Phase composition in cemented carbides and cermets

Phase composition in cemented carbides and cermets

Int. J. of Refractory Metals & Hard Materials 12 (1993-1994) 107-113 1994 ElsevierScienceLimited Printed in GreatBritain 0263-4368/94/$7.00 r ELSEV...

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Int. J. of Refractory Metals & Hard Materials 12 (1993-1994) 107-113

1994 ElsevierScienceLimited Printed in GreatBritain 0263-4368/94/$7.00

r

ELSEVIER

Phase Composition in Cemented Carbides and Cermets H.-O. Andr6n," U. Rolander b & P. Lindahl" "Department of Physics, Chalmers University of Technology,S-412 96 G6teborg, Sweden hAB Sandvik Coromant, R& D Materials and Processes, S-126 80 Stockholm, Sweden

Abstract: Cemented carbides and cermets of types WC-Co, WC-MC-Co,

TiC-TiN-Mo2C-Ni/Co and (Ti,W, Ta)(C,N)-(Co,Ni) have been studied with atom probe field ion microscopy and analytical electron microscopy. All materials were found to have a hard phase skeleton with grain boundary segregation of binder metal. The binder phase contained metal atoms from hard phases but was almost completely free of C and N. Cobalt binder phase contains W and Mo but little Ti, nickel binder phase Ti and Mo. The amount of Ti did not vary much with N content in the alloy, whereas the Mo content increased substantially with total N content. Materials containing more than one hard phase formed a core-rim structure. No diffusion of metal atoms occurred in the hard phases during sintering, but C and N may diffuse so that the system approaches ( p a r a ) - e q u i l i b r i u m with respect to these elements during sintering.

added. Titanium-based cemented carbides are at present winning an increasing market share, and they are now usually known as cermets, especially when they contain nitrogen. It is interesting to note that as Co replaces Ni and W replaces Mo cermets are approaching cemented carbides in composition. Much work has been done on phase diagrams for several of the relevant carbide-metal systems, and detailed work on carbonitride systems is now under way. It could therefore be expected that it should be a straightforward procedure to predict the microstructure of cemented carbides and cermets. However, this is not the case, not even for the simple WC-Co system, as we shall see later. The microstructure of cemented carbides and cermets is more complicated than expected mainly for the following three reasons.

INTRODUCTION Cemented carbides are hard and tough tool materials consisting of carbide grains embedded in a metal binder phase. They are fabricated from carbide and metal powders by liquid phase sintering at about 1300-1500°C. Plain WC-Co was invented in 1923 and was for a long time the dominant type. Additions of TiC and other cubic carbides (TaC, NbC, HfC... ) were later made to WC-Co to improve its steel cutting performance. During the last 20 years wear resistant coatings have dramatically improved the performance of WC-Co-based materials, which are still the most important type of cemented carbides. TiC-Ni-based cemented carbides, with Mo2C additions to improve wetting and densification during sintering, are very hard and had in the past a limited use, e.g. for fine turning, but suffered from a lower toughness than WC-Co-based materials. In the 1970s it was found that TiN additions to TiC-Ni greatly improved toughness. An intense development work followed, which resulted in commercial grades in which Ni was partly replaced by Co, and carbides and carbonitrides of, for example, Ta, W, Nb or V were

• The composition of the material changes during fabrication. During ball-milling small amounts of, for example, Fe, Cr and Ni from stainless-steel walls and W, C and Co from cemented carbide milling bodies are added to the powders. Co may oxidise during mill107

108

H.-O. Andrdn, U. Rolander, P. Lindahl

ing, and the oxide subsequently consumes carbon when reduced during sintering. Carbon may be added from remnants of the pressing wax (PEG). Sintering in vacuum or in an inert gas gives rise to decarburisation and denitrification. • Equilibrium cannot be achieved during sintering if more than one hard phase (i.e. carbide, carbonitride or nitride phase) remains undissolved at the sintering temperature, because of the slow diffusion of metal atoms in the hard phases. Therefore, non-equilibrium phases are retained. However, the diffusivity of C and N is sometimes sufficient to allow the system approach (para)-equilibrium with respect to these elements. • Phase compositions are not given by equilibrium relations at the sintering temperature, since cooling after sintering is generally slow under production conditions. Instead, the equilibrium composition at the temperature where diffusion ceases will be frozen in, and composition gradients will form. Consequently, to understand the microstructure of cemented carbides and cermets--simple model alloys as well as complex commercial grades-detailed microstructural characterisation is necessary. A considerable amount of work using X-ray diffraction and various types of microscopy and microanalysis has therefore been made in the past, and has been reviewed previously. 1-4 However, more conventional methods of microanalysis such as analytical electron microscopy have difficulties in coping with small phase areas, fine scale segregation and accurate analysis of the light elements Cand N. We have therefore used atom probe field ion microscopy, in conjunction with analytical electron microscopy, to study the detailed microstructure of cemented carbides and cermets. Specimens for the atom probe instrument have the shape of a sharp needle. As a high positive voltage is applied to the specimen, the high electric field at the specimen extracts atoms, which are individually analysed in a mass spectrometer. The radial projection nature of the instrument gives it a uniquely high spatial resolution, about 1 nm both laterally and in depth. With its high spatial resolution and (theoretically) equal detectivity for all atomic species the atom probe instrument is well suited for the study of cemented carbides,

