermal activation have also been ana-

lysed before [23]. A comparison of the peak in both

materials together with its modulus change is shown

for 1207 �C in Fig. 10. The peak is partially high and

broad. In the cermet it is present at even higher temper-

ature than that in WC–Co. Therefore, the maximum

cannot be observed in the limited temperature–fre-

quency range of the instrument. The peaks were absentin measurements performed on the hard phase skeletons

of both materials. Furthermore, it shows a strong sensi-

tivity on stress. An example of this effect is displayed in

Fig. 11. The increase of the excitation amplitude leads to

an increase of the high temperature peak in WC–Co,

present here merely as background. The low tempera-

ture peaks, PW2 and PW3, are not affected by the in-

creased stress. As in WC–Co, it was observed that thepeak in the cermet is sensitive to the excitation

amplitude.

The main IF parameters are summarized in Table 1.

Notice that the attempt frequencies of the peaks lie

formed at 500 m/min. In the grain in (a) slip has occurred along one

rain in (b) seems to have deformed along two plane directions.

Fig. 8. Three-point bending measurements at different temperatures

showing transitions from brittle to ductile and to highly plastic

behaviour for (a) the WC–Co and (b) the cermet.

Fig. 9. The transition between bulk deformation and Co enhanced

creep is experimentally shown by comparing the WC–Co (a) and

cermet (b) complete materials with their corresponding hard phase

skeletons.

Fig. 10. IF frequency spectrum and modulus change at 1207 �C of

WC–Co and TiWCN–Co showing the high temperature peak attrib-

uted to GBS by Co infiltration.


much above the Debye frequency (1013 Hz) and there-

fore the activation energy is apparent [14].

4.4. Ab initio calculations

The energy required to transfer a Co atom from the

binder phase to the studied WC/WC grain boundaries

has been calculated. An average value over all transla-

tion states is used. It is found that segregation occurs

to W-terminated WC/WC grain boundaries in sub-monolayer proportions. In the boundary, the Co atom

will replace either a carbon atom or a tungsten atom

in the ð1�210Þ interface layer.

The situation is somewhat different when considering

segregation to a carbon terminated boundary, where Co

atoms preferentially replace all available interface car-

bon atoms. In a generic boundary, which can be thought

of as consisting of regions varying in carbon and tung-sten content, a substitutional segregation of less than

one monolayer is expected.

Segregated Co will have a large effect on the interface

properties, e.g., the ability of the binder phase to pene-

trate the boundary. The driving force for boundary

infiltration, as given by the relation cWC/WC�2cCo/WC > 0, is large for clean boundaries (�1–3 J/m2

for W-terminated boundaries and �2–4 J/m2 for C-ter-

minated boundaries). However, the occurrence of segre-

gated Co in the boundaries substantially increases the

grain boundary resistance to metal infiltration (seeFig. 12). This effect is so large that for at least some

boundaries, no driving force for infiltration can be

expected.

Fig. 11. In frequency spectra, the strain amplitude influences mainly

the high temperature GBS peak (here appearing merely as back-

ground), whereas the other IF peaks are not affected.

Table 1

Characteristic parameters of the high temperature peaks found in WC–

Co and TiCN–WC–Co

Activation energy

(eV)

Attempt frequency

(s�1)

Broadening

factor

PW4 7 1025 0.3

PT4 5.3 1017 0.3

Effect of Co segregation on GB infiltration

0

0.5

1

1.5

2

2.5

3

3.5

0 1 2 4 5 6 8Grain boundary geometry

Incr

ease

in G

B in

filtr

atio

n re

sist

ance

(J/m

2)

Co replaces C

Co replaces W

3 7

Fig. 12. Effect of Co segregation on the ð1�210Þ=ð0001Þ R2 tilt grain

boundary strength for eight different geometries. For each geometry

the effect of replacing a C or W atom with Co in one of the adjacent

grains is given.

Change in Wsep (J/m2) for different GB structures

-1.00

-0.50

0.00

0.50

1.00

1.50

2.00

2.50

3.00

1 3

Grain boundary geometry

(1-210)-C

(0001)-(W,C)

(1-210)-W

2 4 5 6 7 8

Fig. 13. Change in work of separation, DWsep, for the eight different

grain boundary geometries. The strength is increased (DWsep positive)

for almost all grain boundary structures. Grain boundary geometry

1–4 corresponds to W-terminated grain boundaries and geometry 5–8

corresponds to C-terminated grain boundaries. For each geometry the

effect of replacing a C or W atom with Co in the interface plane in

accordance with the caption is given.


