Cemented Carbide Sintering

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Cemented Carbide Sintering:

Constitutive Relations and

Microstructural Evolution

Anders Petersson

Stockholm 2004

Doctoral Dissertation

Royal Institute of Technology

Department of Materials Science and Engineering


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Akademisk avhandling som med tillstånd av Kungliga Tekniska Högskolan framläg-ges till offentlig granskning för avläggande av teknologie doktorsexamen onsdagenden 28 april 2004 kl 10.00 i Kollegiesalen, Administrationsbyggnaden, KungligaTekniska Högskolan, Valhallavägen 79, Stockholm.

Fakultetsopponent är Professor Colette Allibert, Institut National Polytechnique

de Grenoble, Frankrike.

Typsatt i LATEX.

ISBN 91-7283-739-x

ISRN KTH/MSE/--04/21--SE+METO/AVH

c Anders Petersson, March 2004

Högskoletryckeriet, Stockholm 2004


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Abstract

Cemented carbides based on tungsten carbide and cobalt are commonly produced

by a powder metallurgy route including liquid phase sintering. The pressed compactdensifies to almost half its volume during sintering due to pore elimination. Thesintering 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 den-sification, and the microstructural evolution during cemented carbide sintering.Dimensional changes are monitored using dilatometry with and without appliedexternal load. The microstructural evolution is investigated with light optical mi-croscopy and scanning electron microscopy. Thermodynamic calculations are usedas reference.

Constitutive relations are derived for uniaxial viscosity, viscous equivalent of Poisson’s ratio and sintering stress based on relative density and temperature. Therelations are extended to a model describing sintering shrinkage with explicit de-pendencies on carbide grain size and binder content. The model is divided in threestages 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 thestage 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 beenstudied. High carbon content enhances shrinkage rate, but the effect of grain size

distribution is rather small. The mean carbide grain size is insufficient to describedensification 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 simu-lated 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

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Preface

This thesis is based on the following appended papers that will be referred to inthe text by their Roman numerals.

I. A. Petersson and J. Ågren, “Constitutive behaviour of WC-Co materials withdifferent grain size sintered under load”, Acta Mater 52(2004):7, 1847–1858

II. A. Petersson and J. Ågren, “Numerical simulation of shape changes duringcemented 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—Effectsof carbide grain size and cobalt content”, submitted manuscript

IV. A. Petersson, “Sintering shrinkage of WC-CO and WC-(Ti,W)C-Co materialswith different carbon contents”, in manuscript

V. A. Petersson and J. Ågren, “Sintering shrinkage of WC-Co materials withbimodal grain size distributions”, submitted manuscript

VI. A. Petersson and J. Ågren, “Rearrangement and pore size evolution duringWC-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 princi-pal writer, planning most of the experiments and performing all calculations, areparticipation 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 performedmicroscopy for Papers II, V and VI and the main part of computer coding and sim-ulations for Paper VI. The author outlined the main part of analysis and evaluationfor Papers I–III and contributed to the evaluation for Papers V–VII.

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Contents

Introduction 1

Objectives 3

Liquid phase sintering 5

Surface energy relations . . . . . . . . . . . . . . . . . . . . . . . . . . . 6Sintering under applied stress . . . . . . . . . . . . . . . . . . . . . . . . 8

Cemented carbide sintering 9

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

Topics of present work 17

Constitutive modelling . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17Outlining the parameters: Paper I . . . . . . . . . . . . . . . . . . 17Composition and sintering shrinkage: Paper II . . . . . . . . . . . 19A shrinkage model for binder content and grain size: Paper III . . 20

Densification . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 22Influence of chemical composition: Paper IV . . . . . . . . . . . . . 22Effect of particle size distribution: Paper V . . . . . . . . . . . . . 23

Microstructural evolution . . . . . . . . . . . . . . . . . . . . . . . . . . 24Rearrangement: Paper VI . . . . . . . . . . . . . . . . . . . . . . . 24M6C formation and dissolution: Paper VII . . . . . . . . . . . . . . 27

Concluding remarks 29

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Introduction

Cemented carbides is a class of composite materials in which a carbide, often tung-sten carbide (WC), provides hardness and wear resistance and a metal binder,

almost exclusively cobalt (Co), provides toughness [1]. These materials have longbeen used in applications such as cutting, grinding, drilling and wear-resistant me-chanical components. A wide range of material properties can be obtained bychanging the material composition. Although the most often used carbide is thehexagonal WC, atlternatives are available. With cubic carbides such as TiC, NbCand TaC the high temperature properties are improved [2]. Minor additions, typi-cally less than a few percent by mass, of for instance Cr2C3 or VC, are often usedas 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 isremoved in a de-waxing step and the compacts sintered to eliminate porosity andincrease strength.

The microstructural events occuring during sintering are driven by the loweringof excess energy attributed to surfaces and the pressure that develops under acurved interface. The surface area and energy reduction may be accomplishedin different ways. One major division in sintering theory is between solid statesintering and liquid phase sintering, respectively, depending on whether or not anyliquid phase is present. Densification during solid state sintering is assumed tobe due to vacancy flux from pores to grain boundary sinks driven by the vacancy

concentration difference [3, 4]. Some possible mechanisms of matter transport areviscous or plastic flow, evaporation and condensation, volume diffusion and surfacediffusion [3].