and we have worked with these materials since 1980. 5,6 However, to achieve good quantitative accuracy of atom probe analyses is not a simple matter, 7-~° and considerable effort had to be spent to develop the instrument for this p u r p o s e . II Conditions of analysis have to be carefully selected, especially for Ti-based materials. 12,13 A summary of our development of methods for atom probe analysis of cermets was given at the last Plansee seminar.~4 The purpose of this paper is to give an overview of our results on the microstructure of cemented carbides and cermets. Only macroscopically homogeneous materials are included in this paper, which means that, for example, coated materials are not treated. WC-Co It might be expected that plain WC-Co cemented carbide should be a very simple system. Since WC is known to be accurately stoichiometric and does not dissolve any cobalt, the liquid cobalt phase ought to dissolve W and C in atomically equal proportions during liquid phase sintering. During cooling after sintering, W and C would then reprecipitate onto WC until the temperature is too low for any diffusion to occur. However, in practice the total amount of C in the material can be controlled since carbon is lost in the process, as explained above, and C may be added as soot. W-C-Co phase diagrams are well known, and a section at 1250°C is shown in Fig. 1. Included in the diagram are also our atom probe measurements of the binder phase composition in WC-20%Co after heat treatment at fixed carbon activities (ac = 1 and 0.4) and subsequent quenchingJ 5 From the diagram it is evident that the W content in the liquid phase during sintering can be controlled by the total carbon content in t h e material. The hardness of the binder phase is determined by its W content, which means that the total carbon content is important for the hardness and toughness of the material. It should be noted, that too much or too little carbon results in the formation of graphite and r/-phase (M6C), respectively, both of which are detrimental for the toughness of the material. After sintering W and C reprecipitates in atomically equal amounts onto WC grains. Figure 2 shows a W composition profile in the binder phase of a WC-20%Co cemented carbide, sin-

109

Phase composition in cemented carbides and cermets FCC-Co+Graphite+WC

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Fig. 1. The cobalt-rich corner of a calculated isothermal section of the C o - W - C system at 1250°C. Atom probe measurements of the binder phase composition at a carbon activity of 1 and 0"4 are shown. Adapted from Ref. 15.

tered at 1360°C and subjected to the normal slow furnace cooling. A tungsten gradient towards an adjacent WC grain is seen. Calculations showed, that in the interior of the binder phase grain the equilibrium composition freezes in at about 1000°C, and close to the WC grain at a temperature a few hundred degrees lower. Essentially no carbon was found in the binder phase. This can be understood from the low carbon activity in the sintering furnace and the high diffusivity of C in the binder phase. In addition, as the temperature falls an increasing supersaturation of W in the binder increases its carbon activity. 16 Thus, commercial products after normal sintering contain tungsten but no carbon in the Co binder phase. An old matter of debate concerns the exact microstructure of the WC phase in WC-Co. Because of the high rigidity of WC-Co it was long ago suggested that (for normal volume fractions of binder phase, -<20%) the WC grains make up a continuous rigid skeleton. ~7 However, to explain the high toughness of W C - C o it was later postulated that thin binder phase layers existed between the WC grains, is Atom probe analysis through WC-WC boundaries showed that they contain about 0.5 monolayers of Co, strictly localised to

the boundary plane (Fig. 3), and the same result was obtained using analytical electron microscopy) 9-21 Consequently, the WC grains form a skeleton, in which Co is present as grain boundary segregation. The grain size of the binder is very large, about 1 mm. 22 Thus, both the carbide skeleton and the binder phase are continuous through the material. Small amounts of Cr3C 2 and some other carbides are sometimes added to WC-Co to restrict grain growth during sintering. Chromium carbide dissolves during sintering, and Cr segregates to the grain boundaries in the WC skeleton where it replaces some of the Co segregaron. 21 The solubility of both Cr and Co in WC is very low. Presumably, it is more difficult for a grain boundary to drag Cr atoms than Co atoms along with it during migration. This restricts grain boundary migration and grain coalescence during sintering.

