Internal friction in WC-Co hard metals

Materials Science and Engineering, A 105/106 (1988) 313-321

313

Internal Friction in WC-Co Hard Metals* R. SCHA LLER and J. J. A M M A N N

lnstitut de G~nie A tornique, Ecole Polytechnique Fkdkrale de Lausanne, PHB-Ecublens, CH-I OI 5 Lausanne (Switzerland) C. BONJOUR

Stellram SA, CH-1260 Nyon (Switzerland) (Received November 9, 1987)

Abstract

Internal friction measurements have been performed on various grades ,of WC-Co cemented carbides from room temperature to 1000 °C. These composite materials exhibit an internal friction spectrum which is mainly composed of a relaxation peak and a high temperature exponential background. The peak appears in the same temperature range where an increase in toughness has been observed and interpreted as being due to a brittle-to-ductile transition of the material The exponential background can be associated with high temperature creep phenomena. The analysis of the results obtained shows the important role of the cobalt binder phase in the mechanical behaviour of WC-Co. 1. Introduction W C - C o cemented carbides, used as cutting tools, are subjected to severe stress conditions at relatively high temperatures (800-1000°C). During machining, the tool lifetime can be limited by various wear mechanisms such as abrasion, delamination, chemical dissociation and dissolution [1 ], but also by deficiency in the mechanical properties such as toughness or creep resistance [2]. It is important to obtain a better understanding of the fundamental mechanisms which are responsible for the mechanical properties, such as toughness, of these composite materials. For simplicity, W C - C o can be considered as a two-phase material: WC hard particles and cobalt *Paper presented at the 3rd International Conference on the Science of Hard Materials, Nassau, The Bahamas, November 9-13, 1987.

0921-5093/88/$3.50

binder phase. It is evident that the mechanical properties of such a composite will strongly depend on the volumetric fraction and the morphology of the two phases [3, 4]. They are also greatly affected by structural defects such as dislocations, twin boundaries and point defects, which can play an important role in the plasticity of each phase [5, 6] or in grain boundary sliding

[7]. Internal friction is a parameter which is very sensitive to the mobility of structural defects [8]. It results from material anelasticity. Anelastic deformation is due to the reversible movements of structural defects, as dislocations, around their equilibrium position. Such movements are generated by small applied stresses which do not give rise to any modification of the microstructure. Then, because of the very low strain amplitude investigated ( 10 - 8_ 10 - 4), internal friction can be considered as a non-destructive technique well suited to the study of the mechanisms of microplastic deformation which are responsible for the mechanical properties of hard and brittle materials exhibiting a limited plasticity. In this sense, internal friction has been used to study the microscopic mechanisms which may be responsible for the mechanical behaviour of W C - C o cemented carbides from room temperature to 1300 K. The non-destructive character of the technique has revealed itself as very advantageous for measuring brittle specimens in a reproductive manner under various experimental conditions. 2. Internal friction technique Anelasticity of solids [8, 9] occurs in the following simple experiment (Fig. l(a)). A stress o, of low intensity, is applied abruptly to a specimen © Elsevier Sequoia/Printed in The Netherlands

314 applied stress, is called anelastic relaxation and may be defined by two parameters: the relaxation intensity

O

A=

6a ~° _ _ J r - J u

ee

.t

and the relaxation time r. The total strain e can be written as a function of time:

v

11.

'Eo

If the stress is removed after a certain delay time t, we can observe the instantaneous recovery of the elastic strain and with the same relaxation time r the recovery of the anelastic strain. In such an experiment, strain is completely recoverable. This behaviour is representative of the standard anelastic solid, the equation of which can be derived from eqn. ( 1):

t T

//////////////

I

=J o+ rJoo T JR- Ju

A

Ju

Ju

m W

Q

I,i

O" Fig. 1. Anelasticity of a solid: (a) elastic and anelastic strain e of a solid submitted to stress a; (b) rheological model corresponding to the anelastic behaviour in (a).

at time t = 0 and held constant while the strain e is recorded as a function of time. We observe the instantaneous elastic strain ee = Jua (where Ju is the unrelaxed compliance) and the anelastic strain ea which increases with time from zero to an equilibrium value ea °°. When equilibrium is reached, = ~e "4- Ea °°

=Jr a where Jr is the relaxed compliance. This evolution of the solid from one equilibrium state to a new state, under a mechanical

(2)

