Fibre Science and Technology 19 (1983) 37-58
Fracture Toughness (Kic) of Cemented Carbides
J. D. B o l t o n a n d R. J. Keely National Institute for Higher Education, Limerick (Republic of Ireland)
SUMMARY In this paper, a brief review is presented of fracture toughness (Kit) results and Kic test methods for cemented carbides with emphasis on precracking methods, specimen geometry and on how fracture toughness testing of these materials can be rationalised to make it more acceptable for industrial use. Specific emphasis is placed on the advantages of Kic testing as compared to conventional strength test measurements and on the close relationships that exist between Ki~ and microstructure, composition and strength of these materials. Also discussed in this paper are fracture toughness results estimated from transverse rupture strength ( TRS) and Palmqvist tests.
IN'IRODUCTION Cemented carbides are technically and industrially one of the most important composite materials and are renowned for their high levels of hardness, compressive strength and wear resistance. These composites are, however, flaw sensitive and exhibit low toughness and impact strength. This has led to a large number of investigations into the susceptibility of these materials to defect-initiated catastrophic failure and in the last decade considerable attention has been focused on their plane strain fracture toughness (Kic). The principal aims of this paper are to provide an up-to-date assessment of fracture toughness testing of cemented carbides, to show 37 Fibre Science and Technology 0015-0568/83/0019-0037/$03.00 © Applied Science Publishers Ltd, England, 1983. Printed in Great Britain
J. D. Bolton, R. J. Keel)'
how fracture toughness is related to the structure, composition and strength of these materials and, finally, to show how fracture toughness testing of these materials can be rationalised to make it more acceptable for use in industry--with particular emphasis on pure WC~Co alloys. More than half of all cemented carbides in current use are pure WC-Co alloys which are stronger and tougher than any of the other carbide binder combinations. This fact has generated much interest in the fracture toughness testing of these alloys. For this reason, our emphasis on W C ~ o alloys is felt to be justified.
F R A C T U R E T O U G H N E S S T E S T I N G OF C E M E N T E D CARBIDES In the past decade, a large number of workers have independently reported 'Kit' (critical stress intensity factor) and 'Gic' (critical strain energy release rate) results for cemented carbides, using a variety of different test methods. 2 -25 The use of Gic values is not as popular as Ki~ particularly when comparing different materials because in using Gi¢ it is also necessary to know both Young's modulus (E) and Poisson's ratio (v) for the materials. The relationship between Ki¢ and Gi¢ is given by eqn (1). =
For this reason we concern ourselves mainly with K~c in the following sections. Strictly speaking, the designation Kic as used for fracture toughness refers only to plane strain fracture toughness measured in accordance with an accepted standard such as either the British z6 or American 27 Standard Test Method. The fracture toughness of cemented carbides cannot, however, be measured using these methods due mainly to the fact that it is not possible to precrack these materials by fatigue. At the moment, although there has been at least one 'recommended procedure '28 there is no recognised standard for the fracture toughness testing of these materials. To avoid confusion, we will as far as possible refer to published fracture toughness results by the same designation as that given to them by their author. Despite its lack of standardisation, fracture toughness testing has a number of distinct advantages over conventional strength tests, which will now be described.
Fracture toughness (Kic) of cemented carbides
Advantages and disadvantages (i)
Fracture toughness Kic is a quantitative assessment of the ability of a material to resist unstable crack propagation from a critical defect. As such, it is an intrinsic material property, unlike some of the more conventional strength measurements such as impact, tensile and transverse rupture strength (TRS), which are highly susceptible to influences from specimen geometry, surface finish and the size and distribution of microstructural defects. (ii) The fundamental nature and low scatter 14 of K~c makes fracture toughness useful as a tool in quality control, materials evaluation and comparison. (iii) Knowledge of Ki~ permits calculation of the critical size defect that will cause catastrophic failure for a given stress. This is important for cemented carbides because the critical size defect for these materials is relatively small. (iv) Fracture toughness will have a greater application to design performance as more information becomes available about the mechanical behaviour of these materials. 2s The only disadvantages of fracture toughness testing are that most of the methods used are complicated and expensive and therefore not very attractive to industry, 2 while the simpler methods, with one or two exceptions, give results which are either too high or too low. In this next section we review the advantages and disadvantages of the various precracking and fracture toughness test methods.