WC-MC-Co WC-Co with addition of cubic carbides is a very important class of cemented carbides, much used for metal cutting tools. The cubic carbides (yphase) have the general formula MC (NaCI structure) where M stands for Ti, Ta, Nb etc. or a

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combination of these (mixed carbides). It should be noted, that the MC phases may be sub-stoichiometric with respect to C which gives the system one further degree of freedom. The carbide grains form a WC-MC skeleton, and segregation of Co (approximately half a monolayer) is found in WC-WC, WC-MC and M C - M C boundaries. 21 We have studied the microstructure of an 85% WC-8(Ti, W)C-7Co grade with atom probe and analytical electron microscopy. 23 The Co binder phase contained very little Ti, so solid solution hardening of the binder is brought about by its W content, which is controlled by the carbon content during liquid phase sintering. As for plain WC-Co the C content in the binder after sintering is very low. Similarly we have found that the solubility of Ta in the cobalt binder of TaC containing cemented carbides is very low. The MC phase was found to have a core-rim structure (similar to a cermet microstructure, see Fig. 5 below). The cores were found to be undissolved remnants of the (Ti0.8W0.2)C powder. They had the same metal composition as the original powder and contained some impurity nitrogen and dislocations from the milling process. The cores act as nucleation sites for rims, which form during sintering by epitaxial precipitation. (Ti, W)C rims had a higher W content, a low impurity N content and were essentially free of dislocations. 23 They reflect the composition of the binder phase from which precipitation took place. The core-rim boundary is a purely compositional boundary, since the crystal structure is continuous across it. The compositional boundary is very sharp, and the W diffusion profile into the core is less than 1 nm wide (Fig. 4). Consequently, thermodynamic equilibrium is not achieved during sintering.

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Fig. 4. Atom probe composition profile through a core-rim compositional boundary in WC-(Ti,W)C-Co cemented carbide. Diffusion of W into the TiC core is less than 1 nm. The core contains some impurity N. Adapted from Ref. 23.

MODEL CERMETS

A series of model cermets (designated '1' to '5' in figures), containing TiC, 10 vol.% Mo2C, 15% binder metal and between 0 and 20% TiN were prepared and characterised in detail. 2+ The series contains two sets of alloys, one with Ni binder and one with Co binder. All materials contained a skeleton of carbonitride grains with a core-rim structure (Fig. 5). Grain boundary segregation was not measured in this series, but a cermet with a mixed Ni-Co binder contained about half a monolayer of Ni + Co. 23 The development of the core-rim structure was followed by interrupting the sintering at several temperatures and studying the material with X-ray diffraction and scanning electron microscopy. It was found that Mo2Cdissolves in the binder below 1200°C, well before any liquid is formed (eutectic temperatures were in the range 1310-1335°C). At higher temperatures TiN dissolves, so that only TiC cores, rims and liquid is present at the sintering temperature of 1520°C. Rims had the general composition (Ti,Mo)(C,N) and consisted of two parts: the inner rim, richer in Mo and formed already during solid state sintering when Mo2C dissolves; and an outer rim, leaner in Mo and formed during liquid phase sintering. Cores were remnants of the TiC powder. Very little diffusion of Mo into cores was seen (Fig. 6). Mo was found only at planar faults in TiC, which obvi-

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the (Ti,Mo)(C,N) phase (Fig. 9). The binder phase is quite different in Ni and Co based cermets (Fig. 10). The solubility for Ti is dramatically different in Co (8-10at%) and Ni (0.5%) but varies little with total N content of the alloy. The solubility for Mo is not so different (1-5% in Ni, somewhat more in Co) but increases quickly with total N content. Very little N was detected in the binder phase; some little C was observed only in N-free material. From the present results it is obvious that the amount of Ti and Mo in the binder can be controlled both by C and N contents in the alloy and by using a mixed Ni-Co binder. Binder phase compositions in materials with a mixed binder have been the subject of another study which is presented separately at this conference. 2~'

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Fig. 6. C o m p o s i t i o n p r o f i l e s at c o r e - r i m b o u n d a r i e s in TiC-20vol.%TiN- 10%Mo2C-15%Ni cermet. Atom probe analyses accumulated from several specimens. Adapted from Ref. 24.

ously act as diffusion paths for metal atoms in TiC. By contrast, considerable amounts of Ni had diffused into cores and C out of cores, so that the system was approaching (para)-equilibrium with respect to C and N at 1520°C (Fig. 7). The composition of (outer) rims as a function of total N in the alloy is shown in Fig. 8. The N content increased linearly with total N / ( N + C ) content in the alloy, whereas no difference was found in Ti and Mo contents. However, by comparing analyses from different parts of the rims it was found that Mo increases the activity of N in