The presence of intergranular Co will also affect the

mechanical strength of the boundaries. The strength ishere taken to be the ideal work of separation, Wsep, cal-

culated as the difference between the sum of energies of

the free cleavage surfaces and the interface energy of the

intact boundary. This is an important quantity, also in

the context of boundary infiltration, as it should give a

measure of the energy barrier for grain boundary crack

initialization. Results for the work of separation are

shown in Fig. 13. It can be seen that the grain boundary

strength is increased significantly by segregated Co. The

largest effect is found for segregation by substitution ofcarbon in carbon terminated boundaries.

4.5. Discussion

The prediction of the ab initio calculations that Co

segregates to grain boundaries in submonolayer propor-

tions is confirmed by the TEM-EDX measurements. If

the Co concentrations in the grain boundaries are con-verted into atomic layers with the method presented

by Henjered et al. [24] they will correspond to a thick-

ness of less than a monolayer. This segregation increases

the work of separation and the resistance against Co

infiltration, according to the ab initio calculations, but

still we obviously get Co diffusion into the grain bound-

aries during deformation. The orientation of the binder

lamellae parallel to the resulting compressive force in theinserts suggests that the boundaries have been subjected

to a tensile stress perpendicular to the direction of

compression.

An explanation to this may be that the compressive

stress will force some grains to wedge themselves be-

tween other grains which, in turn, will be forced apart.

The local tensile stress arising in these boundaries will

then act as a driving force for Co diffusion and the en-ergy barrier, given by the work of separation, can be

overcome. Thus, it will be more energetically favourable

to form two phase boundaries instead of one grain

boundary. Simultaneously, boundaries oriented roughly

45� to the compressive force will experience shear stres-

ses. In these boundaries GBS will occur and will also be

facilitated by the Co infiltration of some grain

boundaries.


In the cermet, all lamellae appear as straight lines but

in the WC–Co some WC grains apparently continue to

lower their surface energy by forming facets against

the Co lamellae. Since WC grows preferentially along

well defined directions [25] the energy of the Co/WC

interface can be minimized by arranging the atoms inthese low energy planes. The straight lamellae are

formed in WC/WC interfaces already arranged along

low energy planes.

The temperatures in the turning process are unknown

but the three-point bending results show that the cemen-

ted carbide does not deform plastically to any significant

extent at temperatures below 800 �C and the cermet not

below 1100 �C. It can therefore be expected that theWC–Co and the cermet are exposed to temperatures

above 800 �C and 1100 �C, respectively, during turning.

In addition, the cermet most likely reaches a higher tem-

perature than the WC–Co at corresponding cutting

speed since it has a lower heat conductivity. The higher

temperature in the cermet could, thus, be an explanation

for the higher degree of deformation at the intermediate

cutting speeds. Hence, if the plastic deformation wasplotted against temperature, instead of cutting speed,

it would exhibit a more sluggish dependence for the cer-

met compared to the WC–Co over the whole range of

measurement.

A correlation between the mechanisms seen by the IF

measurements [14] and the three-point bending results

indicates that the deformation below 800 �C for the

WC–Co and below 1100 �C for the cermet can be asso-ciated with dislocation movements in the binder phase

and the hard phase skeleton. Dislocation movements

and even subgrain formation in the hard phase have also

been seen in creep tests [8,11,12].

The dislocation movements in the cermet hard phase,

envisaged from the interpretation of IF spectra, were

also found in the corresponding cutting inserts although

they were restricted to the grain boundary regions. Inthe WC skeleton, however, dislocation movements were

not detected by IF whereas the microstructure of the

cutting inserts most definitely exhibit signs of plastic

deformation of some WC grains.

These somewhat contradictory results may be ex-

plained by the difference in stress situation between the

test methods, giving totally different strain rates. The

strain rate in the cutting inserts during the turning testshas been estimated to the order of 10�3 s�1. In the IF

measurements the strain rate is dependent on the fre-

quency and varies between the order of 10�8 at 1 mHz

and 10�4 s�1 at 10 Hz which corresponds well with

strain rates reported from three-point bending and creep

tests [8,11,12,26].