The presence of a small amount of liquid phase is often beneficial for sinteringand three stages are indentified in classical liquid phase sintering theory [5], distin-guished by the microstructure evolution. Firstly, as the liquid forms, it flows out inthe microstructure and causes the solid particles to rearrange into denser packing.Secondly, the particle centres approach and density increases by means of solutionand 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

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2 Introduction

a liquid phase sintering process. Sintering is in practice almost always performedat temperatures where a quite significant volume fraction liquid exists. However,for finer grained cemented carbides much of the shrinkage, if not the most, occursduring heating to the sintering temperature and prior to liquid formation. Thisshrinkage shares however most typical features with the first stage of classical liquidphase sintering rather than solid state sintering processes.

The particle bonding and increased strength during sintering is often, but notalways, accompanied by pore elimination and densification. The term sintering istherefore frequently used synonymously with densification. Since sintering can beperformed without shrinkage, or even with swelling, the rate of increasing densitywill in the following be referred to as densification rate, shrinkage rate or strainrate 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 emphasison fundamental physical and chemical properties in Reference [1], industrial aspectsin Reference [2], mechanical properties in Reference [6], microstructural aspects inReference [7] and sintering in Reference [8].


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Objectives

Present work is intended to improve knowledge of sintering in cemented carbide sys-tems and propose models that can describe shape changes during cemented carbide

sintering. The work is founded on two pillars, namely the constitutive behaviourduring sintering and the concurrent microstructural evolution. These cannot beseparated 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 cementedcarbide sintering, but limited to medium–coarse carbide grain sizes. The main toolfor studying densification is dilatometry, which is a thermo-analytical (TA) tool,in which temperature and length of the sample are simultaneously recorded. Bytaking the time derivative of the length measurement it is possible to obtain theinstantaneous shrinkage rate that is modelled on a physically sound basis. The

microstructural evolution accompanying the shape changes is investigated mainlyusing scanning electron microscopy (SEM). Thermodynamic constraints are crucialfor sintering and thermodynamic calculations are therefore an integral part of thesintering 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 constitutivebehaviour is therefore on densification. However, friction between powder and diewall, as well as within the powder, may lead to a heterogenous packing in thegreen compact. This will in turn lead to heterogenous shrinkage during sinter-ing. Shrinkage calculations from given green densities should therefore be useful

in compaction-tool design aiming for near net shaping. Compaction and sinter-ing densification of a WC-Co material were successfully simulated in the worksof Brandt [9] and Haglund [10]. The composition of a cemented carbide dictatesextensively its shrinkage behaviour and the aim of present work is to extend theconstitutive 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 reasonablyshort-term research cannot be to fully describe the parameters with bearing ondensification. That would be a tremendous task. This work rather strives to isolatethe most crucial aspects of sintering and characterise their variaton with main

compositional parameters. Furthermore, experimental information on some of the

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4 Objectives

key parameters such as interfacial energies and contact angles with their variationswith sintering conditions are still largely lacking. This necessitates a fair amountof parameter fitting to study sintering of a given alloy once constitutive models areobtained. To predict, within narrow tolerances, the outcome of sintering one willprobably 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 tovirtually design materials, give promises for a faster and more cost efficient materialsdevelopment. The desire for sintering models that can be used to simulate sinteringvia computer calculations is founded in this context. With such simulations itshould be possible to explore effects of various compositional aspects and processparameters under different sintering conditions.


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Liquid phase sintering

During sintering the excess surface energy attributed to porosity is decreased bymeans of matter transport that becomes kinetically possible at high temperature.

Thus, sintering deals with a central issue of materials science, i.e. how much isthe total energy of the system lowered and in what way. When penetrating theprocess in order to understand or predict its course, one inevitably finds severalpossible reaction steps, making sintering a quite complicated process. Not only isthe total energy reduced by pore elimination, but also by microstructural coarseningand in many cases through phase transformations. Under such circumstances thedescription of sintering is further complicated.

For complete densification by liquid phase sintering three things are usuallyrequired, namely (i) an appreciable amount of liquid, (ii) solubility of the solid inthe 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 respectivedominant process.

When the liquid phase forms it wets the solid particles resulting in capillaryforces rearranging the particles into a denser packing. The powder compact re-sponds to the capillary pressure as a viscous solid and viscosity increases withdecreasing porosity [5]. Rearrangement is retarded by friction between solid par-ticles [16], but is usually so fast that all other events are overshadowed and fulldensity can be reached by rearrangement if around 35 volume percent liquid ispresent [5].

In the second stage densification occurs with shape accomodation of particlesseparated by liquid films through solution and reprecipitation of solids leading tocloser packing [16]. Material is dissolved at solid-liquid interfaces with high chem-ical potential and reprecipitated at other sites with lower chemical potential. Anumber of solution-reprecipitation mechanisms have been identified, e.g. Ostwaldripening, Ostwald ripening with shape accomodation, contact flattening, particledisintegration, coalescence and pore elimination by liquid flow [17]. The solubilityof the solid in the liquid is inversely related to the particle size. The resultingconcentration gradient causes material flow from finer particles to coarser parti-cles. Thus, the second stage is usually marked by microstructural coarsening [5].

One possible process is illustrated with the initial configuration of two particles

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6 Liquid phase sintering

separated by a liquid film in Figure 1 and after some material from the particlecontact has dissolved in the liquid and precipitated elsewhere leading to a denserconfiguration in Figure 2. In the case of equally sized spherical particles the stressgradient driving dissolution in the contact zone stems from the capillary pressureprovided by the binder. Both figures are redrawn from Reference [11].

r

Liquid

δR

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

h x

r

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-energyconfiguration of the microstructure during liquid phase sintering [5].

Wetting is important for the sinterability of a powder system and characterisedby the contact angle, given by the solid-liquid-vapour equilibrium. This is demon-

strated in Figure 3 drawn after Reference [5].


