Cemented Carbide Sintering: Constitutive Relations and Microstructural Evolution

Anders Petersson

Stockholm 2004 Doctoral Dissertation Royal Institute of Technology Department of Materials Science and Engineering Akademisk avhandling som med tillstånd av Kungliga Tekniska Högskolan framläg- ges till offentlig granskning för avläggande av teknologie doktorsexamen onsdagen den 28 april 2004 kl 10.00 i Kollegiesalen, Administrationsbyggnaden, Kungliga Tekniska Högskolan, Valhallavägen 79, Stockholm.

Fakultetsopponent är Professor Colette Allibert, Institut National Polytechnique de Grenoble, Frankrike.

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ISBN 91-7283-739-x ISRN KTH/MSE/--04/21--SE+METO/AVH

c Anders Petersson, March 2004

Högskoletryckeriet, Stockholm 2004 Abstract

Cemented carbides based on tungsten carbide and cobalt are commonly produced by a powder metallurgy route including liquid phase sintering. The pressed compact densifies to almost half its volume during sintering due to pore elimination. The sintering behaviour changes with material composition, such as carbide grain size, binder fraction, carbon content and addition of cubic carbides. This thesis is devoted to the study of constitutive behaviour, in particular densification, and the microstructural evolution during cemented carbide sintering. Dimensional changes are monitored using dilatometry with and without applied external load. The microstructural evolution is investigated with light optical microscopy and scanning electron microscopy. Thermodynamic calculations are used as reference. Constitutive relations are derived for uniaxial viscosity, viscous equivalent of Poisson’s ratio and sintering stress based on relative density and temperature. The relations are extended to a model describing sintering shrinkage with explicit de- pendencies on carbide grain size and binder content. The model is divided in three stages of which two pertain to the solid state and the third to liquid phase sintering. Solid state shrinkage is suppressed in a material with coarse carbides and in the stage of liquid phase sintering grain size strongly influences the uniaxial viscosity. The binder content affects primarily the later densification. The effects of carbon content and grain size distribution on shrinkage have been studied. High carbon content enhances shrinkage rate, but the effect of grain size distribution is rather small. The mean carbide grain size is insufficient to describe densification for very broad distributions only. Shrinkage occurs through rearrangement and solution-reprecipitation. Rear- rangement is studied through the evolution of the pore size distribution and simulated generically using a discrete element method.

Keywords: Cemented carbides, Sintering, Constitutive relations, Microstructure, Densiﬁcation, Modelling

ISBN 91-7283-739-x • ISRN KTH/MSE/--04/21--SE+METO/AVH

iii iv Preface

This thesis is based on the following appended papers that will be referred to in the text by their Roman numerals. I. A. Petersson and J. Ågren, “Constitutive behaviour of WC-Co materials with different grain size sintered under load”, Acta Mater 52(2004):7, 1847–1858 II. A. Petersson and J. Ågren, “Numerical simulation of shape changes during cemented carbide sintering”, In: F.D.S. Marquis (ed.) Powder Materials: Current Research and Industrial Practices III, TMS, Warrendale, 2003, 65– 75 III. A. Petersson and J. Ågren, “Modelling WC-Co sintering shrinkage—Effects of carbide grain size and cobalt content”, submitted manuscript IV. A. Petersson, “Sintering shrinkage of WC-CO and WC-(Ti,W)C-Co materials with different carbon contents”, in manuscript V. A. Petersson and J. Ågren, “Sintering shrinkage of WC-Co materials with bimodal grain size distributions”, submitted manuscript VI. A. Petersson and J. Ågren, “Rearrangement and pore size evolution during WC-Co sintering below the eutectic temperature”, submitted manuscript

VII. A. Petersson, B. Jansson, J. Qvick and J. Zackrisson, “M6C formation during sintering of cemented carbides containing (Ti,W)C”, Int J Refract Met Hard Mat 22(2004), 21–26 The author’s contributions to the co-authored papers, besides being the principal writer, planning most of the experiments and performing all calculations, are participation in experimental work for Paper I, parts of experimental work for Pa- pers II, III and VII and all experimental work for Paper VI. The author performed microscopy for Papers II, V and VI and the main part of computer coding and simulations for Paper VI. The author outlined the main part of analysis and evaluation for Papers I–III and contributed to the evaluation for Papers V–VII.

v vi Contents

Introduction 1

Objectives 3

Liquid phase sintering 5 Surface energy relations ...... 6 Sintering under applied stress ...... 8

Cemented carbide sintering 9 Thermodynamic aspects ...... 10 Binder spreading and rearrangement ...... 11 Solution-reprecipitation and the ﬁnal stage ...... 13 Microstructural evolution ...... 14

Topics of present work 17 Constitutive modelling ...... 17 Outlining the parameters: Paper I ...... 17 Composition and sintering shrinkage: Paper II ...... 19 A shrinkage model for binder content and grain size: Paper III . . 20 Densification ...... 22 Influence of chemical composition: Paper IV ...... 22 Effect of particle size distribution: Paper V ...... 23 Microstructural evolution ...... 24 Rearrangement: Paper VI ...... 24 M6C formation and dissolution: Paper VII ...... 27

Concluding remarks 29

vii viii Introduction

Cemented carbides is a class of composite materials in which a carbide, often tungsten carbide (WC), provides hardness and wear resistance and a metal binder, almost exclusively cobalt (Co), provides toughness [1]. These materials have long been used in applications such as cutting, grinding, drilling and wear-resistant mechanical components. A wide range of material properties can be obtained by changing the material composition. Although the most often used carbide is the hexagonal WC, atlternatives are available. With cubic carbides such as TiC, NbC and TaC the high temperature properties are improved [2]. Minor additions, typically less than a few percent by mass, of for instance Cr2C3 or VC, are often used as grain growth inhibitors [2]. Cemented carbides are commonly produced using a powder metallurgy route. Powders with the proper compositions are mixed and milled with additives, spray- dried to granules and compacted to green compacts. The added pressing aid is removed in a de-waxing step and the compacts sintered to eliminate porosity and increase strength. The microstructural events occuring during sintering are driven by the lowering of excess energy attributed to surfaces and the pressure that develops under a curved interface. The surface area and energy reduction may be accomplished in different ways. One major division in sintering theory is between solid state sintering and liquid phase sintering, respectively, depending on whether or not any liquid phase is present. Densification during solid state sintering is assumed to be due to vacancy flux from pores to grain boundary sinks driven by the vacancy concentration difference [3, 4]. Some possible mechanisms of matter transport are viscous or plastic flow, evaporation and condensation, volume diffusion and surface diffusion [3]. The presence of a small amount of liquid phase is often beneficial for sintering and three stages are indentified in classical liquid phase sintering theory [5], distinguished by the microstructure evolution. Firstly, as the liquid forms, it flows out in the microstructure and causes the solid particles to rearrange into denser packing. Secondly, the particle centres approach and density increases by means of solution and reprecipitation. Thirdly, the solid particles can sinter to a rigid skeleton. Cemented carbide sintering is in some respects in between solid state sintering and liquid phase sintering, but should for two reasons rather be characterised as