Apparently, the high strain rate present during cut-

ting promotes dislocation movements in the hard phaseof the WC–Co and cause plastic deformation of the WC

grains. This deformation is most likely accommodated

by diffusion and climb in the binder phase which quite

readily occurs at the temperatures in question.

In the cermet, however, the hard phase dislocations

are restricted to the grain boundary regions and plastic

deformation is mainly occurring by GBS. The deforma-

tion rate in the cermet is therefore controlled by move-ment of grain boundary dislocations and the mobility

and redistribution of the binder phase. This is also in

accordance with the creep deformation model presented

by Fahrmann [11] who assumed that plastic deforma-

tion of cermets occurs by rigid carbide grains sliding

against each other via thin binder films.

Above 900 �C, the complete WC–Co material be-

comes weaker than the WC skeleton, indicating that acertain amount of GBS, accommodated by binder phase

deformation, occurs. The presence of binder phase

lamellae in the grain boundaries implies that the binder

has deformed not only by dislocation movements but

also by diffusion which can be expected to occur simul-

taneously with the GBS, since this process breaks up the

skeleton and thereby facilitates GBS. Other deformation

mechanisms of the binder phase, such as fcc- >hcpphase transformation, have also been reported [27,28]

but no unambiguous signs of this could be seen by

TEM in the cutting inserts. Around 1000 �C, the WC–

Co starts to exhibit a highly plastic behaviour together

with a significant deformation of the skeleton. This sug-

gests that deformation by GBS is accommodated by the

WC grain deformation.

For the cermet, a plastic behaviour can be seen fromaround 1200 �C while the skeleton is still brittle at this

temperature, which indicates that the deformation by

GBS is not accommodated by bulk deformation.

The high temperature IF peaks shown in this study

are attributed to GBS of the carbide/carbonitride phase

infiltrated by the cobalt. In fact, they are only found in

the complete hardmetals and not in the ceramic skele-

tons. The presence of these peaks shows that, despite adifference in temperature (possibly related with the melt-

ing point of the Co/hardphase eutectic), the high tem-

perature mechanisms are quite comparable in both

materials. The peaks are characterized by an apparent

activation energy, which is related to a microstructure

that changes within the course of the measurements.

As observed by deformation tests, cobalt enhances the

deformation of the carbide skeleton at high temperature.Cobalt infiltration of the grain boundaries must be dri-

ven by stress. In fact, cobalt lamellae are only found

in deformed specimens. A characteristic of the high tem-

perature peaks reported in this study is that they are

stress sensitive, i.e., their amplitude increases as the IF

excitation amplitude is increased (Fig. 11). This is not

the case for the lower temperature peaks reported in

Fig. 11 [29,30]. Even with the small stresses applied dur-ing internal friction experiments, some microplasticity

may be induced in the grain boundaries leading to


microstructural changes and, thus, to the high apparent

activation energies.

5. Conclusions

• Ab initio calculations predict that Co segregate to

grain boundaries and thereby increases the work of

separation and the resistance to Co infiltration.

• Segregations of Co in submonolayer proportions in

the grain boundaries are confirmed by TEM-EDX.

• During cutting, tensile stresses are induced in the

grain boundaries parallel to the cutting force, creating

a driving force for Co diffusion into the boundaries.• The hard phase skeletons of both the WC–Co and the

cermet are partly broken up by Co grain boundary

infiltration, making GBS possible.

• Three-point bending shows that the WC–Co and the

cermet deform plastically at temperatures above

800 �C and 1100 �C respectively.

• The WC–Co has a gradual transition from brittle to

plastic behaviour. Deformation starts with disloca-tion movements in the binder phase and moves over

to a gradual increase in GBS, accommodated by Co

diffusion and dislocation climb in combination with

dislocation movements and slip in the hard phase.

• The cermet exhibits a more abrupt transition from

brittle to ductile behaviour where dislocation move-

ments in the binder phase precede the GBS accommo-

dated by Co diffusion and dislocation movements inthe hard phase grain boundary regions.

References

[1] Lin J, Liu C-Y. Measurement of cutting tool temperature by

an infrared pyrometer. Meas Sci Technol 2001;12:1243–9.

[2] Yen Y-C, Sohner J, Lilly B, Altan T. Estimation of tool wear in

orthogonal cutting using the finite element analysis. J Mater

Process Technol 2004;146:82–91.

[3] Bolognini S, Feusier G, Mari D, Viatte T, Benoit W. High

temperature mechanical behaviour of Ti(C, N)–Mo–Co cermets.