From a rheological point of view, the solid can be represented by the model shown in Fig. l(b). The spring B in parallel with the dashpot C is essential to the recoverable nature of the anelastic strain. If this spring does not operate, creep behaviour is observed; anelastic strain does not reach an equilibrium value and is then no longer completely recoverable. From a microscopic point of view, anelastic strain can be interpreted as being due to the movements of structural defects (elastic dipoles or dislocations) from one equilibrium position defined at a = 0 to another defined at a ~ 0. The relaxation intensity A is then proportional to the concentration of defects which are relaxing, the relaxation time accounting for their mobility. Measurements performed in the way shown in Fig. 1 are very delicate. Effectively, if the elastic strain amplitude is about 10 -6, anelastic strain values of 10- 9_ 10- 8 must be obtained in order to detect anelastic relaxation phenomena in metals. For this reason, it is more convenient to use dynamic methods in practice for measuring the relaxation parameters. In this case, an alternative stress a = ao exp(ia~t) is applied to the system. The linearity of the stress-strain relationship ensures that strain is periodic with the same frequency: e = e0 exp{i(tot - 6)}

315

where 6 is the phase lag of strain behind stress due to anelasticity. Introducing these expressions for o and e into eqn. (2) leads to the following relationship:

= {Jl( 0) ) - iJ2( 0) )} o

(3)

where Jl and -/2 respectively, are the real and the imaginary part of the complex compliance J*:

L-L J1 = J u -f

2 "

1+0) r-

0)T

4 = (Jr- L)

2 2 1+0) r

(4)

The internal friction Q - I of the material is related to the loss angle d by Q - l = tan 6 =_J2

(5)

JL

Introducing eqn. (4) into eqn. (5) and assuming that Jr -- Ju '~ Ju gives Q t= A

0)r 1+0) 2 r 2

dJ J,(0))-Jo J -A

1 1+

0) 2 1"2

where dJ/J is the variation of the dynamic modulus due to the anelastic relaxation. Equations (6) show that the internal friction Q - 1 has a maximum as a function of 0)r, centred at 0)r = 1. This internal friction peak gives us the relaxation intensity A (the height of the peak) and the relaxation time r (the position of the peak on the 0)r axis). If several independent relaxation mechanisms are activated, the superposition principle allows us to write

Q-I= Z I

0)1"i

Ai 1 + 0)2 ri2

Relaxation i gives rise to a peak located at temperature Tpi (defined by 0)roexp(EJkTpi)=l, where Ei is the activation energy), and the height of which is ½Ai. Therefore, we obtain an internal friction spectrum as a function of the temperature, which is characteristic of the material.

3. Experimental procedure

For metals, tan d (see eqn. (5)) is small and therefore it is generally difficult to measure the phase lag between strain and stress directly. In the low frequency range, it is easier to use resonant systems which vibrate at a natural frequency. The internal friction, Q-~ is deduced from the free decay of the vibrations. When tan 6 is small, it is possible to show that Q L=tan d

(6)

Ju

depends on the temperature), r = roexp(E/kT), and it is possible to measure the internal friction as a function of temperature, keeping 0) constant. Then,

(7)

where i refers to the mechanism. The material then exhibits an internal friction spectrum composed of internal friction peaks due to the various microscopic relaxation mechanisms. When the relaxation mechanisms are thermally activated (the viscosity of the dashpot in Fig. l(b)

ln( )

,9,

where A / a n d Ai+,, , respectively, are the ith and the (i + n)th amplitudes in the free decay of the vibrations. In the present case, the internal friction measurements were performed by means of an inverted torsion pendulum working in the low frequency range (0.2-5 Hz) under vacuum (10-3 Pa). The specimens used for the measurements were sheets of size 0.5 ram× 2.5 ram× 110 mm. They were cut by spark machining into bars of size 2.5 m m × 2 0 m m × 1 1 0 mm. These initial bars of WC-Co, with varying amounts of cobalt and an initial carbide grain size of 3.6/~m, were then sintered. The measurements of the internal friction and the dynamic elastic modulus were performed automatically during heating and cooling at a constant rate of 2 K mm-1. The temperature range investigated was between room temperature and 1300 K. The maximum strain amplitude was approximately 5 × 10- 6.