Precracking and fracture toughness test methods Precracking, prior to fracture toughness testing, consists of introducing into a specimen a specific and measurable critical defect 14 (normally a sharp crack). This is one of the most important and probably most difficult aspects of fracture toughness testing of these materials. In particular, if the precrack used is not sharp, planar, stress-free and of sufficient size to satisfy compliance calculations, then the result obtained cannot be considered a valid fracture toughness result. Precracking is further complicated by the fact that the stress intensity required to initiate a precrack in these materials is often very close to their critical stress intensity factor, Kic.2a Given in Table 1 is a summary of most of the precracking and fracture
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J. D. Bolton, R. J. Keely
toughness testing methods which have been used for cemented carbides. The advantages and disadvantages of these methods can be summarised as follows:
Methods A to G Advantages. Methods A to G are all relatively simple to use and the 'single edge notched beam' (SENB) specimens used are of economical size, unsophisticated design and easily manufactured. Also the use of the SENB specimens minimises the effect of specimen misalignment and gripping which usually causes problems in testing of cemented carbides, z Comparing 3 point and 4 point loading for the fracture toughness testing of these specimens, the latter has the advantage that no shear stresses are present in the central portion of the beam and is therefore superior 33 even though SENB specimens used in 4 point bending are normally slightly larger than for 3 point bending. Disadvantages. The main disadvantage associated with methods A and B is that in machining the notch, thermal damage is introduced at the notch root in the form of microcracks and thermally damaged material. Although some allowance is made for this in method B, there is still some uncertainty as to how this influences crack extension and hence fracture toughness. If these uncertainties could be removed, these methods would be particularly useful in that they are quite simple and are applicable to a wide range of cemented carbides. Methods C and D suffer from a major disadvantage in that the notch root radius (p) produced by diamond slitting is too large with the result that fracture toughness obtained using these methods tend to be too high by a factor of 2 to 3. The main disadvantage associated with methods E and F is the removal of the plastically deformed zone produced on indentation. To obtain a valid fracture toughness result, it is necessary to remove the residual stresses caused by indentation using an annealing treatment after precracking. This, however, results in further complications, in that with full stress relief by annealing, the precrack tends to heal by diffusion bonding and the surrounding plastically deformed zone becomes tougher 2 thus giving fracture toughness results which are too high. If shorter annealing times are used the stresses are not fully relieved and the results obtained are too low. Removal of the deformed zone by machining could solve the problem but, since this zone is large in relation to the
Fracture toughness (Kit) of cemented carbides
size of the precrack, it would be difficult to avoid complete removal of the precrack. 2 Of the seven methods A to G, method G suffers the least number of disadvantages. The precrack, produced by wedge indentation, is sharp, planar, of uniform depth and easily measured. The plastic zone is small relative to the precrack length and can be easily removed by machining. 32 The only apparent disadvantage of this method is that it is difficult to use on the tougher cemented carbides due to their resistance to precrack initiation. Methods H to L Advantages. Valid fracture toughness values should be obtained by careful use of these methods, which is in itself an obvious advantage. Additional advantages of methods H and J are that a number of fracture toughness results can be obtained from one specimen and precracking and fracture toughness testing is performed in one continuous operation using the same jig. Disadvantages. The main drawbacks of methods H to L are that they are complicated and expensive, requiring closely controlled laboratory conditions and are unsuitable for routine fracture toughness testing. A further criticism of methods H and J is that the specimens used are relatively large and of complex design, they are much more expensive to produce than SENB specimens and, because of their size, are more likely to contain non-homogeneous microstructures and composition. 34 Method M While method M has advantages and disadvantages similar to methods H and J it suffers one further drawback in that there is uncertainty about the nature of the stress-field near the crack tip in the specimen used. It is felt that this method may measure a mixed mode of crack opening at the critical stress intensity, rather than pure mode 1 opening. If this is so, then the fracture toughness results produced by this method cannot be considered to be Kic results. There is also some doubt about the method used to calculate Kic. The crack front produced in the specimen is curved and does not penetrate the upper surface of the specimen which remains in compression, while the expression used to calculate Kic assumes a straight crack front perpendicular to the upper and lower surfaces of the specimen.