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Surface energy relations 7

solid

liquid

vapour

θ θγ LV

γ SVγ SL

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 equa-tion [5]

γ SV = γ SL + γ LV cos θ, (1)

where the respective interfacial energies are solid-vapour, γ SV , solid-liquid, γ SL,and liquid-vapour, γ LV . The contact angle is denoted θ. Good wetting is char-acterised by a small contact angle and conversely poor wetting by a large contactangle. 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 energythe net interfacial energy is reduced by spreading the liquid over the solid surfaceas demonstrated in Figure 4 after Reference [5]. Chemical affinity and solubilitybetween the phases generally contribute to efficient spreading [5].

liquid

solid

vapour

γ LV

-γ SV

+γ SL

< 0

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 r1and r2 the energy change becomes [16]

dE

dV = γ

dA

dV = γ 1r1 +

1

r2, (2)


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8 Liquid phase sintering

where γ is the surface energy and dA the change in surface area. Identifying theterm dE/dV as a stress σ, the Laplace equation

σ = γ

1r1

+ 1r2

(3)

is obtained. If the surface is convex the stress is tensile and if the surface is concavethe stress is compressive. It is this stress gradient that drives material transportduring sintering [16].

Sintering under applied stress

It is self evident that densification during sintering is enhanced by an applied pres-sure if it is sufficiently large. On a micro-level such a pressure enhances not onlyrearrangement, but also solution at particle contacts and plastic flow in the solidparticles [18]. By applying an external pressure less sinterable powders can be sin-tered to high density and strength [16]. The applied pressure is often modelledas additive to the capillary pressure [5, 18]. For an extensive review on sinteringmodelling, including application of load, see Reference [19].


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Cemented carbide sintering

In industrial practice cemented carbides are sintered by liquid phase sintering attemperatures in the range 1350–1650◦C leading to virtually pore-free compacts [2].

The foundation of the process in terms of thermodynamic and kinetic aspectsof the alloy systems has been studied extensively, as well as their interaction in thedeveloping microstructures, see for instance References [1, 2, 6–8] mentioned in theintroduction.

Freytag and Exner [20] studied sintering through changes in the magnetic prop-erties of a WC-12 mass-% Co material. It was found that below 800◦C practicallyno shrinkage is observed but stresses introduced by milling are relieved and cobaltoxide on surfaces are reduced. In the range 800–1100◦C shrinkage is limited, tung-sten 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 inthe cobalt-saturated carbide grain boundaries should have a decisive role in solidstate sintering. When the binder melts a substantial amount of carbide is dissolvedin the liquid. Rearrangement and pore filling eventually lead to full density.

Surface oxides, practically always present in a powder, are reduced in a widetemperature range 650–850◦C with a maximum around 750◦C with an additionalpeak in CO/CO2 emission at 950◦C for doped WC materials [21]. Freytag andExner [20] relate the main carbon loss during sintering to reaction of oxygen at thecarbide 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 andCO2 from adhesive oxygen and oxides with carbon from WC, which leads to theconclusion that shrinkage starts when all surface oxides are reduced [22]. This washowever recently questioned by de Macedo et al. [23] who studied wetting of Co, Niand 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 reductionalthough they appear in sequence. Göthelid et al. [24] studied Co spreading on WCsubstrates with scanning tunneling microscopy (STM). It was found that at firstthe substrate is partially covered by a so called 2×2 structure on top of which Cothen 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

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10 Cemented carbide sintering

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

Thermodynamic aspects

The simplest cemented carbide and the base for most alloys is the mixture of WCand 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 car-bides [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 foronly small deviations additional phases will form. For low carbon contents the M6Cmay become stable and for high carbon contents graphite. Neither one is usuallydesired 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 thenarrow compositional range for the two-phase field the carbon activity increasesfrom 0.3 to 1 with respect to graphite.

0

0.2

0.4

0.6

0.8

1.0

0 0.2 0.4 0.6 0.8 1.0

x C

xW

M6C

fcc-Co

W2C

Co7W6

WC

liq-Co M12C

graphite

bcc-W

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 desiredin 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


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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-%. Sincesome carbon inevitably reacts with oxygen during sintering it is in practice oftennecessary to add carbon to the WC-Co mixture to obtain the desired compositionof the finished product [2].

1100

1200

1300

1400

1500

5.2 5.4 5.6 5.8

t e m p e r a t u r e ,

°

C

carbon content, mass-%

M6C

WC

fcc-Co

WC

liquid Co

WC

fcc-Co

graphite

WC

fcc-Co

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 thecobalt-rich liquid becomes stable at lower temperature [26].

It is believed that conditions close to equilibrium prevail at temperatures typ-ically above 1000◦C owing to the short diffusion distances involved in cementedcarbide 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 forcefor 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 rearrange-ment 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


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density obtained in the very first stage of sintering even before the liquid is formeddistinguishes the WC-Co system from other liquid phase sintering systems [29].

Froschauer and Fulrath [30] studied the microstructural evolution in situ andconfirmed that rearrangement is the main densification process. Haglund and co-workers [31] showed for a WC-10 mass-% Co material that densification in solidstate can be treated as a rearrangement process like in the first stage of liquidphase sintering. No drastic change in densification rate occurred as the bindermelted.

During ball milling of the powder the carbide particles become partially coveredby cobalt and small carbide fragments are embedded in Co, which is important forthe initial sintering [32]. Milling and cold pressing also create highly deformed con-tact areas with high internal stresses that may deform by creep or other transportmechanisms 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 thanWC-Co milled separately and mixed. Densification started locally in cobalt-richareas packed with carbides from milling.