1 2 Introduction a liquid phase sintering process. Sintering is in practice almost always performed at temperatures where a quite significant volume fraction liquid exists. However, for finer grained cemented carbides much of the shrinkage, if not the most, occurs during heating to the sintering temperature and prior to liquid formation. This shrinkage shares however most typical features with the first stage of classical liquid phase sintering rather than solid state sintering processes. The particle bonding and increased strength during sintering is often, but not always, accompanied by pore elimination and densification. The term sintering is therefore frequently used synonymously with densification. Since sintering can be performed without shrinkage, or even with swelling, the rate of increasing density will in the following be referred to as densification rate, shrinkage rate or strain rate rather than sintering rate. A vast amount of literature is available in the technically important field of cemented carbides and their sintering. Extensive reviews are given with emphasis on fundamental physical and chemical properties in Reference [1], industrial aspects in Reference [2], mechanical properties in Reference [6], microstructural aspects in Reference [7] and sintering in Reference [8]. Objectives

Present work is intended to improve knowledge of sintering in cemented carbide systems and propose models that can describe shape changes during cemented carbide sintering. The work is founded on two pillars, namely the constitutive behaviour during sintering and the concurrent microstructural evolution. These cannot be separated since it is through the microstructural changes the macroscopical dimen- sions of the compacts change. This is largely an experimental work aiming for a broad description of cemented carbide sintering, but limited to medium–coarse carbide grain sizes. The main tool for studying densification is dilatometry, which is a thermo-analytical (TA) tool, in which temperature and length of the sample are simultaneously recorded. By taking the time derivative of the length measurement it is possible to obtain the instantaneous shrinkage rate that is modelled on a physically sound basis. The microstructural evolution accompanying the shape changes is investigated mainly using scanning electron microscopy (SEM). Thermodynamic constraints are crucial for sintering and thermodynamic calculations are therefore an integral part of the sintering studies. First and foremost, a dense compact is required since retained porosity is detri- mental to most of the mechanical properties [6]. The main focus of the constitutive behaviour is therefore on densification. However, friction between powder and die wall, as well as within the powder, may lead to a heterogenous packing in the green compact. This will in turn lead to heterogenous shrinkage during sintering. Shrinkage calculations from given green densities should therefore be useful in compaction-tool design aiming for near net shaping. Compaction and sintering densification of a WC-Co material were successfully simulated in the works of Brandt [9] and Haglund [10]. The composition of a cemented carbide dictates extensively its shrinkage behaviour and the aim of present work is to extend the constitutive modelling to enable simulations for alloys covering different grain sizes, varying carbon contents and addition of cubic carbides. Due to the complexity of cemented carbide sintering, the aim of a reasonably short-term research cannot be to fully describe the parameters with bearing on densification. That would be a tremendous task. This work rather strives to isolate the most crucial aspects of sintering and characterise their variaton with main compositional parameters. Furthermore, experimental information on some of the

3 4 Objectives key parameters such as interfacial energies and contact angles with their variations with sintering conditions are still largely lacking. This necessitates a fair amount of parameter fitting to study sintering of a given alloy once constitutive models are obtained. To predict, within narrow tolerances, the outcome of sintering one will probably still have to perform a number of experiments, but it is in the scope of present work to limit the number of such experiments by giving a sense of directions. The ability to simulate processes and to predict properties, and ultimately to virtually design materials, give promises for a faster and more cost efficient materials development. The desire for sintering models that can be used to simulate sintering via computer calculations is founded in this context. With such simulations it should be possible to explore effects of various compositional aspects and process parameters under different sintering conditions. Liquid phase sintering

During sintering the excess surface energy attributed to porosity is decreased by means of matter transport that becomes kinetically possible at high temperature. Thus, sintering deals with a central issue of materials science, i.e. how much is the total energy of the system lowered and in what way. When penetrating the process in order to understand or predict its course, one inevitably finds several possible reaction steps, making sintering a quite complicated process. Not only is the total energy reduced by pore elimination, but also by microstructural coarsening and in many cases through phase transformations. Under such circumstances the description of sintering is further complicated. For complete densification by liquid phase sintering three things are usually required, namely (i) an appreciable amount of liquid, (ii) solubility of the solid in the liquid and (iii) complete wetting [11–15]. Moreover, densification in a typical liquid phase sintering system is normally divided in three stages: (i) rearrangement, (ii) solution-reprecipitation and (iii) solid phase sintering based on the respective dominant process. When the liquid phase forms it wets the solid particles resulting in capillary forces rearranging the particles into a denser packing. The powder compact re- sponds to the capillary pressure as a viscous solid and viscosity increases with decreasing porosity [5]. Rearrangement is retarded by friction between solid particles [16], but is usually so fast that all other events are overshadowed and full density can be reached by rearrangement if around 35 volume percent liquid is present [5]. In the second stage densification occurs with shape accomodation of particles separated by liquid films through solution and reprecipitation of solids leading to closer packing [16]. Material is dissolved at solid-liquid interfaces with high chemical potential and reprecipitated at other sites with lower chemical potential. A number of solution-reprecipitation mechanisms have been identified, e.g. Ostwald ripening, Ostwald ripening with shape accomodation, contact flattening, particle disintegration, coalescence and pore elimination by liquid flow [17]. The solubility of the solid in the liquid is inversely related to the particle size. The resulting concentration gradient causes material flow from finer particles to coarser particles. Thus, the second stage is usually marked by microstructural coarsening [5]. One possible process is illustrated with the initial configuration of two particles

5 6 Liquid phase sintering separated by a liquid film in Figure 1 and after some material from the particle contact has dissolved in the liquid and precipitated elsewhere leading to a denser configuration in Figure 2. In the case of equally sized spherical particles the stress gradient driving dissolution in the contact zone stems from the capillary pressure provided by the binder. Both figures are redrawn from Reference [11].

R δ Liquid

Figure 1. Two particle held together by capillary pressure. After Reference [11].

Figure 2. Centre-to-centre approach due to solution-reprecipitaion. After Refer- ence [11].

In the third stage the solid phase may sinter into a rigid skeleton. The densiﬁ- cation is then described with typical solid state sintering models [16].

Surface energy relations

Surface energy is the driving force for the events that leads to a minimum-energy conﬁguration of the microstructure during liquid phase sintering [5]. Wetting is important for the sinterability of a powder system and characterised by the contact angle, given by the solid-liquid-vapour equilibrium. This is demonstrated in Figure 3 drawn after Reference [5]. Surface energy relations 7

vapour γ liquid LV θ θ γ γSL SV

solid

Figure 3. Interfacial energies and the deﬁnition of a contact angle. After Refer- ence [5].

Interfacial energies and the wetting angle are related by means of Young’s equation [5]

γSV = γSL + γLV cos θ, (1) where the respective interfacial energies are solid-vapour, γSV , solid-liquid, γSL, and liquid-vapour, γLV . The contact angle is denoted θ. Good wetting is characterised by a small contact angle and conversely poor wetting by a large contact angle. Wetting is promoted by solubility of the solid in the liquid, formation of intermediate compounds and interdiffusion [5]. If the sum of the liquid-vapour and solid-liquid interfacial energies is lower than the solid-vapour interfacial energy the net interfacial energy is reduced by spreading the liquid over the solid surface as demonstrated in Figure 4 after Reference [5]. Chemical affinity and solubility between the phases generally contribute to efficient spreading [5].