J. D. Bolton, R. J. Keel)'
Method N Little is known of this recently announced method except that it has been used to produce results which are comparable with those obtained using some of the other precracking methods. There are some doubts, however, about the validity of precracks produced by this method particularly since the depth of precrack may be inadequate to satisfy a/w (a = crack length, and w = specimen width) compliance requirements and because TiC and eta-phase in the coated layer may subsequently affect crack extension in the bulk specimen. Apart from those methods in Table 1, there are a number of other methods which have been used to evaluate the toughness of these materials, the two most notable being 'TRS~lefect size correlation' method and the Palmqvist toughness method. Other methods used f o r toughness evaluation TRS-defect size correlation method. TRS tests on cemented carbides 35'a6 have been shown to give a direct relationship between the size of the fracture initiating defect and the stress acting on it at the time of failure (af), of the form a f = Y Kicd - 1/2
where Y is a test geometry constant and d is either half the length of the major axis of the defect 35 or the square root of the area of the defect. 36 a j can be calculated if the position of the critical defect relative to the load application points is measured. A number of workers have used this relationship in conjunction with TRS tests to produce fracture toughness results which they described as being either equal t o 21 o r directly related to Kic. 8'17'36 Although this method helps explain the dependence of TRS and fracture toughness, the evaluation of Ki¢ using this method would be tedious 42 and not entirely accurate. The Palmqvist method o f toughness measurement. The Palmqvist method gives a measure of the surface toughness of a brittle material 3s by measuring its resistance to crack propagation from the corners of a Vickers Diamond Pyramid indentation made on a flat stress-free surface. 21 Palmqvist toughness (W) measured in Jm -2 is defined 39 as the
Fracture toughness (Ki¢) of cemented carbides
indenting load (P) divided by the sum of the four crack lengths (L) at the indentation corners: W = P/L
This method has been used for a number of years in both research and industry, particularly for the comparative ranking of different grades of cemented carbides. 4° Its popularity is mainly due to the easy manner of performing tests on simple test pieces without complicated equipment. The Palmqvist method, however, suffers from a number of disadvantages, especially when compared with some of the better Kic methods. One of the most serious drawbacks is that there is as yet no simple correlation between the measured Palmqvist surface toughness and bulk fracture toughness. It has been shown that there is an approximately linear correlation between W and Gic for W C - C o alloys containing up to 10wt ~o Co and this has been used to calculate 'Gic' values for a number of W C - C o alloys. 22 There is some doubt about this approach. Because the W results of one author 22 were correlated to the Gic results of another author.18 (ii) The usefulness of these G~c results are in themselves questionable, being subject to uncertainties relating to the measurement of Kic, E and v. (i)
Other disadvantages associated with the Palmqvist method are: (i)
The specimen preparation required to achieve a stress-free surface is time-consuming and expensive. (ii) Due to the high loads used in indenting the specimens, diamond indenters tend to fail after a relatively small number of tests. (iii) This method is only applicable to a limited range of cemented carbide materials. (iv) Although two separate analyses 38'44 have been used to elucidate the nature of Palmqvist cracking, as yet there exists no fully comprehensive theoretical basis on which to explain this phenomenon. 42 With these disadvantages in mind, it is unlikely that this method will have much to offer in the fracture toughness testing of cemented carbides although it may be of use in understanding the behaviour of these materials in applications where they are subject to indentation cracking 42 for example in drilling and mining.
J. D. Bolton, R. J. Keely
Having looked at the relative merits of the various methods used to evaluate fracture toughness of these materials, the results obtained will now be examined.