At temperatures above one half of the melting temperature, 650–710◦C for finegrained cemented carbides, the dispersed binder particles behave like a viscous massspreading by wetting onto the surrounding carbide surfaces [33,35–37]. The spread-ing of the interfacial zones immediately after surface oxide reduction constitutes thefirst stage of densification [21].

The surface free energy is lowered by the cobalt spreading [38] and exchangingWC-gas and Co-gas interfaces for WC-Co lowers the interfacial energy to approxi-

mately one fifth [35]. Solid state spreading of cobalt is driven by the surface energyrelations in a similar way as desribed by the Young’s equation, Equation 1, al-though 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 observedthat it is only a film that spreads over the surface and not the bulk metal. Themechanism behind binder spreading is believed to be viscous flow of a defect-richcontact zone [23]. Meredith and Milner [36] suggested that the viscosity in thecarbide-metal interface is further reduced by carbide dissolution owing to enhanceddiffusion.

The rate of binder spreading is limited by factors such as area, curvature and

chemistry of the carbide particles and content, dispersion, chemistry and physicalproperties of the binder. Binder spreading is connected to the Laplace forces actingalong the wetting front line between carbide and binder. These microscopic forcesrearrange the carbide particles and reduce the mean distance between neighbouredparticles, resulting in densification [35].

The smaller the grain size is the more interfacial area is contained within thecompact and the higher is the stored interfacial energy. As a result sintering startsearlier for fine-grained cemented carbides and the share of solid state shrinkageincreases [35]. Moreover, milling and mixing introduce more defects for finer WCmaterials, which results in higher internal stress, shorter distance between source

and sinks of diffusion and thus lower viscosity [39]. The maximum in shrinkage rate


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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 priorto sintering. For ultrafine cemented carbides the solid state sintering can contributewith 90 percent of the total densification [35]. Earlier shrinkage is observed forattritor-milled powders compared to balled-milled powders [37].

Grain growth inhibitors such as VC and Cr2C3 in small amounts shift shrinkageonset to higher temperature, decrease the solid state shrinkage rate and lower themelting point of the binder [21, 35]. Addition of either carbon or tungsten to afine-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 dividedin two stages. The first stage of densification, typically below 1100–1200◦C, isbelieved to be homogeneous shrinkage due to binder spreading and creep, whilein 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 forcein driving densification, while in the second stage the binder is annealed and itsintrinsic stresses neglible compared to the capillary stress. Moreover, in the firststage the binder is nearly free to creep, while in the second stage creep is restrictedby the large number of carbide-binder interfaces [35]. Densification in the first stagedepends strongly on processing prior to sintering and only weakly on grain size,while in the second stage a weak dependence on processing but strong dependenceon grain size are observed. The solid state shrinkage appears more complex forultrafine 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 accomplishedthrough grain shape accomodation and for cemented carbides with contributionsfrom contact flattening where carbide particles are in contact, dissolution of smallgrains 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 pris-matic shape is not necessarily compatible with the aggregated structure formed inthe initial stage. With a sufficient amount of binder this may lead to additionalrearrangement [36].

Practically full density is normally reached for most WC-Co materials by re-arrangement and shape accomodation. Sintering of the solid phase skeleton, i.e.the last stage of classical liquid phase sintering is usually not observed in cementedcarbide sintering. Other microstructural changes may however occur, for instancegrain size and grain size distribution, grain shape and binder phase distribution

may all change [29].


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Since full density is reached through continuous repacking of clusters leavingcobalt free to flow further into the structure Snowball and Milner [34] argue that itis not necessary to postulate a third stage of shrinkage, but with low cobalt contenta slow final densification pertaining to solid state sintering of the carbides may beobserved [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 duringand after sintering, while Nelson and Milner [38] claimed that there is a cobalt filmbetween carbide particles for short sintering times. This would allow for particlesliding during densification, but a rigid carbide network will eventually form duringprolonged sintering. Gurland and Norton [14] ruled out the possibilty of a carbidenetwork based on infiltration experiments. Andrén [41] refers to measurementswith atom probe field ion microscopy (APFIM) and analytical electron microscopy

(AEM) showing approximately half a monolayer Co is present between carbidegrains. Accordingly, there is a continous carbide skeleton, but half the contactsurface 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 ad-dition to retained porosity are a few of the more important microstructure char-

acteristics that influence the properties such as transverse rupture strength andhardness [6]. Primary interest here is, however, the microstructural evolution con-current with densification.

Along with dimensional change, i.e. shrinkage and creep deformation, the mi-crostructure changes too. If it was not for the microstructural change, compactshape would not actually change since it is through the rearrangement of carbideparticles and grain shape accomodation that densification occurs.

The carbides are rearranged from a rather homogenous spatial distribution inthe green body to an aggregated structure during the initial stage of shrinkage, andMeredith and Milner [36] attributed the rapid initial densification to the formation

of close-packed boundaries between carbides when cobalt diffuses into the carbideinterfaces and bringing the particles closer. In a second stage of densification thebinder bridges can rearrange the carbide particles [21]. During the intense shrinkagein solid state WC-Co clusters that have formed upon cobalt spreading behave likea viscous mass, and the clusters creeps into the surrounding void area. SmallerCo particle size and better dispersion leads to more WC-Co clusters formed duringbinder spreading, shorter diffusion distances and more uniform capillary forces. Theresult 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


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Microstructural evolution 15

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

Figure 7. Micrograph obtained after heating a WC-Co material to 1280◦C andcooling. 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 evidencedby the prismatic grain shape typical for WC-Co materials after prolonged sinter-ing [36]. The preference for prismatic growth stems from the WC crystal structureand interfacial energy anisotropy [1, 42]. Prismatic carbide shapes develop faster

in fine-grained materials than coarse-grained materials [38] and the shape of thefacetted WC grain changes with the carbon content [42–44]. The carbide facettingis evident in Figure 8 which was obtained in SEM with secondary electron modeafter sintering the same material as in Figure 7 at 1430◦C for 60 min. This sinteringcycle is similar to the one used in industrial practice.