γ -γ +γ < 0 LV SV SL

vapour liquid

solid

Figure 4. Spreading of a liquid ﬁlm over a solid surface driven by a reduction of net interfacial energy. After Reference [5].

If a volume element dV is removed or added from a surface of principal radii r1 and r2 the energy change becomes [16]

dE dA 1 1 = γ = γ + , (2) dV dV r1 r2 8 Liquid phase sintering where γ is the surface energy and dA the change in surface area. Identifying the term dE/dV as a stress σ, the Laplace equation

1 1 σ = γ + (3) r1 r2 is obtained. If the surface is convex the stress is tensile and if the surface is concave the stress is compressive. It is this stress gradient that drives material transport during sintering [16].

Sintering under applied stress

It is self evident that densification during sintering is enhanced by an applied pressure if it is sufficiently large. On a micro-level such a pressure enhances not only rearrangement, but also solution at particle contacts and plastic flow in the solid particles [18]. By applying an external pressure less sinterable powders can be sintered to high density and strength [16]. The applied pressure is often modelled as additive to the capillary pressure [5, 18]. For an extensive review on sintering modelling, including application of load, see Reference [19]. Cemented carbide sintering

In industrial practice cemented carbides are sintered by liquid phase sintering at temperatures in the range 1350–1650◦C leading to virtually pore-free compacts [2]. The foundation of the process in terms of thermodynamic and kinetic aspects of the alloy systems has been studied extensively, as well as their interaction in the developing microstructures, see for instance References [1, 2, 6–8] mentioned in the introduction. Freytag and Exner [20] studied sintering through changes in the magnetic properties of a WC-12 mass-% Co material. It was found that below 800◦C practically no shrinkage is observed but stresses introduced by milling are relieved and cobalt oxide on surfaces are reduced. In the range 800–1100◦C shrinkage is limited, tungsten carbide partially dissolves in the binder that spreads over the carbide surfaces. Further heating, but still below the temperature of liquid formation, results in a large portion of the total shrinkage. It is stressed that enhancement of diffusion in the cobalt-saturated carbide grain boundaries should have a decisive role in solid state sintering. When the binder melts a substantial amount of carbide is dissolved in the liquid. Rearrangement and pore filling eventually lead to full density. Surface oxides, practically always present in a powder, are reduced in a wide temperature range 650–850◦C with a maximum around 750◦C with an additional ◦ peak in CO/CO2 emission at 950 C for doped WC materials [21]. Freytag and Exner [20] relate the main carbon loss during sintering to reaction of oxygen at the carbide surfaces in the temperature range 700–1000◦C. Dilatometry coupled with thermogravimetry, calorimetry and mass spectrome- try has further shown that the shrinkage onset occurs after formation of CO and CO2 from adhesive oxygen and oxides with carbon from WC, which leads to the conclusion that shrinkage starts when all surface oxides are reduced [22]. This was however recently questioned by de Macedo et al. [23] who studied wetting of Co, Ni and Fe on W and WC. It was found that wetting occurs also on oxidised surfaces, and it is concluded that the onset of shrinkage is not related to oxide reduction although they appear in sequence. Göthelid et al. [24] studied Co spreading on WC substrates with scanning tunneling microscopy (STM). It was found that at first the substrate is partially covered by a so called 2×2 structure on top of which Co then spreads rather easily. Clusters of cobalt oxide form in the presence of oxygen rather than the 2×2 structure. It appears as spreading is impeded until oxygen is

9 10 Cemented carbide sintering removed and the 2×2 structure restored. This was found to occur at a temperature of 800◦C, which is the temperature where shrinkage typically starts during WC- Co sintering. Although sintering conditions diﬀer substantially from the ultra-high vacuum employed in STM studies this points to a correlation between an existing oxide ﬁlm and poor wetting conditions.

Thermodynamic aspects

The simplest cemented carbide and the base for most alloys is the mixture of WC and Co. An isothermal section of the ternary phase diagram Co-W-C at 1350◦C, calculated with the Thermo-Calc software [25] and a database for cemented carbides [26], is shown in Figure 5. Compositions are given in atomic fractions. It is apparent that the carbon content is important for the material constitution and for only small deviations additional phases will form. For low carbon contents the M6C may become stable and for high carbon contents graphite. Neither one is usually desired in the ﬁnished product. It is also demonstrated that the two-phase regions {fcc-Co+WC} and {liq-Co+WC} narrow with decreasing Co content. Within the narrow compositional range for the two-phase ﬁeld the carbon activity increases from 0.3 to 1 with respect to graphite.

graphite 1.0

0.8

0.6

C WC x 0.4 W2C

0.2 M C liq-Co 6 M12C

0 Co W bcc-W fcc-Co 0 0.2 0.4 7 60.6 0.8 1.0

Figure 5. Isothermal section of the ternary phase diagram Co-W-C at 1350◦C.

The narrow range of carbon content permissible if only WC and Co are desired in the sintered structure is further demonstrated in Figure 6. This is a calculated vertical section through the Co-W-C system at 10 mass-% Co. Here the stable Binder spreading and rearrangement 11 phases can be viewed as function of temperature. A stoichiometric mixture of WC- 10 mass-% Co corresponds to 5.52 mass-% C. The two-phase regions {fcc-Co+WC} and {liquid Co+WC} are limited to the carbon range 5.38–5.54 mass-%. Since some carbon inevitably reacts with oxygen during sintering it is in practice often necessary to add carbon to the WC-Co mixture to obtain the desired composition of the ﬁnished product [2].

1500 WC liquid Co 1400 C °

1300

temperature, graphite M6C WC 1200 WC WC fcc-Co fcc-Co fcc-Co

1100 5.2 5.4 5.6 5.8 carbon content, mass-%

Figure 6. Vertical section of the ternary phase diagram Co-W-C at 10 mass-% Co.

With the addition of other carbides such as TiC, NbC, TaC, VC and Cr2C3 the cobalt-rich liquid becomes stable at lower temperature [26]. It is believed that conditions close to equilibrium prevail at temperatures typically above 1000◦C owing to the short diﬀusion distances involved in cemented carbide sintering [27]. It should perhaps be stated that porosity present in the compact is not ther- modynamically stable. Porosity contributes excess energy that is the driving force for sintering. Indeed, all interfaces in the compact contribute to the total energy. By reducing the interfacial area the total energy is lowered. This is the cause of coarsening when small constituents shrink on behalf of larger ones that grow.