316

4. Results

WC-Co cemented carbides exhibit a characteristic internal friction spectrum as a function of temperature (Fig. 2) [10]. It is mainly composed of a peak located at about 950 K in WC-11wt.%Co (vibration frequency, about 2.6 Hz). In the temperature range lower than the peak, i.e. from room temperature to 700 K, the internal friction is weak but constant. In the temperature range above the peak, the internal friction is high (approximately 15 x 10-3 at about 1100 K) and increases exponentially with the temperature. This part of the internal friction spectrum will be referred to as "the exponential background". Associated with the peak, as well as with the exponential background, we observe anomalies in the decrease in the vibrational frequency as a function of temperature. These anomalies are due to the effects of the anelastic relaxations on the dynamic modulus (eqn. (6)). The internal friction peak is effectively a relaxation peak, because it shifts in temperature when the frequency is modified. The activation energy, deduced from the Arrhenius plot, is 2.7 eV, which is very close to the self-diffusion energy in cobalt. In contrast, the internal friction (mainly the exponential background component) increases with increasing cobalt content (Fig. 3). In addition, the relaxation peak shifts towards lower temperatures. These results tend to show that the internal friction phenomena, peak and exponential background, are associated with mechanisms which take place in the cobalt binder phase. In order to localize more precisely the relaxation mechanisms, specimens of pure WC and

pure cobalt were prepared. WC samples were sintered without binder at 2000 °C. Cobalt samples were obtained by melting cobalt industrial powders in an alumina crucible. The internal friction results obtained from these specimens are compared to the spectrum for WC-11wt.%Co in Fig. 4. It clearly appears that the relaxation peak in WC-Co (curve a) is not due to a microscopic mechanism in the WC particles or at the WC-WC interfaces. The internal friction of pure sintered WC presents only a slight monotonic increase in this temperature range (curve b). Also, the peak is not due to the allotropic phase transition of cobalt which is responsible for the sharp peak located at about 700 K [11] (curve c). Pure cobalt exhibits a similar increase in the internal friction at high temperatures as WC-Co, but no significant peak appears in the temperature range at around 900 K, although the change in slope at about 1100 K is

20.0

/ .-23%

15.0

Co

/

....~/ 11°/o Co

10.0 •" ./'~ :" / / :" / ,7 • /" "

5.0

0.0

~

5 % Co

,

500

G00

~T

700

8a00

(K) I

9~00 10~00 1100 1200

Fig. 3. Effects of the cobalt content on the internal friction of W C - C o cemented carbides.

30,0

O-~* 103

F(Hz) 2.80

-t*103

30.0

Q-t. 10 a

25 . 0

a)

W [ - 11°/'o [ o

2.75 25.0 2.70 2.G5

20.0

" ' c)

15.0 ...../

2.G0 10. 0

//"/ i "

10.0

0.0

~ 500

-/ G00

' 700

800

,.

,," ""/----~-(~~v- ~ ' ~

2.55

""" \,

2.50

/ //

5.0

Eo

2.45 i 900

i 1000

i 1100

, T (K 1200

2,40 1300

Fig. 2. Internal friction Q - i and oscillation frequency F (giving approximately the dynamic elastic modulus) of WC- 11 wt.%Co as a function of temperature.

/

o.~0 °

__~_--~G00

700

_____-/ 800

900

1000

1100

T cK~ 1200

1300

Fig. 4. Internal friction spectra of WC-1 lwt.%Co (curve a), pure WC sintered without binder (curve b) and a cobalt specimen obtained by melting cobalt industrial powders (curve c).

317

certainly due to the presence of a tiny peak located at about 1000 K. From the results shown in Figs. 3 and 4, it is possible to conclude that the relaxation peak and the high temperature exponential background in WC-Co are due to the cobalt binder phase, but the relaxation mechanism is not related to the properties of pure bulk cobalt. The cobalt binder phase in WC-Co exhibits mechanical properties completely different from those of pure bulk cobalt. This also means that the behaviour of the cobalt phase is affected by the presence of the surrounding WC skeleton. 4.1. Thermal treatments The internal friction spectrum of WC-Co is sensitively modified by heat treatment of the specimen (Fig. 5). In the just-sintered condition (curve 1), the internal friction peak is located at a relatively high temperature (about 900 K) and the exponential background is relatively high. After it has been annealed for 12 h at 1000 K, a decrease in the high temperature exponential background and a shift of the peak towards a lower temperature (about 800 K) are observed (curve 2). During subsequent aging at 1000 K, the peak temperature remains constant, but the peak height decreases somewhat. After 40 h at 1000 K, it reaches the stable state shown in curve 3. If the temperature is increased above 1000 K, a rapid increase in the internal friction with increasing temperature is observed (curve 3). After the specimen has been annealed at 1200 K for 1 h, curve 4 is obtained; this is very similar to curve 1. Annealing at 1200 K leads to a microstructural