The grain size of the sintered structure basically depends on the initial sizedistribution and thus the milling conditions [45]. Large grains may act as startingpoint for rapid discontinuous grain growth widening the grain size distribution.

One trend in more recent materials development is towards finer carbide grainsizes, since this improves hardness and wear resistance [46]. This has naturally

led to studies devoted to sintering mechanisms in submicron and nanocrystalline


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Figure 8. Micrograph obtained after sintering a WC-Co material at 1430◦C for60 min. WC appears light grey, Co dark grey and pores black.

cemented carbide materials. Since this study is limited to medium–coarse grainsizes the reader is referred to for instance References [35,47,48].


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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 capillaryaction of the binder. Rearrangement is driven, in the absence of an applied stress,by the sintering stress inversely proportional to grain size and linearly propotionalto the pore fraction. Binder spreading was considered practically instantaneousand diffusional creep of the binder rate limiting. A linear combination of surfacediffusion 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 intoaccount 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. Aslightly different approach, used for instance by Gillia et al. [51] for WC-Co, waschosen in which the uniaxial viscosity and the viscous equivalent of Poisson’s ratioare derived from sintering under applied axial load. The sintering stress can becalculated from these two parameters and the shrinkage rate under free-sinteringconditions, i.e. without the application of load. The evaluation of the constitutiveparameters followed in much the approach by Bordia and Scherer [52,53].

The constituive parameters were modelled as functions of temperature and rela-tive density assuming that the material is Newtonian, i.e. the strain rate is linearlyrelated 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 duringliquid phase sintering and marked by an initially decreasing viscosity. Note thenthe drastic increase as the materials approaches full density during liquid phasesintering 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


26/45

18 Topics of present work

roughly twice as high in the second stage. For liquid phase sintering a little dif-ferent approach was chosen than usual. With the analogy of free volume in glasstransition and percolation theory [55,56] a description of rearrangement facilitatedby the present porosity was suggested. Combining the deformation pertaining tothe dense material and the rearrangement an effective viscosity was obtained withthe desired properties of a fast increase with relative density, but still allowing afinite deformation rate of the dense material. It should be pointed out that thesuggested model perhaps does not give the easiest numerical treatment. If theviscosity of the dense material is neglected an exponential function describing theproper increase with relative density should be suitable.

0 500 1000 1500 2000 2500

800

900

1000

1100

1200

1300

1400


° C

time, s

fine

coarse

103

104

105

106

u n i a x i a l v i s c o s i t y ,

M P a

s

103

104

105

106

u n i a x i a l v i s c o s i t y ,

M P a

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 anexponential relation, as used by Gillia et al. [51], Scherer’s model [57]

ν p ≈ 1

2

ρ

3− 2ρ, (4)

was used. This was derived for a particular pore arrangement for a different ma-terial, but captured the essence of derived data in Paper 2. Moreover, it has thebenefits of a rather simple relation to relative density as sole parameter and thecorrect limits of ν p → 0 for ρ → 0 and ν p → 0.5 for ρ → 1. Although a satisfactoryrepresentation is found with this model it is naturally a source of uncertainty inthe analysis.

The sintering stress is the thermodynamic force associated with sintering. Itcan 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


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Constitutive modelling 19

view the sintering stress is a pressure of equal magnitude as the hydrostatic tensilestress that would stop shrinkage. The calculation of sintering stress in Paper Ishowed a value slightly increasing with decreasing porosity, which contrasts theassumptions in References [49, 50]. The sintering stress was described by a powerfunction of relative density.

In Paper I the influence of an applied stress on densification was studied aswell. It is a commonly held opinion that the hydrostatic component of an appliedstress enhances densification [5, 16]. Though the applied uniaxial load increasesaxial shrinkage and decreases radial shrinkage the total densification is expected toincrease. The first series of experiments with axial stresses up to 0.7 MPa influenceddensification rate only marginally, and a second series was performed under muchhigher 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 arelative 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 stressof 0.7 MPa and 150 N in 6 MPa at most. The carbide grain size was 1.7 µm. Similardiagrams obtained under free sintering for a variety of compositions together withcalculated solidus and liquidus temperatures indicate that the first local maximumis related to solid state shrinkage and the second to liquid phase shrinkage.

0.5 0.6 0.7 0.8 0.9 10

0.2

0.4

0.6

0.8

1

1.2x 10

-3

relative density

d e n s i f i c a t i o n r

a t e

1 N

15 N

150 N

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

Composition and sintering shrinkage: Paper II

The general constitutive behaviour was outlined in Paper I, although some pa-rameters scattered. Moreover, a model of more practical interest should accountexplicitly for parameteters of the material composition. The aim of Paper II was

therefore twofold. Firstly, the suggested models were used to describe shrinkage


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data from dilatometry where experimental certainty normally is rather high. Sec-ondly, results from sintering a number of different materials should point to thegeneral influence of respective parameter on shrinkage. Microstructures were exam-ined at various stages of sintering before and after liquid formation using scanningelectron microscopy.