Binder spreading and rearrangement

Densification during cemented carbide sintering occurs initially through rearrangement followed by solution-reprecipitation [28,29]. With higher amount of liquid and smaller carbide particle size most densification is achieved rapidly [28] and the high 12 Cemented carbide sintering density obtained in the very first stage of sintering even before the liquid is formed distinguishes the WC-Co system from other liquid phase sintering systems [29]. Froschauer and Fulrath [30] studied the microstructural evolution in situ and confirmed that rearrangement is the main densification process. Haglund and co- workers [31] showed for a WC-10 mass-% Co material that densification in solid state can be treated as a rearrangement process like in the first stage of liquid phase sintering. No drastic change in densification rate occurred as the binder melted. During ball milling of the powder the carbide particles become partially covered by cobalt and small carbide fragments are embedded in Co, which is important for the initial sintering [32]. Milling and cold pressing also create highly deformed contact areas with high internal stresses that may deform by creep or other transport mechanisms at elevated temperature relaxing the defect structure [21, 33]. Snow- ball and Milner [34] showed that co-milled WC-Co densifies to a higher extent than WC-Co milled separately and mixed. Densification started locally in cobalt-rich areas packed with carbides from milling. At temperatures above one half of the melting temperature, 650–710◦C for fine grained cemented carbides, the dispersed binder particles behave like a viscous mass spreading by wetting onto the surrounding carbide surfaces [33,35–37]. The spreading of the interfacial zones immediately after surface oxide reduction constitutes the first stage of densification [21]. The surface free energy is lowered by the cobalt spreading [38] and exchanging WC-gas and Co-gas interfaces for WC-Co lowers the interfacial energy to approximately one fifth [35]. Solid state spreading of cobalt is driven by the surface energy relations in a similar way as desribed by the Young’s equation, Equation 1, although the contact angle between carbide and binder is more complex to define. Very high spreading rates are reported for Co on WC substrates and it was observed that it is only a film that spreads over the surface and not the bulk metal. The mechanism behind binder spreading is believed to be viscous flow of a defect-rich contact zone [23]. Meredith and Milner [36] suggested that the viscosity in the carbide-metal interface is further reduced by carbide dissolution owing to enhanced diffusion. The rate of binder spreading is limited by factors such as area, curvature and chemistry of the carbide particles and content, dispersion, chemistry and physical properties of the binder. Binder spreading is connected to the Laplace forces acting along the wetting front line between carbide and binder. These microscopic forces rearrange the carbide particles and reduce the mean distance between neighboured particles, resulting in densification [35]. The smaller the grain size is the more interfacial area is contained within the compact and the higher is the stored interfacial energy. As a result sintering starts earlier for fine-grained cemented carbides and the share of solid state shrinkage increases [35]. Moreover, milling and mixing introduce more defects for finer WC materials, which results in higher internal stress, shorter distance between source and sinks of diffusion and thus lower viscosity [39]. The maximum in shrinkage rate Solution-reprecipitation and the final stage 13 occurs at lower temperature and the shrinkage rate curve is wider but flatter [21,22]. Densification behaviour also becomes increasingly sensitive to processing steps prior to sintering. For ultrafine cemented carbides the solid state sintering can contribute with 90 percent of the total densification [35]. Earlier shrinkage is observed for attritor-milled powders compared to balled-milled powders [37].

Grain growth inhibitors such as VC and Cr2C3 in small amounts shift shrinkage onset to higher temperature, decrease the solid state shrinkage rate and lower the melting point of the binder [21, 35]. Addition of either carbon or tungsten to a fine-grained WC-Co material decreases the shrinkage rate with stronger effect of tungsten. The influence on a coarse-grained material is smaller [20]. Except for the ultrafine materials solid state shrinkage can kinetically be divided in two stages. The first stage of densification, typically below 1100–1200◦C, is believed to be homogeneous shrinkage due to binder spreading and creep, while in the second stage unsymmetical forces become strong enough to introduce twist, rotation and shift among particles, although still governed by binder creep [21]. In the first stage intrinsic stresses of the binder contribute to the capillary force in driving densification, while in the second stage the binder is annealed and its intrinsic stresses neglible compared to the capillary stress. Moreover, in the first stage the binder is nearly free to creep, while in the second stage creep is restricted by the large number of carbide-binder interfaces [35]. Densification in the first stage depends strongly on processing prior to sintering and only weakly on grain size, while in the second stage a weak dependence on processing but strong dependence on grain size are observed. The solid state shrinkage appears more complex for ultrafine materials for which it is not possible to divide it in only two stages [35].

Solution-reprecipitation and the ﬁnal stage

Densification in the intermediate stage of liquid phase sintering is accomplished through grain shape accomodation and for cemented carbides with contributions from contact flattening where carbide particles are in contact, dissolution of small grains and reprecipitation on large grains and by coalescence via grain boundary migration and solution-reprecipitation [29]. Microstructural coarsening leads to a prismatic grain shape. However, the prismatic shape is not necessarily compatible with the aggregated structure formed in the initial stage. With a sufficient amount of binder this may lead to additional rearrangement [36]. Practically full density is normally reached for most WC-Co materials by rearrangement and shape accomodation. Sintering of the solid phase skeleton, i.e. the last stage of classical liquid phase sintering is usually not observed in cemented carbide sintering. Other microstructural changes may however occur, for instance grain size and grain size distribution, grain shape and binder phase distribution may all change [29]. 14 Cemented carbide sintering

Since full density is reached through continuous repacking of clusters leaving cobalt free to flow further into the structure Snowball and Milner [34] argue that it is not necessary to postulate a third stage of shrinkage, but with low cobalt content a slow final densification pertaining to solid state sintering of the carbides may be observed [38]. The actual existence of a continuous solid skeleton in WC-Co has been debated. For instance, Warren and Waldron [40] suggested that there is a skeleton during and after sintering, while Nelson and Milner [38] claimed that there is a cobalt film between carbide particles for short sintering times. This would allow for particle sliding during densification, but a rigid carbide network will eventually form during prolonged sintering. Gurland and Norton [14] ruled out the possibilty of a carbide network based on infiltration experiments. Andrén [41] refers to measurements with atom probe field ion microscopy (APFIM) and analytical electron microscopy (AEM) showing approximately half a monolayer Co is present between carbide grains. Accordingly, there is a continous carbide skeleton, but half the contact surface is covered by Co.

Microstructural evolution

Many key properties of cemented carbides are determined by the microstructure. Carbide grain size and shape, binder fraction, composition and distribution in addition to retained porosity are a few of the more important microstructure characteristics that influence the properties such as transverse rupture strength and hardness [6]. Primary interest here is, however, the microstructural evolution concurrent with densification. Along with dimensional change, i.e. shrinkage and creep deformation, the microstructure changes too. If it was not for the microstructural change, compact shape would not actually change since it is through the rearrangement of carbide particles and grain shape accomodation that densification occurs. The carbides are rearranged from a rather homogenous spatial distribution in the green body to an aggregated structure during the initial stage of shrinkage, and Meredith and Milner [36] attributed the rapid initial densification to the formation of close-packed boundaries between carbides when cobalt diffuses into the carbide interfaces and bringing the particles closer. In a second stage of densification the binder bridges can rearrange the carbide particles [21]. During the intense shrinkage in solid state WC-Co clusters that have formed upon cobalt spreading behave like a viscous mass, and the clusters creeps into the surrounding void area. Smaller Co particle size and better dispersion leads to more WC-Co clusters formed during binder spreading, shorter diffusion distances and more uniform capillary forces. The result is a more homogeneous void evolution that will influence the later stages of densification [37]. The early rearrangement into locally dense areas occurs at the expense of cre- ating micropores due to inhomogeneities in the green compact [35]. With binder Microstructural evolution 15

flow into the surrounding matrix more and more carbide particles are cemented in a growing agglomerate [34]. Later densification proceeds by creep of the WC-Co material to fill the micropores [35]. Spreading of Co along carbide surfaces and rearrangement into an aggregated structure is shown in Figure 7 which was obtained in SEM with secondary electron mode after heating a WC-Co material with 3 µm WC and 10 mass-% Co to 1280◦C and cooling.