FRACTURE T O U G H N E S S RESULTS In this section we examine the relationships between fracture toughness and the structure and composition of cemented carbides which are of prime importance in understanding the strength of these materials. Dependence of Kic on microstructural parameters The main parameters used to describe the microstructure of cemented carbides, shown typically in Fig. 1 are as follows: (i)
Volume fraction of binder phase (VFB), or Volume percent of binder phase (V%B) (ii) Mean carbide grain size (dc) (iii) Mean free path of the binder phase (2) (iv) Contiguity of the carbide phase (C) all of which can be measured by quantitative metallographic analysis. 45 The interrelation of these four parameters is given by eqn (4) 23 =
( 1 - v B)(1 - c )
From eqn (4), it can be seen that to fully describe the microstructure it is necessary to know at least three of these four parameters. Despite this, cemented carbides are often described in terms of V%Rand dc only. This is because the contiguity changes only slowly (i.e. with the logarithm of sintering time) 46 so that the microstructure of cemented carbides prepared under normal conditions can be adequately described by V%B and dc. 4s However, in the case of cemented carbides, with the wide grain size distributions or with alloys prepared using unconventional methods, it is important that a third parameter be used to describe the microstructure, either 2 or C. Some workers have published Kic results for cemented carbides giving only V%Bwhile defining grain size, dc, merely as fine, medium or coarse. By describing the microstructure in this manner, there is insufficient
Fracture toughness (Kic) of cemented carbides
Fig. 1. A crack traversing a typical WC-Co microstructure ( x 5000).
information available to enable a realistic comparison to be made between results from different sources. There is also some doubt about the practice of quoting cobalt content of these materials in weight percent, as it is felt that these values, which are sometimes notional, may not give a very accurate indication of the true binder content. It has been shown that the fracture toughness Kic of WC4So alloys increases with increasing binder content, 5'15'16'19'2°'23 mean carbide grain size2' 5,1 2 , 1 5 , 1 6 , 1 9 , 2 0,2 3 and mean free path of binder phase, 5,15,16,19 while it decreases with increasing contiguity. 5'2°'23 Similar relationships have also been shown to exist between Gic and these parameters.5'16'19 A linear model, regression analysis used on two sets of published results 2°'23 has shown that there is significant correlation between Kic and each of these four microstructural parameters, while the overall correlation (R) at 0-98, is very high, for both sets of results (see Table 2, columns 1 and 2). The relationship between K~¢and these four parameters can be expressed by eqn (5) gic = a o + a I V%B -[- a2dc q- a3,~ -t- a a C
It should also be noted from Table 2 that the partial correlation coefficients (r) between V%B, 2 and Ki¢ are relatively high and approximately equal, whereas the coefficients for dc, C and Ki~ are somewhat lower. Previous investigations of the relationship between the fracture toughness and microstructure of these materials have concentrated
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mainly on the relationship between Ki¢ and 2. Although the correlation between Ki¢ and 2 is quite high ( ,-~0.9), the overall correlation between Kic and these four parameters is obviously more useful, being so much higher. We have also found, however, that the relationship between Kic and these parameters can be simplified, without loss in overall correlation, by expressing K~ in terms of V%a and dc only, which suggests that these two parameters alone are adequate to describe the dependence of K~¢ on microstructure as shown in eqn (6) (see Table 2, columns 3 and 4) Kic = ao + al V%B + a2clc
Using this approach, with a total of 106Kic results, from eleven d i f f e r e n t s o u r c e s 2'3'5'8'12'15'16'19'2°'23'25 gives eqn (7), with an overall
correlation coefficient of 0.89, which is again quite significant (see column 5, Table 2). Kic = 5.286 + 0.291V%B + 1.683dc
Equation (7) is shown in Fig. 2 superimposed on a graphical summary of most of the Kic results published to date for WCq2o alloys. 2-25 From
dc -- g r a i r l size
Fracture toughness (Kic) of WCq2o alloys. (a) Diamond slitting methods; (b) hardness indentation methods.
J. D. Bolton, R. J. Keely
this it can be seen to be in good agreement with the overall trend of the results, while at the same time Kic results obtained by using either hardness indentation or diamond slitting prepacking techniques, differ significantly from the other results. This equation is not intended as a definitive model for the dependence of Kic on microstructural parameters, however, it is expected to be useful in that: (i)
It gives a good indication of the dependence of K~ on volume ~o cobalt and mean carbide grain size for WC Co alloys, (ii) It provides a rough estimate of the expected Ki~ value given V%B and &', (iii) It provides a method whereby Kic results produced by various workers, using different precracking methods, can be compared with a view towards the standardisation of fracture toughness testing of these alloys. One reservation we have about eqn (7) is that for some of the results used in producing the equation, cobalt contents were given only in weight percent and needed to be converted to V%B.