COMMERCIAL CERMETS An alloy was studied (designated W in the figures) with the total composition 36 at.% Ti-5 W-3 Ta26 C-17 N-6 Co-6 Ni (balance impurity O, Fe and M o ) . 25'27 It was fabricated from several different powders and sintered in vacuum at 1480°C for 2 h. Two types of cores were observed: (W, Ti)(C,N) and Ti(C,N). The (W, Ti)(C,N) cores were remnants of a (W, Ti)C powder, which had been strongly decarburised during sintering. Only a part of the carbon loss was replaced by N. As expected from the low diffusivity of metal atoms in the hard phases, no change in metal content had occurred. Ti(C,N) cores stemmed from a Ti(C,N) powder in which the C content had decreased. No diffusion of metal atoms had

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Fig. 10. Atom probe analysis of the binder phase in model cermets as a function of total N/(N + C) in the alloy. Adapted from Ref. 24.

note that although the commercial alloy (A) and the model alloys (1-5) were prepared from very different powders and had quite differentcompositions, the fact that the Ti + Ta and Mo + W content was approximately equal gave equal Ti + Ta and Mo + W contents in the rims. The N content was given by the total N/(N + C) in the alloys, just as was the case for Ti(C,N) cores. CONCLUSIONS

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commercial alloy(A). Adapted from Ref. 25. occurred, except in planar faults, where W and Ta was observed. In Fig. 6 the core composition is compared with that of the model cermet alloys. The linear relationship between core N content and total N/(N + C) in the alloy is still obeyed. Rims had the general formula (Ti,W, Ta)(C,N) and had approximately the same composition at both types of cores. Sometimes enrichment of W was observed close to rim-core boundaries, in analogy with Mo enrichment in the inner rims of the model cermet alloys. The composition of Ti(C,N) cores is given in Fig. 7. It is interesting to

• The detailed microstructure of cemented carbides (WC-Co and WC-MC-Co) and cermets (TiC-TiN-Mo2C-Ni/Co and (Ti, W, Ta)(C,N)-(Co,Ni)) was studied with atom probe field ion microscopy and analytical electron microscopy. • All materials were found to have a carbide (or carbonitride) skeleton. About half a monolayer of binder metal segregation was found in grain (and phase) boundaries of the skeleton. • The binder phase contains metal atoms from hard phases which dissolve during sintering. The metal content is not the eqiulibrium concentrations at the sintering temperature, but the composition at a lower temperature is frozen in. Depleted zones exist close to hard phases, and the binder is almost completely free of C and N. • Solid solution hardening of the binder is due to metal atoms only. Co binder phase contains W and Mo but little Ti, Ni binder phase Ti and Mo. The amount of Ti does not vary much with N content in the alloy, whereas the Mo content increases substantially with total N content.

Phase composition in cemented carbides and cermets



Materials containing m o r e than o n e hard phase that is undissolved during sintering f o r m a c o r e - r i m structure. Rims precipitate epitaxially o n t o cores during sintering, and c o m p o s i t i o n a l gradients reflect the s e q u e n c e of hard phase dissolution. N o diffusion of metal atoms occurs during sintering (except in planar faults in cores), so that the material is far f r o m t h e r m o d y n a m i c equilibrium. • C and N m a y diffuse in the hard phases during sintering so that the system a p p r o a c h e s (para)-equilibrium with respect to these elements. A linear relationship b e t w e e n N content in carbonitride phases and total N / ( N + C) in the material was observed. M o increases the activity of N in carbonitride phases. • In c o m p l i c a t e d cermets, Ta m a y to a first a p p r o x i m a t i o n b e treated as Ti and W as Mo. ACKNOWLEDGEMENTS

Prof. B. A r o n s s o n of Sandvik C o r o m a n t is t h a n k e d for introducing us to the exciting field of c e m e n t e d carbides and for e n c o u r a g e m e n t and support. Parts of this w o r k was p e r f o r m e d b y D r A. H e n j e r e d and D r M. Hellsing, f o r m e r m e m b e r s of the a t o m p r o b e g r o u p at Chalmers University. C o l l a b o r a t i o n with D r C. Chaffield and D r B. U h r e n i u s of Sandvik C o r o m a n t and D r H. Jonsson and M r T. M a i n e r t of Seco Tools is gratefully acknowledged. Financial s u p p o r t was received from the Swedish National B o a r d for Industrial and Technical D e v e l o p m e n t ( N U T E K ) , the Swedish R e s e a r c h Council for Engineering Sciences (TFR), A B Sandvik C o r o m a n t and Seco Tools A B .