Figure 7. Micrograph obtained after heating a WC-Co material to 1280◦C and cooling. WC appears light grey, Co dark grey and pores black.

The carbide particles are typically rounded after milling the starting powders. The wide extent of the solution-reprecipitation process during sintering is evidenced by the prismatic grain shape typical for WC-Co materials after prolonged sintering [36]. The preference for prismatic growth stems from the WC crystal structure and interfacial energy anisotropy [1, 42]. Prismatic carbide shapes develop faster in ﬁne-grained materials than coarse-grained materials [38] and the shape of the facetted WC grain changes with the carbon content [42–44]. The carbide facetting is evident in Figure 8 which was obtained in SEM with secondary electron mode after sintering the same material as in Figure 7 at 1430◦C for 60 min. This sintering cycle is similar to the one used in industrial practice. The grain size of the sintered structure basically depends on the initial size distribution and thus the milling conditions [45]. Large grains may act as starting point for rapid discontinuous grain growth widening the grain size distribution. One trend in more recent materials development is towards ﬁner carbide grain sizes, since this improves hardness and wear resistance [46]. This has naturally led to studies devoted to sintering mechanisms in submicron and nanocrystalline 16 Cemented carbide sintering

Figure 8. Micrograph obtained after sintering a WC-Co material at 1430◦C for 60 min. WC appears light grey, Co dark grey and pores black. cemented carbide materials. Since this study is limited to medium–coarse grain sizes the reader is referred to for instance References [35,47, 48]. Topics of present work

Constitutive modelling

Outlining the parameters: Paper I Ågren and co-workers [49,50] devised a model for the simulation of shape changes. The model was based on rearrangement of the carbide particles under capillary action of the binder. Rearrangement is driven, in the absence of an applied stress, by the sintering stress inversely proportional to grain size and linearly propotional to the pore fraction. Binder spreading was considered practically instantaneous and diffusional creep of the binder rate limiting. A linear combination of surface diffusion and volume diffusion was introduced. Since a strong influence of heating rate on shrinkage rate had been observed [31] the effect of heating was taken into account explicitly. The first attempt to extend this model to describe shrinkage of materials with different grain sizes failed, and the study reported in Paper I was initiated. A slightly different approach, used for instance by Gillia et al. [51] for WC-Co, was chosen in which the uniaxial viscosity and the viscous equivalent of Poisson’s ratio are derived from sintering under applied axial load. The sintering stress can be calculated from these two parameters and the shrinkage rate under free-sintering conditions, i.e. without the application of load. The evaluation of the constitutive parameters followed in much the approach by Bordia and Scherer [52, 53]. The constituive parameters were modelled as functions of temperature and relative density assuming that the material is Newtonian, i.e. the strain rate is linearly related to stress. No direct relation to heating rate was assumed. The calculated uniaxial viscosity showed three stages as demonstrated in Fig- ure 9. Of these are the first two before liquid formation. The third stage is during liquid phase sintering and marked by an initially decreasing viscosity. Note then the drastic increase as the materials approaches full density during liquid phase sintering and the relatively large difference towards the end of sintering for fine (1.7 µm) and coarse (5.4 µm) carbide grain sizes. Solid state sintering was modelled with diffusional viscosity similarly as Her- ring [54]. The apparent activation energies in the two stages differed and were

17 18 Topics of present work roughly twice as high in the second stage. For liquid phase sintering a little different approach was chosen than usual. With the analogy of free volume in glass transition and percolation theory [55,56] a description of rearrangement facilitated by the present porosity was suggested. Combining the deformation pertaining to the dense material and the rearrangement an eﬀective viscosity was obtained with the desired properties of a fast increase with relative density, but still allowing a ﬁnite deformation rate of the dense material. It should be pointed out that the suggested model perhaps does not give the easiest numerical treatment. If the viscosity of the dense material is neglected an exponential function describing the proper increase with relative density should be suitable.

6 1400 10

1300

5 10 s 1200 fine C °

4 1100 10

coarse temperature, 1000 uniaxial viscosity, MPa 3 10

900

800 0 500 1000 1500 2000 2500 time, s

Figure 9. Uniaxial viscosity calculated from sintering under axial load.

The experimentally determined viscous Poisson’s ratio was unfortunately bur- dened with large scatter. Two modelling approaches were tested. In addition to an exponential relation, as used by Gillia et al. [51], Scherer’s model [57]

1r ρ ν ≈ , (4) p 2 3 − 2ρ was used. This was derived for a particular pore arrangement for a different material, but captured the essence of derived data in Paper 2. Moreover, it has the benefits of a rather simple relation to relative density as sole parameter and the correct limits of νp → 0 for ρ → 0 and νp → 0.5 for ρ → 1. Although a satisfactory representation is found with this model it is naturally a source of uncertainty in the analysis. The sintering stress is the thermodynamic force associated with sintering. It can be viewed as the collective action of interfacial tension throughout the compact. One common definition seemingly derives from the work of Gregg and Rhines [58] who measured the tensile force that was necessary to counteract shrinkage. In this Constitutive modelling 19 view the sintering stress is a pressure of equal magnitude as the hydrostatic tensile stress that would stop shrinkage. The calculation of sintering stress in Paper I showed a value slightly increasing with decreasing porosity, which contrasts the assumptions in References [49, 50]. The sintering stress was described by a power function of relative density. In Paper I the influence of an applied stress on densification was studied as well. It is a commonly held opinion that the hydrostatic component of an applied stress enhances densification [5, 16]. Though the applied uniaxial load increases axial shrinkage and decreases radial shrinkage the total densification is expected to increase. The first series of experiments with axial stresses up to 0.7 MPa influenced densification rate only marginally, and a second series was performed under much higher stress of up to 6 MPa for which a marked effect on densification was found. Under all sintering conditions a densification-rate maximum was found near a relative density of 0.8 such as shown in Figure 10, although temperature differed. 1 N is believed to be close to free sintering. 15 N load resulted in a maximum stress of 0.7 MPa and 150 N in 6 MPa at most. The carbide grain size was 1.7 µm. Similar diagrams obtained under free sintering for a variety of compositions together with calculated solidus and liquidus temperatures indicate that the first local maximum is related to solid state shrinkage and the second to liquid phase shrinkage.

-3 x 10 1.2

0.8 1 N

15 N

0.6 150 N densification rate 0.4

0.2

0 0.5 0.6 0.7 0.8 0.9 1 relative density

Figure 10. Densiﬁcation rate versus relative density for three levels of uniaxial load.