20.0

~_ C~-1

"

103

WC - 11%

//

Co

.//

1. A f t e r sintering 2 A f t e r 12hrs at IO00K 3

Atfer

/

///1

l+0hrs a, ,000K

t+. A f t e r lhr a, 1200K

4.2. Changes in the binder The cobalt binder phase can be modified by the addition of some alloying elements. It is known that additions of ruthenium or chromium to WC-Co cemented carbides can reduce grain coarsening during sintering [13], but they can also modify the mechanical properties of the cobalt binder phase. In Fig. 6, the effects of chromium or ruthenium additions on the internal friction peak are presented. Each specimen was measured after the same heat treatment (after an anneal for 12 h at 1000 K). Curve a was obtained for the standard WC-1 lwt.%Co; curves b and c

Q-~* 1 0 a

26.6

,

/ /

/~.~-

state similar to the sintered condition. The transition from curve 2 to curve 4 is reversible. Another anneal at 1000 K of the specimen exhibiting curve 4 leads to the reappearance of curve 2. Such a behaviour is very close to that associated with precipitation and resolution processes. Effectively, other researchers [12] have detected, by high temperature X-ray analysis, the formation of C%W precipitates during aging at about 700 °C (973 K). Therefore, 1000 K anneals can activate the precipitation of tungsten atoms in a supersaturated Co-W solid solution, and 1200 K anneals lead to a resolution process because of an increase in the solubility. If the solute atoms are controlling the mobility of defects, which are responsible for the anelastic relaxation in cobalt, it is easy to understand that changes in concentration of the solid solution by precipitation lead to shifts of the internal friction peak along the temperature axis during heat treatments where the relaxation time ~ is modified.

15.0

a) WC- 11% Co b) WE- 9.5% C o - 1.5% c) W C - 9 . 5 % C o - 1 . 5 %

Ru Cr

10.0 / ; 7

-

-

f ~.~ _-~:

10.8

3

./" a

5.0 0.0

~

L

600

L

800

i

I

1000

i

.:

..

/

T (K)

..-

im

,,.,,, ..J"""

" ."

¢ //

..

1200

Fig. 5. Evolution of the internal friction spectrum of W C - 1 lwt.%Co during subsequent heat treatments: curve 1, after sintering; curve 2, after annealing for 12 h at 1000 K; curve 3, after annealing for 40 h at 1000 K; curve 4, after annealing for 1 h at 1200 K.

566

666

766

06~

s'~6

l~'e6

t1~

12~

Fig. 6. Internal friction as a function o f temperature, for W C - ] ] w t . % C o (curve a), W C - I ] w t . % C o - R u (curve b) and

W C - I I w t . % C o - C r (curve c).

318

were obtained for specimens containing 1.5 wt.% Ru and 1.5 wt.% Cr respectively. With respect to the internal friction of pure WC-11wt.%Co, the additions of ruthenium and chromium shift the curves towards higher temperatures. Moreover, a decrease in the peak height is observed. For W C - C o - C r (curve c) the internal friction above the peak temperature is reduced by a factor of 2 in comparison with that for pure WC-Co. In addition, a large decrease in the peak height has been observed for the W C - C o - C r system during anneals at 1000 K [10], accompanied by an increase in the elastic modulus. This can be a characteristic behaviour of precipitation hardening [ 14]. Even if the internal friction spectrum is rather sensitive to modifications of the cobalt binder phase, it does not change drastically when the nature of the binder phase is completely different (Fig. 7). Two other types of binder were tested. Figure 7 shows the internal friction spectra of WC-11wt.%Fe and WC-11wt.%Ni measured under the same experimental conditions as WC-11wt.%Co. These spectra are qualitatively similar. They are composed of a relaxation peak and a high temperature exponential background. The main differences are the position of the peak along the temperature axis and the intensity of the high temperature internal friction. When the peak temperature is lower, the high temperature exponential background is higher for a given temperature. It seems that the relaxation mechanism which gives rise to the peak is responsible for the rate of increase in the internal friction at high temperature. In Fig. 7, it is possible to note that the peak which appears for WC-Fe is 200 K O.-1 " 10 3

//we-N,

30.0

wc-~ , / //

/ /

///

20.0

/

//

/ ~// i

/~WC-Co

/ /

10.0

0.0

.