The microstructural evolution was found generally coherent with findings in forinstance References [31,36,59], with the main contribution to densification from re-arrangement, although leading to clustering, and subsequent solution-reprecipitation.

Different carbide grain sizes were shown to still result in similar temperaturesof shrinkage onset. However, shrinkage rate during solid state sintering decreaseswith increasing carbide grain size, and for a material with approximately 10 µmWC the main solid state densification was suppressed. The latter can be linked tothe microstructural evolution showing that although the binder spreading onto the

neighbouring carbide particles is fast the continuous linking of more carbides into anagglomerate is impeded. Following the decrease in shrinkage rate during solid statesintering the relative amount of shrinkage during liquid phase sintering increaseswith carbide grain size. With a medium carbide grain size, typically 2 µm, only afraction of the densification is due to liquid phase sintering. For coarse materials,with carbide grain sizes around 5 µm, densification before and after liquid formationare 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 carbidegrain 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 outlinedin Paper I can be combined to successfully describe free sintering shrinkage ratein three stages as shown in Figure 11 for a WC-Co material with 2 µm WC and10 mass-% Co. At the transition from the first stage to the second stage themodelled curves overlap and cannot be distinguished. It was also shown how thesemodels can be used to calculate sintering shrinkage by numerical integration for agiven sintering cycle.

A shrinkage model for binder content and grain size: Paper III

The constitutive relations outlined in Paper I were shown suitable for describingshrinkage in Paper II. With phenomenological relations for the effects of cobaltcontent and carbide grain size on the respective parameters uniaxial viscosity, vis-cous equivalent of Poisson’s ration and sintering stress the three stages of shrinkagecan be described with no less than 13 adjustable parameters. In Paper III these13 parameters were reduced to only four by fitting model parameters to dilatometrydata for three cobalt contents and three carbide grain sizes. The remaining fourparameters are enough to reasonably well represent shrinkage for other materials.


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Constitutive modelling 21

0 1000 2000 3000 4000 5000

-2.5

-2

-1.5

-1

-0.5

0

x 10 -4

time, s

f r e e s

i n t e r i n g s

t r a i n

r a t e ,

1 / s

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

The obtained model of linear shrinkage with adjusted parameters reads

ε̇1 = −Aρ2.1

d0.063

exp(−Q1/RT ) (1− 2ν p)RT (5)

ε̇2 = −Bρ2.1d−0.035 exp(−Q2/RT ) (1− 2ν p)

RT

1− f WC1 + 3f WC

(6)

ε̇3 = −Cρ4.1f 0.69

Co (1− 2ν p)

1 +

0.96exp(−0.096/ (1− ρ))

d0.45f 0.18Co

(7)

for the respective stages. ρ is the relative density, R the gas constant, T absolutetemperature, ν p 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, f WC volume fraction carbide and f Co volume fraction cobaltin the dense material. The change from the first to the second stage occuredaround 1100◦C. In the solidus-liquidus interval the shrinkage rate is expected to bea combination of the solid fraction of the material shrinking according to Equation 6and the liquid fraction shrinking according to Equation 7.

Two adjustable parameters, A and Q1 refer to the first stage of shrinkage. Theapparent activation energy Q1 was found in the range 60–140 kJ/mol. In this stageof sintering Gille et al. [35] observed activation energies of densification rangingfrom 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


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diffusion of Co in the WC-Co interfaces discussed in the literature [20,36] may beresponsible for the very low activation energies observed. In this stage of sinteringdeformation is most likely governed by surface diffusion of Co over WC surfaces.

In the second stage all materials showed practically the same apparent activationenergy of 250 kJ/mol in Reference [35], which was also observed for most materialsin Papers I and only a little higher for uniaxial viscosity in Paper III. This valueis typical for volume diffusion of cobalt [61], and it is reasonable to assume thatdeformation is governed by diffusional creep pertaining to volume diffusion of thecobalt-rich binder. Missiaen and Roure [62] analysed WC-Co shrinkage with amorphological description of shrinkage. They conclude that densification and binderspreading are coupled and that volume diffusion in the cobalt phase is the governingmechanism.

As discussed in Paper III an accurate model of densification would preferablydescribe shrinkage for a material with only a few, if any, adjustable parameters. Itwas also pointed out, as well as in Reference [35], that the early shrinkage is sensitiveto processing prior to sintering, and accordingly the first stage of shrinkage differsbetween different materials, which then must be reflected by model parameters.Moreover, Gille et al. [22] investigated a new ultrafine WC material where thecarbide had a relatively small lattice distortion compared to other ultrafine grades.Given the lower solution pressure for the less distorted lattice the solution andreprecipitation was slower resulting in a more stable sintering, less coarsening anda more homogenous microstructure [22]. Accordingly, processing influences also the

later stages of densification and it seems difficult to realise one common descriptionof 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 shrink-age. Three WC-Co and and three WC-(Ti,W)C-Co materials with different carboncontents were investigated.

Additions of the cubic carbide (Ti,W)C delays shrinkage to higher temperatureand high carbon content above the limit of graphite formation enhances shrinkagefor both WC-Co and WC-(Ti,W)C-Co materials. A low carbon content towardsthe limit of M6C formation delays shrinkage for a material with (Ti,W)C but not

for WC-Co.