Composition and sintering shrinkage: Paper II The general constitutive behaviour was outlined in Paper I, although some parameters scattered. Moreover, a model of more practical interest should account explicitly for parameteters of the material composition. The aim of Paper II was therefore twofold. Firstly, the suggested models were used to describe shrinkage 20 Topics of present work data from dilatometry where experimental certainty normally is rather high. Sec- ondly, results from sintering a number of different materials should point to the general influence of respective parameter on shrinkage. Microstructures were exam- ined at various stages of sintering before and after liquid formation using scanning electron microscopy. The microstructural evolution was found generally coherent with findings in for instance References [31,36,59], with the main contribution to densification from rearrangement, although leading to clustering, and subsequent solution-reprecipitation. Different carbide grain sizes were shown to still result in similar temperatures of shrinkage onset. However, shrinkage rate during solid state sintering decreases with increasing carbide grain size, and for a material with approximately 10 µm WC the main solid state densification was suppressed. The latter can be linked to the microstructural evolution showing that although the binder spreading onto the neighbouring carbide particles is fast the continuous linking of more carbides into an agglomerate is impeded. Following the decrease in shrinkage rate during solid state sintering the relative amount of shrinkage during liquid phase sintering increases with carbide grain size. With a medium carbide grain size, typically 2 µm, only a fraction of the densification is due to liquid phase sintering. For coarse materials, with carbide grain sizes around 5 µm, densification before and after liquid formation are practically equal. In extra-coarse materials, up to 10 µm WC, only a fraction of the densification is caused by solid state sintering. For the terminology of carbide grain sizes, see for instance the discussion by Roebuck [60]. Cobalt content mainly influences the later stages of shrinkage from a relative density of approximately 0.7 where shrinkage rate increases with cobalt content, although microstructural evolution occured in similar ways. It was shown that the constitutive models with adjustable parameters outlined in Paper I can be combined to successfully describe free sintering shrinkage rate in three stages as shown in Figure 11 for a WC-Co material with 2 µm WC and 10 mass-% Co. At the transition from the first stage to the second stage the modelled curves overlap and cannot be distinguished. It was also shown how these models can be used to calculate sintering shrinkage by numerical integration for a given sintering cycle.

A shrinkage model for binder content and grain size: Paper III The constitutive relations outlined in Paper I were shown suitable for describing shrinkage in Paper II. With phenomenological relations for the eﬀects of cobalt content and carbide grain size on the respective parameters uniaxial viscosity, viscous equivalent of Poisson’s ration and sintering stress the three stages of shrinkage can be described with no less than 13 adjustable parameters. In Paper III these 13 parameters were reduced to only four by ﬁtting model parameters to dilatometry data for three cobalt contents and three carbide grain sizes. The remaining four parameters are enough to reasonably well represent shrinkage for other materials. Constitutive modelling 21

-4 x 10 0

-0.5

-1

-1.5

free sintering strain rate, 1/s -2

-2.5

0 1000 2000 3000 4000 5000 time, s

Figure 11. Shrinkage rate (—) measured and (---) described by the shrinakage model. After Paper II.

The obtained model of linear shrinkage with adjusted parameters reads

Aρ2.1d0.063 exp (−Q /RT ) (1 − 2ν ) ε˙ = − 1 p (5) 1 RT 2.1 −0.035 Bρ d exp (−Q2/RT ) (1 − 2νp) 1 − fWC ε˙2 = − (6) RT 1 + 3fWC 4.1 0.69 0.96 exp (−0.096/ (1 − ρ)) ε˙3 = −Cρ fCo (1 − 2νp) 1 + 0.45 0.18 (7) d fCo for the respective stages. ρ is the relative density, R the gas constant, T absolute temperature, νp the viscous equivalent of Poisson’s ratio described by Equation 4, d carbide grain size, k the Boltzmann constant, Q2 apparent activation energy of typically 250 kJ/mol, fWC volume fraction carbide and fCo volume fraction cobalt in the dense material. The change from the first to the second stage occured around 1100◦C. In the solidus-liquidus interval the shrinkage rate is expected to be a combination of the solid fraction of the material shrinking according to Equation 6 and the liquid fraction shrinking according to Equation 7. Two adjustable parameters, A and Q1 refer to the first stage of shrinkage. The apparent activation energy Q1 was found in the range 60–140 kJ/mol. In this stage of sintering Gille et al. [35] observed activation energies of densification ranging from 50 kJ/mol for attritor-milled powders to 100 kJ/mol for ball-milled powders. Apparent activation energies of 70 kJ/mol and 120 kJ/mol was found for uniaxial viscosity in Paper I. All these are about the same order of magnitude. The enhanced 22 Topics of present work diffusion of Co in the WC-Co interfaces discussed in the literature [20, 36] may be responsible for the very low activation energies observed. In this stage of sintering deformation is most likely governed by surface diffusion of Co over WC surfaces. In the second stage all materials showed practically the same apparent activation energy of 250 kJ/mol in Reference [35], which was also observed for most materials in Papers I and only a little higher for uniaxial viscosity in Paper III. This value is typical for volume diffusion of cobalt [61], and it is reasonable to assume that deformation is governed by diffusional creep pertaining to volume diffusion of the cobalt-rich binder. Missiaen and Roure [62] analysed WC-Co shrinkage with a morphological description of shrinkage. They conclude that densification and binder spreading are coupled and that volume diffusion in the cobalt phase is the governing mechanism. As discussed in Paper III an accurate model of densification would preferably describe shrinkage for a material with only a few, if any, adjustable parameters. It was also pointed out, as well as in Reference [35], that the early shrinkage is sensitive to processing prior to sintering, and accordingly the first stage of shrinkage differs between different materials, which then must be reflected by model parameters. Moreover, Gille et al. [22] investigated a new ultrafine WC material where the carbide had a relatively small lattice distortion compared to other ultrafine grades. Given the lower solution pressure for the less distorted lattice the solution and reprecipitation was slower resulting in a more stable sintering, less coarsening and a more homogenous microstructure [22]. Accordingly, processing influences also the later stages of densification and it seems difficult to realise one common description of sintering shrinkage for WC-Co materials. Though the suggested shrinkage models in Paper III use the mean grain size, microstructural coarsening can be accounted for in simulations by solving the equa- tions of shrinkage and grain growth simultaneously, should the latter be available.

Densiﬁcation

Inﬂuence of chemical composition: Paper IV

Paper IV aimed at investigating the influence of chemical composition on shrinkage. Three WC-Co and and three WC-(Ti,W)C-Co materials with different carbon contents were investigated. Additions of the cubic carbide (Ti,W)C delays shrinkage to higher temperature and high carbon content above the limit of graphite formation enhances shrinkage for both WC-Co and WC-(Ti,W)C-Co materials. A low carbon content towards the limit of M6C formation delays shrinkage for a material with (Ti,W)C but not for WC-Co. Densification 23