Kic models There have been a number of theories and models put forward to explain the dependence of Ki~ for W C ~ o alloys on microstructural parameters, s'ls'1<19'2°'23'47 However, these have met with only limited success in that, although some of these models contribute qualitatively towards a better understanding of the interaction between microstructural parameters and fracture processes, there are still a number of fundamental problems. One of the main problems encountered is that there is still a large degree of uncertainty as to the exact nature of the fracture paths that result on critical crack extension. There are four principal types of fracture path that occur: (i)
De-cohesion of carbide/carbide interfaces, which occurs mainly in alloys of low cobalt content. (ii) De-cohesion of carbide/binder interfaces which predominates at higher cobalt content. (iii) Transgranular cleavage of carbide grains--normally associated with large carbide grains.
Fracture toughness (Kic) of cemented carbides
(iv) Intergranular rupture--where the fracture passes through the binder phase. This becomes noticeable with increasing binder content. These four types occur in differing relative proportions, depending on the microstructural parameters, but it is often quite difficult to distinguish between them even for any one alloy. Another problem is that although there is reasonable agreement between the K~c results of those authors who have produced models, there are still differences of 30-40 % between some of their Kic results for similar alloys. These differences must be resolved before an acceptable overall model can be produced.
Relationship between fracture toughness and strength of WC-Co alloys The compressive, tensile and transverse rupture strengths of these materials each pass through a maximum with increasing cobalt content. 48 These maxima, while usually occurring at a low cobalt concentration, are achieved at some optimum 2 value, which is a function of V%B, dc and C. There is, however, difficulty in relating these mechanical properties to microstructural parameters, as the results obtained represent the sensitivity of the material to defect initiated failure rather than the true strength, particularly with the TRS and tensile tests on low cobalt alloys. Attempts have been made to relate strength test results to defect initiation both by statistical methods such as the Weibull Distribution 49'5° and the previously mentioned 'TRS defect-size correlation method', 35'36 but have proved to be of limited use. While the use of Weibull statistics gives a reasonable correlation between microstructure and susceptibility to failure, it does not help identify the fracture initiation source nor does it define how the actual strength varies. Using the 'TRS defect-size correlation method', a measure of the 'intrinsic strength' of the material can be found by relating the stress acting on the defect at the time of failure (af) to the stress concentration caused by the defect size and shape. By analysing TRS tests for the type and location of fracture initiating defects, it is possible to determine the effect of various microstructural parameters on the intrinsic strength. This approach is, however, of limited value in explaining the strength of these materials, because intrinsic strength is calculated on the unlikely assumption that elastic conditions exist around the defect at failure. Intrinsic strength values may not therefore reflect true strength-structure relationships, but
J. D. Bolton, R. J. Keel),
may instead indicate fracture under the varying conditions of plasticity that may exist in the region surrounding the defect tip. By considering the interaction of defect-related strength (ad), calculated from eqn (8), and the ultimate strength (au) of these materials, we have found that these strength maxima can be easily explained. ad = Y K i c a
The maximum achievable strength of these materials is controlled by such factors as yield strength, work hardening rate, and ductility with the result that it decreases with increasing 2 as shown schematically in Fig. 3. Theoretically, therefore, these alloys with low 2 values should have highest strength. However, because the ability of these low 2 materials to produce crack tip blunting, by plastic deformation in the binder phase, is severely diminished, the strength of these alloys is controlled by the size (a) of the microstructural defects present in the material as shown by eqn (8) and Fig. 3. Therefore at low 2 values, although the potential ultimate strength of these materials is quite high, this is not achieved due to defect initiated failure. It is also worth noting that with decreasing cobalt content, below the strength maximum, the number and size of defects present in the materials is likely to increase, due to poor sintering behaviour, so that the fall in strength of these alloys can be quite steep below optimum ,~values. od -
strength (o U) Mean free path
Fig. 3. Schematicrepresentationof the relationship betweenthe maximum achievable strength (ultimate strength) and the fracture toughness of WC-Co alloys.