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Densification 23

The densification model of Paper III with four adjustable parameters describedshrinkage reasonably well for WC-Co materials with low and medium carbon con-tent 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 solidstate shrinkage of WC-(Ti,W)C-Co with low and medium carbon contents. ForWC-(Ti,W)C-Co the three stages of shrinkage is reduced to two stages where thefirst stage of lower activation energy in WC-Co is absent. The observed variationof shrinkage rate with temperature points to a much higher apparent activationenergy than the 250 kJ/mol observed in Papers I and III and Reference [35].

Freytag and Exner [20] related the cobalt spreading and thereby increasingspecific surface in the initial stage of WC-Co sintering to the change in coerciveforce occuring in the temperature range 800–1100◦C. Since continued heating doesnot change the magnetic properties much [20], it may be concluded that cobalt

spreading has been completed at this stage. At this temperature densification ratestarts to increase considerably and here was also the transition from the first stageof shrinkage to the second stage found in Paper III and Reference [35].

However, with a significant amount of WC exchanged for the mixed cubic car-bide (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 isin opposition to the prevailing view that wetting is significantly poorer on oxidesthan on carbides. It is plausible that relatively stable surface oxides present onthe carbide surfaces effectively prevent binder spreading. With an oxide stabilityincreasing 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 etal. [23] that Co spreads on a WC as a very thin layer, and by Göthelid et al. [24] thata pre-cursor layer of Co must form on top of the WC surface before spreading canoccur. It seems reasonable that surface diffusion in the first stage of shrinkage formsa Co film over the carbide with spreading completed around 1100◦C. This leads alsoto a little rearrangement and densification. When temperature is high enough forvolume 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 filmis lacking. Thus, shrinkage behaviour of the two less sinterable types of materialsstudied, WC-Co with coarse carbide and WC-(Ti,W)C-Co, are different. In theformer case the second stage is suppressed as discussed above and in the latter casethe first stage.

Effect of particle size distribution: Paper V

This work was initiated to investigate how the width of a particle size distributionaffects shrinkage and if the mean carbide grain size is sufficient in models such as

those devised in Papers I–III.


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Shrinkage of materials with different unimodal and bimodal particle size distri-butions were compared. It was found that the mean grain size, indeed, is sufficientto describe shrinkage for a bimodal distribution with initially 75 percent fine par-ticles and 25 percent coarse by mass. With 50 percent each of fine and coarseparticles, 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 clearlydemonstrated. With bimodal particle size distribution an extreme coarsening wasobserved during liquid phase sintering in addition to marked facetting. A unimodalparticle size distribution showed much less coarsening but still exclusively facettedcarbides.

Microstructural evolution

Rearrangement: Paper VI

This study was started in light of the observation by Haglund [10] that the averagepore size in WC-Co seems to shrink during heating but grow during isothermalholding although macroscopic shrinkage is observed.

The experimental basis for the work was measurements of the pore size distri-bution for various stages of a heating cycle with and without imposed isothermal

holding. Sintering was performed below the temperature of liquid formation. Thepore size distribution, if plotted with number frequencies on a logarithmic scalechanged in a nearly self-similar way as indicated by Figure 12 where Gaussian dis-tribution curves have been adjusted to measured data for the ease of comparison.With isothermal holding time bimodal pore size distributions developed as shown inFigure 13 for the holding times 60, 180 and 480 min. Isothermal holding after slowheating resulted in a pore size distribution with less pronounced bimodal characterthan after fast heating.

The formation of bimodal pore size distributions was viewed in light of particlerearrangement into clusters where pores inside clusters shrink and pores between

clusters grow. The direct cause of clustering in liquid phase sintering is that theinter-particle attraction force increases with decreasing particle distance, which en-hances irregularities in particle packing [63]. The formation of clusters was verifiedin simulations using a Discrete Element Method (DEM), where the capillary at-traction between particles connected by a liquid bridge was modelled. The particleattraction forces, contact forces and viscous drag forces were used to solve the equa-tion 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. Withincreasing temperature the solubility increases monotonously, while at isothermal

conditions the equilibrium solute content is rather soon established and a more rigid


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-2 -1 0 10

1

2

3

4

5x 10

4

logarithmic pore size

N u m b e r f r e q u e n c y ( m m - 2 )

1150oC

1200oC

1250oC

Figure 12. Pore size distributions measured after heating.

structure forms. Moreover, DEM simulations of isothermal holding following heat-ing showed that the microstructural evolution retards considerably simply by therealisation of constant viscosity. The DEM simulations failed however to accountfor solubility effects leading to significant densification.

There has been a question about a possible influence of heating rate on shrinkagerate. 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 etal. [37]. Roura et al. [64] investigated theoretically commonly used sintering modelsand came to the conclusion that a decreasing shrinkage rate during isothermalholding is inherent in standard theory of sintering. In Paper VI it was shown withDEM simulations in Paper VI that the establishment of a constant viscosity resultsin a retarding rate of rearrangment. These are certainly parts of the explanationto slow densification during isothermal holding. However, in experiments on Cospreading over WC surfaces de Macedo et al. [23] found very high spreading ratesand that spreading rate is related to heating rate. During fast heating the metalspread 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 alsoannihilate. With high heating rate the formation of defects is faster than theannihilation resulting in an increased defect concentration, lower viscosity of thecontact zone and faster spreading of the binder [23].

The same argument may hold also for shrinkage. Mixing, milling and pressingof 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, althoughFischmeister and Exner [65] did not observe any drastic influence of milling, but re-lated shrinkage behaviour solely to carbide grain size and surface activity. However,

if a correlation between milling, defect concentration and shrinkage exist, a lower


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-2 -1 0 10

1

2

3x 10

4

logarithmic pore size

N u m b e r f r e q u e n c y ( m m - 2 )

1250oC, 0 min

480 min

180 min

60 min

Figure 13. Pore size distributions measured after heating and isothermal holding.

heating rate would allow for a longer annealing time during heating to the sinteringtime, which in turn would result in slower sintering kinetics. With a higher heatingrate the material would reach the sintering temperature with higher defect concen-tration 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 mi-crographs 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 thefeatures 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 morehomogeneous 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 presentin 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.