REFERENCES 1. Exner, H. E., Int. Metals Rev., 4 (1979) 149-73. 2. Fischmeister, H. F., Science of Hard Materials, ed. R. K. Viswanadham, D. J. Rowcliffe & J. Gurland. Plenum, New York, USA, 1983, pp. 1-45.

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3. Doi, H., Proc. 2nd Int. Conf. Science of Hard Materials, ed. E. A. Almond, C. A. Brookes & R. Warren. Inst. Phys. Conf. Ser. No. 75, Adam Hilger, Bristol, UK, 1986, pp. 489-523. 4. Roebuck, B. & Gee, M. G., Proc. 12th Plansee Seminar (Vol. 2), ed. H. Bildstein & H. M. Ortner. Metallwerk Plansee GmbH, Reutte/Tirol, Austria, 1989, pp. 1-29. 5. Henjered, A., Kjellson, L., Andr6n, H.-O. & Nord6n, H., Scripta Metall., 15 (1981) 1023-7. 6. Hellsing, M., Henjered, A., Nord6n, H. & Andr6n, H.-O., Science of Hard Materials, ed. R. K. Viswanadham, D. J. Rowcliffe & J. Gurland. Plenum, New York, USA, 1983, pp. 931-43. 7. Hellsing, M., Karlsson, L., Andr6n, H.-O. & Nord6n, H., J. Phys. E: Sci. Instrum., 18 (1985) 920-5. 8. Andr6n, H.-O., J. Phys. (Paris), 47 (1986) C7-483-8. 9. Rolander, U. & Andr6n, H.-O., J. Phys. (Paris'), 50 (1989) C8-529-34. 10. Lundin, L. & Rolander, U., Appl. Surf. Sci., 1993, 67 (1993) 459-66. 11. Rolander, U. & Andr6n, H.-O., AppL Surf. Sci., 76/77 (1994) 392-402. 12. Rolander, U. & Andr6n, H.-O., J. Phys. (Paris), 49 (1988) C6-299-304. 13. Rolander, U. & Andr6n, H.-O., J. Phys. (Paris), 50 (1989) C8-371-6. 14. Rolander, U. & Andr6n, H.-O., Proc. 12th Int. Plansee Seminar (Vol. 2), ed. H. Bildstein & H. M. Ortner. Metallwerk Plansee GmbH, Reutte/Tirol, Austria, 1989, pp. 379-88. 15. Hellsing, M. & Andr6n, H.-O., Proc. 2nd Int. Conf. Science Hard Mater., ed. E. A. Almond, C. A. Brookes & R. Warren. Inst. Phys. Conf. Set. No. 75, Adam Hilger, Bristol, UK, 1986, pp. 311-20. 16. Hellsing, M., Mater. Sci. Technol., 4 (1988) 824-9. 17. Dawihl, W., Z. Tech. Phys., 21 (1940) 336-45. 18. Gurland, J. & Norton, J. T., J. Metals', 4 (1952) 1051-6. 19. Henjered, A., Hellsing, M., Andr6n, H.-O. & Norddn, H., J. Phys. (Paris), 45 (1984) C9-349-53. 20. Henjered, A., Hellsing, M., Andrdn, H.-O. & Nord6n, H., Proc. 2nd Int. Conf. Science Hard Mater., ed. E. A. Almond, C. A. Brookes & R. Warren. Inst. Phys. Conf. Ser. No. 75, Adam Hilger, Bristol, UK, 1986, pp. 303-9. 21. Henjered, A., Hellsing, M., Andrdn, H.-O. & Norddn, H., Mater. Sci. Technol., 2 (1986) 847-55. 22. Sarin, V. K. & Johannesson, T., On the deformation of WC-Co cemented carbides. Metal Sci., 9 (1975) 472-6. 23. Rolander, U. &Andrdn, H.-O., Mater. Sci. Engng, AI05/ 106 (1988) 283-7. 24. Rolander, U., Chatfield, C. & Andr6n, H.-O., Acta Metall. Mater. (submitted). 25. Lindahl, R, Rolander, U. & Andrdn, H.-O., J. Hard Mater., 3 (1992) 259-67. 26. Lindahl, R, Rolander, U. & Andr6n, H.-O., Atom-probe analysis of the binder phase in TiC-TiN-Mo2C-(Ni,Co) cermets. Refractory Metals and Hard Metals, XX. 27. Lindahl, R, Rolander, U. & Andr6n, H.-O., Surf. Sci., 246 (1991) 319-22.