Fracture toughness (Ki¢) of cemented carbides
At higher cobalt contents, beyond the maximum strength value, the strength of these materials is not as sensitive to defect initiated failure, due to increased plasticity of the binder phase and is therefore controlled by their ultimate strength. F r o m this, the strength of these materials is seen to be related to their microstructural parameters and inherent defect size. Furthermore, it can be seen that the potential for increased maximum strength value, rests with the possibility of increasing K~¢ and tru and decreasing the inherent defect size of these materials.
Effect of carbon content of Ki~ To date very little conclusive evidence has been published concerning the dependence of fracture toughness on carbon content. Palmqvist tests have shown that the surface toughness of W C - C o alloys reaches a maximum at stoichiometric carbon content, that is 6.13wt% C in WC, and tends to decrease with increasing carbon deficiency 22'4° or carbon excess. 4° Slightly sub-stoichiometric alloys containing scattered patches of eta phase have, however, been shown to possess similar fracture toughness to the stoichiometric alloy. 23 Preliminary indications from fracture toughness Ki¢ tests recently performed by us on WC-8-5 vol % Co alloys 32 show that the effect of carbon content on the toughness of these is similar, at least qualitatively, to the effect of carbon content on TRS. 48 That is, that fracture toughness remains approximately constant within the range 0-5 vol % r/phase to 0.5 vol % free graphite, with the maximum fracture toughness occurring at o p t i m u m carbon content or when a very small amount of structurallyfree graphite is present. Outside these limits there is a drop in fracture toughness--the decrease being more pronounced for sub-stoichiometric carbon content than for excess carbon.
Temperature dependence of Kic A b o v e r o o m t e m p e r a t u r e . Fracture toughness Kic for WC-Co alloys has been shown to be relatively insensitive to temperature between 20 °C and 300°C. J8 Two other investigations 1v'3° have shown fracture toughness to be relatively unaffected by temperature up to 800 °C and while Dawihl 3° has shown that fracture toughness drops sharply above 800 °C, Mail 7 has shown fracture toughness to rise significantly between
J. D. Bolton, R. J. Keely
800 °C and 1000 °C. There are, however, some doubts about the validity of either of these results because of the precracking methods used. As temperature increases, mechanical strength will decrease and so it is to be expected that above a certain temperature fracture toughness will also decrease particularly in low cobalt alloys. There are, however, two other factors which will tend to oppose this drop in toughness:
Crack-tip blunting due to plastic deformation of the binder phase--which will tend to predominate in alloys with high V%~. (ii) Multiple microcrack formation caused by carbide/carbide decohesion which will tend to occur mainly in alloys with low 2. Because of these uncertainties it is difficult to say precisely what effect temperature will have on fracture toughness above 300°C. B e l o w r o o m temperature. No appreciable difference has been found 2° between fracture toughness results at - 196 °C and at room temperature for a number of different WC~Co alloys. Fracture paths at - 196 °C were also reported as being very similar to those at room temperature, z° Suzuki51 has reported that the TRS of W C - C o alloys actually increases at temperatures below room temperature, which is incompatible, however, with the fact that no increase in K~c is indicated. Heat treatment
There is very little documented evidence on the effect of heat treatment on fracture toughness of these materials. There are, however, at least two effects of heat treatment which would be likely to alter their fracture toughness: (i)
Carbide grain growth--which occurs mainly with prolonged sintering times and slow cooling from sintering temperature. As already shown, increased grain size will increase fracture toughness. (ii) Structural alterations within the binder p h a s e ~ h i c h tend to occur mainly between 600 ° and 1 1 0 0 ° C . 5 2 - 5 4 Changes which reduce binder phase plasticity, e.g. precipitation, could reduce fracture toughness, whereas changes which increase plasticity e.g. compositional changes, could improve fracture toughness. Whatever changes do occur, however, are likely to be more significant in alloys of high cobalt content.