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0 2000 4000 6000 8000 10000 12000800

900

1000

1100

1200

1300

1400

1500


° C

time, s

0.5

0.6

0.7

0.8

0.9

1

r e l a t i v e d e n s i t y

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-Comaterial with carbon content above the limit of M6C stability after interruptingthe heating cycle and cooling. Heating to a higher temperature and then coolingdid 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 hasa carbon content above the equilibrium content at the sintering temperature, butthis excess carbon and the free carbon seems kinetically trapped to the carbide inthe early stage of sintering. Accordingly, the material behaves as if it is below theM6C limit until carbon mobility is higher in the later stages of heating and theM6C 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 tem-peratures above 1000◦C, but this is evidently not true for WC-(Ti,W)C-Co mate-rials.


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28


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Concluding remarks

This work has explored compositional effects of cemented carbide sintering. Themain processes of densification have been confirmed rearrangement under capillary

pressure provided by the binder and solution-reprecipitation. For more sinterablematerials, such as WC-Co with fine grain size, rearrangement in solid state providesthe most densification. Less sinterable materials, e.g. WC-Co with coarse carbidesor containing cubic carbides, show also a considerable densification during liquidphase sintering.

A continuum-type of model has been devised to describe shape changes duringsintering based on the constitutive parameters uniaxial viscosity, viscous Poisson’sratio and sintering stress. Explicit relations to cobalt content and carbide grainsize 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 sinteringconditions are unavailable. It was shown that high carbon content has a positiveeffect on shrinkage rate. Solid state spreading and liquid wetting with their de-pendencies on carbon activity should therefore preferably be addressed in futureresearch.

The results obtained in this work supports the common view, see for instanceReference [29], of rearrangement of carbide grains and grain shape accomodation asthe two sources of densification in cemented carbide sintering. Solid state sinteringat 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 anindication of little solution-reprecipitation. On the other hand, heating above theliquidus temperature leads rather quickly to an extensive facetting.

It seems to exist a limit of solid state rearrangement at a relative densityslightly above 0.8. Finer grained materials show a local densification rate max-imium around this value, which is typically reached some ten degrees below thebinder melting temperature. A decreasing densification rate is then measured fora monotonously increasing temperature. It is reasonable to assume that this iscaused by the two partly overlapping densification mechanisms, rearrangement andsolution-reprecipitation, with different kinetic characteristics. The possibility that

the event is related to liquid formation can be ruled out since the maximum occurs

29


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30 Concluding remarks

at different temperatures for different applied loads. Since solution-reprecipitationis rather slow in solid state the densification rate decreases until the liquid forms andsolution-reprecipitation can occur significantly faster. In addition, the rearrange-ment process is facilitated by a more fluid binder and carbide contacts are dissolveddue to the higher solubility. Coarse-grained materials show no minimum in densi-fication rate during continuous heating. These materials reach the solidus-liquidusinterval with a relative density below 0.8 and rearrangement occurs simultaneouslyas shape accomodation.

It is known that powder metallurgy processing, cemented carbide sintering in-cluded, carries a certain degree of sensitivity towards changed conditions. Oneexample is how a barely detectable amount of impurity can bring the outcome of sintering far from the expected [31]. Thus, truly predictive models and simulationsseem 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 den-sification mechanisms there are still open questions regarding the more complexsintering behaviour of materials based on ultrafine and nanocrystalline WC. An-other interesting field to explore further would be the carbide grain growth thatperhaps does not have a decisive role in densification, but is increasingly importantfor finer-grained materials and not the least important for the mechanical proper-ties.


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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 authoris grateful for their support enabling this work.

The author also wishes to extend his sincere appreciation to a number of peoplewho contributed in one way or another. This applies first and foremost to ProfessorJohn Ågren for his professional guideance, co-operation, encouragement and trustin the direction of the work.

Within the BRIIE sintering group the author is particularly grateful to BoJansson 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 atvarious stages of the work. Jenni Zackrisson at Secotools AB is thanked for help inpowder manufacturing and co-authoring one paper. Sven Haglund at the SwedishInstitute for Metals Research is acknowledged for support during the initial stage.Contributions by Per Lindskog and Lars Tholander at AB Sandvik Coromant madethe experiments with sintering under load possible. Stefan Wikman at the SwedishInstitute for Metals Research is thanked for help with the dilatometer and KjellJansson at Stockholm University for performing the last series of dilatometry.

Some present and former colleges deserve to be mentioned. Joakim Odqvist isacknowledged in particular for coping with numerous interruptions for questions

and discussions. The author is thankful to Lu Xiaogang for advices regardingthermodynamic calculations. Martin Schwind contributed in numerous ways, forinstance with valuable pointers in scanning electron microscopy. Henrik Strand-lund is acknowledged particularly for discussions relating to numerical methodsand Henrik Larsson for general discussions. Lars Höglund merits a special thanksfor help with computers and software. The technical and administrative staff atthe Department of Materials Science and Engineering deserve recognition for theirattempts to make practical matters run smoothely.

My parents, brother, sisters and friends, thank you for being there. Finally, mywife and daughters are true sources of inspiration. Gali, Eden and Yuval—this is

to you with love.

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Cemented Carbide Sintering

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