600

,

800

j

1000

,

tT

(K)

1200

Fig. 7. Internal friction spectra of WC-11wt.%Fe WC- 11wt.%Ni compared with the spectrum WC- 11 wt.%Co.

and for

lower than for WC-Co. In the case of WC-Ni, the shift towards lower temperature is about 130 K. The internal friction spectrum of WC-Fe exhibits a second peak at about 980 K. This peak is not a relaxation peak but a phase transition peak due to the allotropic transformation a-Fe to 7-Fe. The transition temperature T~appears to be lower than in pure bulk iron (T~= 910 °C). This could be due to the presence of carbon in the iron binder phase and also to the stresses created by the surrounding WC particles. Nevertheless, the detection of the peak at about 980 K assures us that the internal friction exponential background originates mainly in the binder phase.

5. Discussion

WC-Co cemented carbides exhibit a characteristic internal friction spectrum mainly composed of two components: the relaxation peak and the high temperature exponential background. These two components appear only when the WC skeleton is immersed in a ductile binder phase. From a rheological point of view, a relaxation peak appears as the solution of a three-parameter model containing a Voigt unit (Fig. l(b)) [8]. In such a model, the spring A of compliance Ju would be mostly due to the elasticity of the WC skeleton because the elastic modulus of WC-Co is much greater than the modulus of cobalt. However, the dashpot C is certainly relevant to the binder phase because the damping of WC, sintered without a binder, is weak (Fig. 4). Spring B of compliance J r - J , is essential to the recoverable character of the anelastic deformation, and then to the appearance of a peak in the internal friction curve. If spring B does not work when dashpot C is activated, the anelastic deformation does not reach an equilibrium value. Creep behaviour appears. Then the internal friction does not reach an equilibrium value and an exponential increase as a function of temperature is observed. The restoring force associated with spring B, which limits the anelastic strain, could be due to the WC skeleton which limits the deformation of the binder or to pinning points in the binder which limit the movements of structural defects responsible for the relaxation. It is interesting to compare the internal friction results with the mechanical properties of WC-Co cemented carbides measured by classical tech-

319

niques. First, we can observe that the internal friction peak is located in the same temperature range in which an increase in the critical stress intensity factor has been observed [15, 16]. Effectively, in WC-6wt.%Co, K k starts to increase at about 700 °C (about 973 K) and the relaxation peak in WC-5wt.%Co is centred at around 900 K (Fig. 3). Others [15, 16] have interpreted this increase in toughness as being due to a brittle-to-ductile transition of the material. It is therefore possible to admit that the peak is associated with such a transition and may account for the increase in toughness due to a mechanism controlled by the cobalt binder phase. The increase in toughness results from the activation of the dashpot C (Fig. l(b)) which dissipates a part of the energy used for crack propagation. This dissipation can be achieved by movements of structural defects, such as dislocations, under the effect of the applied stress. It is known that two deformation paths may occur in the cobalt binder phase [6, 17]. At room temperature, the crystal structure of the cobalt binder phase is predominantly f.c.c, with a certain amount of h.c.p, faults. Effectively, the normal allotropic transition from f.c.c, to h.c.p, of pure bulk cobalt does not take place in WC-Co. This is confirmed by our results (Fig. 4) and can be interpreted as being due to the effects of carbide stresses and solutes acquired during sintering. Deformation is then achieved mostly by a stressinduced f.c.c.-to-h.c.p, martensitic phase transformation. As a consequence, the number of slip systems decreases during deformation, leading to macroscopic brittle behaviour. However, in this case, toughness is higher than in pure WC because of the binder ligament rupture mechanism behind the crack tip [18, 19]. As nickel is added to the cobalt, the stability of the f.c.c, structure at room temperature increases, and Vasel et al. [6] have shown that dislocation activity and twinning become the prevalent modes of deformation. However, also, the f.c.c. crystal structure of pure cobalt is stable at high temperatures. This suggests that the binder deformation can be activated by dislocation movements in the f.c.c, phase at high temperatures, giving rise to an increase in the ability of cobalt to relax strain in the composite material. As a consequence, toughness increases. In this way, the internal friction peak may account for the appearance of the second mode of deformation of the binder, i.e. by dislocation movements, and conse-