Fracture toughness (Kit) of cemented carbides
Sintering temperature It has been shown 31 for a WC-10Wt~o Co alloy, that sintering temperature has a pronounced effect on fracture toughness KQ, i.e. fracture toughness, (measured on notched specimens with notch root radius (p) = 0.1 mm). KQ reaches a maximum at sintering temperatures between 1400 and 1500°C, and thereafter falls by about 20% between 1500 and 1600 °C. While there are some doubts about this method of fracture toughness measurement, the trend shown by these results is in good agreement with the sintering behaviour shown by these materials. Below the melting points of the ternary W ~ o - C eutectic (1280 °C) and the binary Co-C eutectic (1315°C), WC~Eo powder compacts have negligible strength. 5z Above these temperatures cobalt infiltration causes a steep rise in their strength, and optimum sintering behaviour is achieved somewhere between 1400 and 1500 °C. At temperatures above 1500 °C pronounced anisotropic growth of some of the carbide grains and 'swelling' results in a loss of mechanical strength. 52 Hot isostatic press (HIP) This has been shown to have almost no effect on either Kics'19 or Palmqvist toughness 4° of W C - C o alloys, which might be expected, as H I P serves only to reduce defect population, and does not alter the fundamental microstructure. Marked improvements in the TRS of these alloys after H I P can therefore be traced directly to the reduction in defect concentration. F R A C T U R E T O U G H N E S S OF C E M E N T E D CARBIDES O T H E R T H A N P U R E W C ~ o ALLOYS Very little published information exists on the fracture toughness of these alloys, and that which does exist tends to be inconclusive. 17,30,31 General trends for these alloys show the fracture toughness with similar dependence on binder content and mean carbide grain size, but their fracture toughness and Palmqvist toughness 4° values tend to be significantly lower than for pure W C - C o alloys. For mixed carbide compositions (WC/Tic/TaC/NbC-Co), TiC tends to be most detrimental to toughness, while TaC tends to have least effect. 40 Cemented carbides containing binders other than cobalt, such as Ni, Fe, NiFe and NiMo also show, on average, lower toughness. 48'52
J. D. Bolton, R. J. Keely
CONCLUSIONS The conclusions we have come to in relation to fracture toughness of cemented carbides are as follows: (i)
There is at the moment a definite need for a standardised approach to the fracture toughness testing of these materials in order that it may become more acceptable for industrial use. Of the methods considered, there are two methods of potential in this area. These are wedge indentation precracking, and ' E D M plus thermal microcracking'. The range of fracture toughness (Ki¢) values found for WC-Co alloys between 5 and 40 volume ~o cobalt is quite limited--at most 8 - 2 6 M P a m 1/2. Because of this it is most important (a) that reliable precracking and fracture toughness methods be used, and (b) that the fracture toughness (Kic) results be measured accurately. This is particularly important if it is intended to find meaningful correlations between the fracture toughness results obtained, the microstructure of the alloy, and the observed fracture processes. Other methods of assessing fracture that have been used, such as Palmqvist toughness testing and TRS~tefect size correlation, have little potential when compared with some of the better Kic methods. Fracture toughness in WC-Co alloys depends strongly on each of the four parameters used to define microstructure and a useful correlation exists between K~ and the two microstructural parameters--volume ~o binder phase and mean carbide grain size. For a more precise definition of the microstructure of these alloys, however, at least one other microstructural parameter must be known, e.g. contiguity. Although a qualitative understanding of the interaction between microstructural parameters and fracture processes exists, there is a need for a fuller understanding of the relative importance of individual fracture mechanisms so that models can be developed to explain the overall dependence of K~¢ on microstructural parameters. The use of fracture toughness theory in clarifying the nature of the strength of these materials has distinct advantages over the methods that have been previously used.
Fracture toughness ( K~c) of cemented carbides
(vii) Little is known about the dependence of fracture toughness of the W C - C o alloy materials on temperature and carbon content. There is also little known about the fracture toughness of cemented carbides other than pure W C - C o alloys. It is expected, however, that future work in these areas will contribute significantly in our understanding of these materials. (viii) Finally, considering the numerous advantages and the potential uses o f the fracture toughness (Kit) approach in helping to define and improve the strength o f cemented carbides, there is little doubt but that it will play an important role in the future industrial development and application o f these materials.
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