The densification model of Paper III with four adjustable parameters described shrinkage reasonably well for WC-Co materials with low and medium carbon content and for WC-(Ti,W)C-Co with high carbon content. The model failed to cap- ture the enhanced densification in WC-Co with high carbon content and the solid state shrinkage of WC-(Ti,W)C-Co with low and medium carbon contents. For WC-(Ti,W)C-Co the three stages of shrinkage is reduced to two stages where the first stage of lower activation energy in WC-Co is absent. The observed variation of shrinkage rate with temperature points to a much higher apparent activation energy than the 250 kJ/mol observed in Papers I and III and Reference [35]. Freytag and Exner [20] related the cobalt spreading and thereby increasing specific surface in the initial stage of WC-Co sintering to the change in coercive force occuring in the temperature range 800–1100◦C. Since continued heating does not change the magnetic properties much [20], it may be concluded that cobalt spreading has been completed at this stage. At this temperature densification rate starts to increase considerably and here was also the transition from the first stage of shrinkage to the second stage found in Paper III and Reference [35]. However, with a significant amount of WC exchanged for the mixed cubic carbide (Ti,W)C almost no shrinkage is yet observed. Although de Macedo et al. [23] recently reported that spreading can occur also on an oxidised substrate this is in opposition to the prevailing view that wetting is significantly poorer on oxides than on carbides. It is plausible that relatively stable surface oxides present on the carbide surfaces effectively prevent binder spreading. With an oxide stability increasing with decreasing carbon content oxide reduction and onset of shrinkage would then be delayed to higher temperature as observed. This leads to the suggestion that the two stages of surface diffusion and volume diffusion are sequential and not simultaneous. It was observed by de Macedo et al. [23] that Co spreads on a WC as a very thin layer, and by Göthelid et al. [24] that a pre-cursor layer of Co must form on top of the WC surface before spreading can occur. It seems reasonable that surface diffusion in the first stage of shrinkage forms a Co film over the carbide with spreading completed around 1100◦C. This leads also to a little rearrangement and densification. When temperature is high enough for volume diffusion to dominate further rearrangement occurs through binder creep. If, like it was suggested for WC-(Ti,W)C-Co, the surface diffusion is limited in the presence of oxides also the volume diffusion is hindered since the necessary film is lacking. Thus, shrinkage behaviour of the two less sinterable types of materials studied, WC-Co with coarse carbide and WC-(Ti,W)C-Co, are different. In the former case the second stage is suppressed as discussed above and in the latter case the first stage.

Effect of particle size distribution: Paper V This work was initiated to investigate how the width of a particle size distribution affects shrinkage and if the mean carbide grain size is sufficient in models such as those devised in Papers I–III. 24 Topics of present work

Shrinkage of materials with different unimodal and bimodal particle size distributions were compared. It was found that the mean grain size, indeed, is sufficient to describe shrinkage for a bimodal distribution with initially 75 percent fine particles and 25 percent coarse by mass. With 50 percent each of fine and coarse particles, however, shrinkage was initially dominated by the fine fraction and sub- sequently by the coarse. The extent of solution-reprecipitation during liquid phase sintering was clearly demonstrated. With bimodal particle size distribution an extreme coarsening was observed during liquid phase sintering in addition to marked facetting. A unimodal particle size distribution showed much less coarsening but still exclusively facetted carbides.

Microstructural evolution

Rearrangement: Paper VI

This study was started in light of the observation by Haglund [10] that the average pore size in WC-Co seems to shrink during heating but grow during isothermal holding although macroscopic shrinkage is observed. The experimental basis for the work was measurements of the pore size distribution for various stages of a heating cycle with and without imposed isothermal holding. Sintering was performed below the temperature of liquid formation. The pore size distribution, if plotted with number frequencies on a logarithmic scale changed in a nearly self-similar way as indicated by Figure 12 where Gaussian distribution curves have been adjusted to measured data for the ease of comparison. With isothermal holding time bimodal pore size distributions developed as shown in Figure 13 for the holding times 60, 180 and 480 min. Isothermal holding after slow heating resulted in a pore size distribution with less pronounced bimodal character than after fast heating. The formation of bimodal pore size distributions was viewed in light of particle rearrangement into clusters where pores inside clusters shrink and pores between clusters grow. The direct cause of clustering in liquid phase sintering is that the inter-particle attraction force increases with decreasing particle distance, which enhances irregularities in particle packing [63]. The formation of clusters was veriﬁed in simulations using a Discrete Element Method (DEM), where the capillary attraction between particles connected by a liquid bridge was modelled. The particle attraction forces, contact forces and viscous drag forces were used to solve the equation of motion and simulate the microstructural evolution during rearrangement. The absence of clustering during heating was interpreted as a result of diﬀer- ent particle attraction attributed to the carbide solubility in the binder. With increasing temperature the solubility increases monotonously, while at isothermal conditions the equilibrium solute content is rather soon established and a more rigid Microstructural evolution 25

4 x 10 5

4 1150oC ) -2 1200oC 3

1250oC

2 Number frequency (mm 1

0 -2 -1 0 1 logarithmic pore size

Figure 12. Pore size distributions measured after heating. structure forms. Moreover, DEM simulations of isothermal holding following heating showed that the microstructural evolution retards considerably simply by the realisation of constant viscosity. The DEM simulations failed however to account for solubility effects leading to significant densification. There has been a question about a possible influence of heating rate on shrinkage rate. Haglund et al. [31] noted that shrinkage rate for WC-Co decreases drastically as isothermal conditions are reached. This has also been observed by Schubert et al. [37]. Roura et al. [64] investigated theoretically commonly used sintering models and came to the conclusion that a decreasing shrinkage rate during isothermal holding is inherent in standard theory of sintering. In Paper VI it was shown with DEM simulations in Paper VI that the establishment of a constant viscosity results in a retarding rate of rearrangment. These are certainly parts of the explanation to slow densification during isothermal holding. However, in experiments on Co spreading over WC surfaces de Macedo et al. [23] found very high spreading rates and that spreading rate is related to heating rate. During fast heating the metal spread on the carbide surface as much as in longer time [23]. Spreading was believed to be caused by viscous flow of a defect-rich contact zone of the metal and carbide. Defects are created in the surface zone as discussed by Schatt [33], but may also annihilate. With high heating rate the formation of defects is faster than the annihilation resulting in an increased defect concentration, lower viscosity of the contact zone and faster spreading of the binder [23]. The same argument may hold also for shrinkage. Mixing, milling and pressing of powders before sintering introduce a defect structure in the cobalt binder. Ac- cording to Gille et al. [35] this is important for the early stage of sintering, although Fischmeister and Exner [65] did not observe any drastic influence of milling, but related shrinkage behaviour solely to carbide grain size and surface activity. However, if a correlation between milling, defect concentration and shrinkage exist, a lower 26 Topics of present work

4 x 10 3 ) -2 2

1250oC, 0 min

60 min 1 180 min Number frequency (mm 480 min

0 -2 -1 0 1 logarithmic pore size

Figure 13. Pore size distributions measured after heating and isothermal holding. heating rate would allow for a longer annealing time during heating to the sintering time, which in turn would result in slower sintering kinetics. With a higher heating rate the material would reach the sintering temperature with higher defect concentration in the binder and higher sinterability. This would point to the existence of a thermal-history eﬀect. In Figure 14 the simultaneous densiﬁcation and microstructural evolution is shown for a WC-10 mass-% Co material with 3 µm carbide grain size. The micrographs were obtained after cooling from 1150◦C, 1200◦C, 1280◦C and 1380◦C, respectively. WC appears light grey, Co dark grey and pores black. Some of the features discussed above are demonstrated:

· Shrinkage starts around 800◦C,

· rearrangement initially leads to an aggregated WC-Co structure,

· solution-reprecipitation and facetting is limited during solid state sintering, i.e. below 1310◦C, but clearly visible after liquid phase sintering, and

· formation of the liquid phase leads to further rearrangement and a more homogeneous binder distribution.