quently for the increase in toughness as being due to an increase in the binder phase plasticity which may play an important role in front of the crack tip [20]. Above the peak temperature, the internal friction increases exponentially as a function of the temperature without reaching an equilibrium value. Such internal friction behaviour is normally observed when the restoring force (spring B in Fig. l(b)) does not work, leading to creep of the material. Effectively non-negligible creep associated with pore formation along the WC-Co and WC-WC interfaces has been observed above 1100 K by Schmid et al. [21] in the same WC-1 lwt.%Co measured by internal friction. Even if the exact nature of the internal friction mechanisms is not yet fully determined, it is possible to hypothesize that damping is due to dislocation movements in the binder phase. The mobility of these defects is limited by a viscous force which may be due to various interaction mechanisms with the solute atoms (pinning, depinnning and dragging) [9]. In fact, the diffusion energy of tungsten in cobalt is 2.95 eV atm-1 [22], which is close to the self-diffusion energy of cobalt (2.8 eV) or to the activation energy of the internal friction peak (2.7 eV). However, a peak in internal friction requires a restoring force opposed to the applied stress, which is essential to the recoverable nature of the relaxation (spring B in Fig. l(b)). Here the restoring force could be due to the WC skeleton which prevents plastic deformation of the binder or to pinning points which limit dislocation movements. Pinning points on the dislocation loops can be precipitates in the binder or the WC-Co interfaces. When the pinning points are strong enough, the dislocations can move only around an equilibrium position and a peak is observed. When the pinning points become weaker, the dislocations can escape and creep is initiated. The weakening of the dislocation pinning points may be responsible for the exponential increase in the internal friction above the peak temperature. If the pinning points are WC-Co interfaces, the weakening may be due to grain boundary sliding. At high temperatures the amplitude of grain boundary sliding can be large enough to lead to cavitation {2O[. Internal friction is good for the mechanical properties, because damping can increase toughness (for a discussion of toughening by martensitic transformation see ref. 23), but the

320

exponential increase means that creep occurs. This behaviour is undesirable in cutting tools. It is possible to reduce the high temperature exponential background, therefore reducing creep, by introducing hard pinning points (i.e. precipitates) on the dislocations. Our results (Fig. 5) show that heat treatments at about 1000 K, which can activate the formation of Co3W precipitates, lead to a decrease in the high temperature exponential background. The shift of the peak towards lower temperatures due to such an annealing can be explained by a decrease in the solute atom concentration. These solute atoms (tungsten) leave the dislocations for the precipitates. As a consequence, the dragging force decreases and the peak appears at lower temperature. Annealing at 1200 K gives rise to the resolution of tungsten in the binder phase. The dragging force then increases and at high temperatures the restoring force decreases. The dragging force depends on the nature of the binder. This may be the reason why the peak shifts with temperature when cobalt is replaced by nickel or iron (Fig. 7). For a rough approximation, it is possible to assume that the high temperature dislocation mobility is inversely proportional to the self-diffusion energy. In this sense, the position of the peaks along the temperature axis is understandable for the cases of WC-Ni and WC-Fe. Self-diffusion energy in nickel is higher than in iron, and consequently the relaxation peak in WC-Ni appears at a higher temperature than in WC-Fe. If only the diffusion energies are considered, the peak in WC-Co would appear between the peaks observed in WC-Fe and WC-Ni. However, cobalt is a special case since, contrary to nickel, the fault energy in the cobalt f.c.c, structure is low at low temperatures. Then the dislocations can sprit into partials, limiting an h.c.p, fault. This means that the dislocation cores are extended and they can be decorated by tungsten atoms in a rather high concentration. (It seems that h.c.p. cobalt holds more W C i n solution than f.c.c. cobalt does [24].) Then the dragging force which acts on dislocations is higher in the cobalt binder phase than in the nickel or iron phases. Also, in f.c.c, cobalt, dislocations can move only when the h.c.p, structure is no longer stable. For instance, ruthenium stabilizes the h.c.p, structure. Consequently, the peak in W C - C o - R u is shifted towards higher temperatures (Fig. 6). Chromium additions lead to a similar behaviour,

but the high temperature exponential background is also weak in this case (Fig. 6(c)). This must result in a better creep resistance of the material.

6. Conclusions WC-Co cemented carbides exhibit a characteristic internal friction spectrum mainly composed of a relaxation peak and a high temperature exponential background. The peak is located in the same temperature range in which an increase in toughness has been observed by other workers [15, 16] and is associated with a brittle-to-ductile transition of the material. The analysis of the internal friction peak allows us to attribute the origin of this increase in toughness to an increase in the damping capacity of the binder phase. In contrast, creep of the material seems to be connected with the high temperature exponential background. When the exponential background is high, the creep resistance is poor. It is possible to reduce the exponential background by introducing hard pinning points, such as precipitates, into the binder. The influence of creep would then be less.

Acknowledgments This work was financially supported by the Swiss Commission pour l'Encouragement de la Recherche Scientifique.

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