Int J Refract Met Hard Mater 1998;16:257–68.

[4] Feusier G, Cutard T, Verdon C, Viatte T, Benoit W. High

temperature properties of TiCN–Mo2C–Co cermets studied by

mechanical spectroscopy. J Phys IV 1996;6:751–4.

[5] Viatte T, Cutard T, Feusier G, Benoit W. High tempera-

ture mechanical properties of Ti(C, N)–Mo2C–Ni cermets stud-

ied by internal friction measurements. J Phys IV 1996;6:743–6.

[6] Sakuma T. High temperature plastic flow in cemented carbides

and cermets. Key Eng Mater 1995;108–110:435–48.

[7] Schmid HG, Mari D, Benoit W, Bonjour C. The mechanical

behaviour of cemented carbides at high temperatures. Mater Sci

Eng A 1988;105/106:343–51.

[8] Lay S, Vicens J, Osterstock F. High temperature creep of WC–Co

alloys. J Mater Sci 1987;22:1310–22.

[9] Sakuma T, Hondo H. Plastic flow in WC–13 wt.%Co at high

temperatures. Mater Sci Eng A 1992;156:125–30.

[10] Cutard T, Feusier G, Viatte T, Benoit W. Microstructure and

high temperature mechanical properties of TiC0.7N0.3–Mo2C–Ni

cermets. Mater Sci Eng A 1996;209:218–27.

[11] Fahrmann M. High temperature creep of Ti(C, N)-based cermets.

Int J Refract Met Hard Mater 1989;8:219–23.

[12] Wirmark G, Dunlop GL, Chatfield C. Tensile creep of WC–Co

cemented carbides at 800–900 �C. Int J Refract Met Hard Mater

1986;5:153–7.

[13] Ostberg G, Andren H-O. TEM specimen preparation of plasti-

cally deformed WC–Co cutting inserts using focused ion beam

(FIB), submitted for publication.

[14] Mari D, Bolognini S, Feusier G, Viatte T, Benoit W. Experimen-

tal strategy to study the mechanical behaviour of hardmetals

for cutting tools. Int J Refract Met Hard Mater 1999;17:

209–25.

[15] Nowick AS, Berry BS. Anelastic relaxation in crystalline sol-

ids. New York: Academic Press; 1972.

[16] Kresse G, Furthmuller J. Phys Rev B 1996;54:11169.

[17] Kresse G, Hafner J. Phys Rev B 1993;47:558.

[18] Kresse G, Hafner J. Phys Rev B 1994;49:14251.

[19] Kresse G, Furthmuller J. Comput Mater Sci 1996;6:15.

[20] Perdew JP, Chevary JA, Vosko SH, Jackson KA, Pederson MR,

Singh DJ. Phys Rev B 1992;46:6671.

[21] Kresse G, Hafner J. J Phys: Condens Matter 1994;6:8245.

[22] Vanderbilt D. Phys Rev B 1990;41:7892.

[23] Buss K, Mari D. High temperature deformation mechanisms in

cemented carbides and cermets studied by mechanical spectros-

copy. Mater Sci Eng A 2004;370:163–7.

[24] Henjered A, Hellsing M, Andren H-O, Norden H. Quantitative

microanalysis of carbide/carbide interfaces in WC–Co-base

cemented carbides. Mater Sci Technol 1986;2:847–55.

[25] Exner HE. Physical and chemical nature of cemented carbides. Int

Met Rev 1979;243:149–73.

[26] Bolognini S, Feusier G, Mari D, Viatte T, Benoit W. TiMoCN-

based cermets: high-temperature deformation. Int J Refract Met

Hard Mater 2003;21:19–29.

[27] Erling G, Kurosawe S, Luyckx S, Sockel HG. Stable and unstable

fracture surface features in WC–Co. J Mater Sci Lett 2000;19:

437–8.

[28] Vasel CH, Krawitz AD, Drake EF, Kenik EA. Binder deforma-

tion in WC–(Co, Ni) cemented carbide composites. Metall Trans

A 1985;16A:2309–17.

[29] Buss K. High temperature deformation mechanisms of cemented

carbides and cermets, Lausanne, 2004.

[30] Buss K, Mari D. In: Proceedings of EuroPM 2002, Lausanne,

Switzerland, 2002.

Documents

Mechanisms of plastic deformation of WC–Co and Ti(C, N)–WC–Co