Although high density is achieved in solid state for the more sinterable materials, liquid phase sintering still brings changes to the microstructure. Cobalt is present in the compacts as quite large entities all to the melting temperature. Accordingly, it is only with liquid phase sintering one can expect a fairly homogenous cobalt distribution. Microstructural evolution 27

1500 1

1400 0.9 1300 C ° 0.8 1200

1100 0.7 relative density temperature, 1000 0.6 900

800 0.5 0 2000 4000 6000 8000 10000 12000 time, s

Figure 14. Densiﬁcation and microstructural evolution.

M6C formation and dissolution: Paper VII

The work reported in Paper VII started after ﬁnding M6C in a WC-(Ti,W)C-Co material with carbon content above the limit of M6C stability after interrupting the heating cycle and cooling. Heating to a higher temperature and then cooling did not show any M6C. The explanation was found in the (Ti,W)C used for the material. (Ti,W)C has, unlike WC, a compositional range for carbon. As manufactured the carbide has a carbon content above the equilibrium content at the sintering temperature, but this excess carbon and the free carbon seems kinetically trapped to the carbide in the early stage of sintering. Accordingly, the material behaves as if it is below the M6C limit until carbon mobility is higher in the later stages of heating and the M6C can dissolve. ◦ Since the M6C carbide was dissolved only at 1350 C the sintering conditions con- stitute one major diﬀerence compared to WC-Co materials. Haglund and Ågren [27] found that conditions close to thermodynamic equilibrium are established at temperatures above 1000◦C, but this is evidently not true for WC-(Ti,W)C-Co materials. 28 Concluding remarks

This work has explored compositional effects of cemented carbide sintering. The main processes of densification have been confirmed rearrangement under capillary pressure provided by the binder and solution-reprecipitation. For more sinterable materials, such as WC-Co with fine grain size, rearrangement in solid state provides the most densification. Less sinterable materials, e.g. WC-Co with coarse carbides or containing cubic carbides, show also a considerable densification during liquid phase sintering. A continuum-type of model has been devised to describe shape changes during sintering based on the constitutive parameters uniaxial viscosity, viscous Poisson’s ratio and sintering stress. Explicit relations to cobalt content and carbide grain size have been derived. The driving force for shrinkage is provided by the interfacial energy relations. It is therefore troublesome that quantitative relations for these as function of sintering conditions are unavailable. It was shown that high carbon content has a positive effect on shrinkage rate. Solid state spreading and liquid wetting with their de- pendencies on carbon activity should therefore preferably be addressed in future research. The results obtained in this work supports the common view, see for instance Reference [29], of rearrangement of carbide grains and grain shape accomodation as the two sources of densification in cemented carbide sintering. Solid state sintering at finite sintering times (up to a few hours) leads in practice only to rearrangement. This is evident from microstructures obtained after sintering just below the solidus temperature. Such micrographs display only limited carbide facetting, which is an indication of little solution-reprecipitation. On the other hand, heating above the liquidus temperature leads rather quickly to an extensive facetting. It seems to exist a limit of solid state rearrangement at a relative density slightly above 0.8. Finer grained materials show a local densification rate max- imium around this value, which is typically reached some ten degrees below the binder melting temperature. A decreasing densification rate is then measured for a monotonously increasing temperature. It is reasonable to assume that this is caused by the two partly overlapping densification mechanisms, rearrangement and solution-reprecipitation, with different kinetic characteristics. The possibility that the event is related to liquid formation can be ruled out since the maximum occurs

29 30 Concluding remarks at different temperatures for different applied loads. Since solution-reprecipitation is rather slow in solid state the densification rate decreases until the liquid forms and solution-reprecipitation can occur significantly faster. In addition, the rearrangement process is facilitated by a more fluid binder and carbide contacts are dissolved due to the higher solubility. Coarse-grained materials show no minimum in densification rate during continuous heating. These materials reach the solidus-liquidus interval with a relative density below 0.8 and rearrangement occurs simultaneously as shape accomodation. It is known that powder metallurgy processing, cemented carbide sintering in- cluded, carries a certain degree of sensitivity towards changed conditions. One example is how a barely detectable amount of impurity can bring the outcome of sintering far from the expected [31]. Thus, truly predictive models and simulations seem difficult to realise. Nevertheless, it is the author’s hope that present work, or at least parts therefrom, will prove useful for the powder metallurgy community. One suggestion for future research has already been mentioned, i.e. the effect of sintering conditions on the wetting and spreading of Co on WC. Regarding densification mechanisms there are still open questions regarding the more complex sintering behaviour of materials based on ultrafine and nanocrystalline WC. An- other interesting field to explore further would be the carbide grain growth that perhaps does not have a decisive role in densification, but is increasingly important for finer-grained materials and not the least important for the mechanical properties. Acknowledgements

This work has been performed within the Brinell Centre Inorganic Interfacial Engi- neering (BRIIE), supported by the Swedish Agency for Innovation Systems (VIN- NOVA) and the following industrial partners: Erasteel Kloster AB, Höganäs AB, AB Sandvik Coromant, Seco Tools AB and Atlas Copco Secoroc AB. The author is grateful for their support enabling this work. The author also wishes to extend his sincere appreciation to a number of people who contributed in one way or another. This applies ﬁrst and foremost to Professor John Ågren for his professional guideance, co-operation, encouragement and trust in the direction of the work. Within the BRIIE sintering group the author is particularly grateful to Bo Jansson and Jan Qvick at Secotools AB for invaluable feedback, powder prepara- tion and co-authorship. Anders Olsson at Atlas Copco Secoroc AB ably provided powders. Björn Uhrenius formerly at AB Sandvik Hard Materials contributed at various stages of the work. Jenni Zackrisson at Secotools AB is thanked for help in powder manufacturing and co-authoring one paper. Sven Haglund at the Swedish Institute for Metals Research is acknowledged for support during the initial stage. Contributions by Per Lindskog and Lars Tholander at AB Sandvik Coromant made the experiments with sintering under load possible. Stefan Wikman at the Swedish Institute for Metals Research is thanked for help with the dilatometer and Kjell Jansson at Stockholm University for performing the last series of dilatometry. Some present and former colleges deserve to be mentioned. Joakim Odqvist is acknowledged in particular for coping with numerous interruptions for questions and discussions. The author is thankful to Lu Xiaogang for advices regarding thermodynamic calculations. Martin Schwind contributed in numerous ways, for instance with valuable pointers in scanning electron microscopy. Henrik Strand- lund is acknowledged particularly for discussions relating to numerical methods and Henrik Larsson for general discussions. Lars Höglund merits a special thanks for help with computers and software. The technical and administrative staﬀ at the Department of Materials Science and Engineering deserve recognition for their attempts to make practical matters run smoothely. My parents, brother, sisters and friends, thank you for being there. Finally, my wife and daughters are true sources of inspiration. Gali, Eden and Yuval—this is to you with love.

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