, Volume 1, Issue 3, pp 198–212 | Cite as

Development of manufacturing technology on WC–Co hardmetals

  • Hongbo NieEmail author
  • Taiquan Zhang
Review Paper


Hardmetals are tungsten carbide (WC)-based composites, which are made of WC as a hard phase and transition metals such as Co, Fe, or/and Ni as ductile binder matrices. Their properties can be mainly tailored through the grain sizes of the sintered carbides and the amount of metallic binder. As successful tool materials, hardmetals are widely applied in metal cutting, wear applications, chipless forming, stoneworking, wood, and plastic working. In 2017, about two-thirds of tungsten consumption (including recycled materials) were produced for hardmetals in the world. This paper briefly introduces the development of manufacturing technology on WC–Co hardmetals from three aspects: powder preparation, bulk densification, and performance characterization. Two special WC–Co hardmetals are also described: cobalt-enrichment zone (CEZ) hardmetals, and binderless hardmetals. Furthermore, the development prospects for manufacturing techniques of hardmetals are also presented in the end.


Hardmetal Ultrafine WC powder Ultrafine cobalt powder Extrusion press Liquid sintering Cobalt-enrichment zone hardmetal Binderless hardmetal 

1 Introduction

Hardmetals (also be named after cemented carbides) are one of the most successfully applied tool materials and the most widespread powder metallurgy products worldwide because of their outstanding combination of hardness and toughness compared to other tool materials, such as diamonds or high-speed steels [1, 2]. Many industries, such as automotive, aerospace, electronics, energy, mining, construction, and consumer durables, depend on tools made from hardmetals. Without hardmetal based tools, these industries could not succeed. After almost one century of development since 1923, hardmetals have come a long way from their original composition and structures to very sophisticated materials with tailored microstructures and complex compositions. García et al. [3] displayed the microstructure of 20 kinds of hardmetals only in one paper, and the actual types are much more than these. Nevertheless, two elements still dominate the hardmetal industry tonnagewise, as they have since the very beginning: tungsten and cobalt. In 2017, the total global production of tungsten is 1.05 × 108 kg, and 64% of it (about 6.7 × 107 kg of tungsten) was used to produce hardmetals [4], as shown in Fig. 1.
Fig. 1

Tungsten first-use in 2017

Varieties of hardmetals or cermets are developed and more than a dozen of them are commercially produced, but straight WC–Co hardmetals, which also include some grades with small amounts of grain growth inhibitors, still account for more than 80% of the market share [5]. In addition, WC–Co hardmetal manufacturing technology is also the basis manufacturing technology for other carbides. This paper reviews the progress of WC–Co hardmetal manufacturing technology from three aspects: powder preparation, bulk densification, and performance characterization. Finally, it also introduces the manufacturing technology of two other hardmetals: cobalt-enrichment zone hardmetals and binderless hardmetals.

2 Powder preparation techniques

Although WC–Co hardmetals can be made directly with mixing and reactive sintering of tungsten, carbon, and cobalt powders [6], the most commonly used method for preparing hardmetals is still using WC and cobalt powders as initial raw materials. This makes it easier to control the WC grain size and the resulting bulk material properties.

It can also be said that the history of hardmetals is the history of a steady widening range of available WC grain sizes for manufacturing. The main reason for this widening of the spectrum of WC grades is that the properties, such as the hardness, strength, toughness, wear resistance, thermal conductivity, etc. can be widely varied by means of the WC grain size, except for those variations achieved by the cobalt content. The finer the WC grain size is, the harder and more strength the material is, and contrariwise [7]. In the early 1990s, Fukatsu et al. [8] developed a micro-grained hardmetal with transverse rupture strength (TRS) of 5.0 GPa, using 0.2 μm WC powders as starting materials in their works. For current manufacturing technology, this strength is still high.

Although fine-grained hardmetals were attempted to develop in the late 1920s, it was not until the 1970s that high-performance fine-grained hardmetals were successfully developed, which is ascribed to that the ultrafine WC powders could be stably manufactured at that time. There are two techniques for producing ultrafine WC powders: production of WC from W or production of WC from WO3.

Producing WC from W is a conventional process. Bock et al. [9] believed that the conventional process, which was named after the conventional calcination–reduction–carburization process in his works, had a high potential and flexibility. R&D powders with scanning electron microscope (SEM) grain sizes of 0.15–0.20 μm were successfully prepared in full-scale and characterized. The tailor-made ultrafine WC powders pre-doped or undoped chromium carbide could be prepared with this process, depending on whether chromium oxide was added before the carbonization process [9, 10]. Figure 2 shows an SEM image of a pre-doped WC powder with the nominal particle size of 0.2 μm on the Chinese market, and it was produced with the conventional process.
Fig. 2

SEM image of a commercial superfine WC powder in China

Miyake et al. [11] and Asada et al. [12] reported a continuous direct carburizing process, other than the conventional process, to prepare the ultrafine WC powders. A continuous two-stage rotary carburization furnace was developed. In the first furnace, there is a nitrogen atmosphere and in the second one there is a hydrogen atmosphere was developed. The finest powders available (Cr3C2 doped or undoped) by this process were characterized by a close particle size distribution and a specific surface area between 3.0 and 3.5 cm2 g−1, and the real particle size was ~ 0.15 μm. Similarly, Wang et al. [13] also reported that ultrafine WC powders were prepared by a two-step process: carbothermic reduction of WO3 and then carbonization reaction of mid products. The single-phase WC powder with a particle size of ~ 0.20 μm could be obtained.

Another reason for continuous refinement of WC grains in ultrafine-grained hardmetals developed was grain growth inhibitors [14, 15]. Cr3C2, VC, TaC, NbC, TiC, Mo2C, ZrC, and HfC were studied as grain growth inhibitors of hardmetals, but only Cr3C2 and VC are proven to be the most effective and are most commonly adopted in the production of hardmetals [14, 15, 16, 17, 18]. In industrial production for fine-grained and submicron-grained WC–Co hardmetals, Cr3C2 is usually added separately as an inhibitor owing to brittleness of VC, but for ultrafine-grained hardmetals, a mixture of Cr3C2 + VC is still selected. The smaller the WC grain size is, the larger the total amount of grain growth inhibitors added in manufacturing. In general, when the mean WC grain sizes in commercial WC–Co hardmetals are ≥ 0.9 μm, the amount of grain growth inhibitor Cr3C2 added is less than 0.5 wt%. When the WC grain sizes are 0.6–0.9 μm, inhibitor Cr3C2 is added in an amount of 0.5–0.7 wt%. In addition, when the WC grain sizes are less than 0.5 μm, the total amount of inhibitor Cr3C2 and VC added is more than 0.7 wt%. Except for the WC grain size, the amount of grain growth inhibitors added also needs to be adjusted according to the content of metal binder. However, it is strictly prohibited to produce a brittle phase in sintered carbides due to the excess addition.

To date, the exact way of how inhibitors work is not fully understood, and controversially discussed in some respects. The most consistent opinion of grain growth inhibition is that Cr/V/Ta dissolute within the cobalt binder and position the growth sites on the tungsten carbide crystal. The dissolution of Cr/V/Ta will limit the ability for W to dissolve within cobalt, leading to a decreased solution and precipitation of W during liquid-phase sintering. In addition, the position of growth sites can hinder tungsten atoms to be integrated in the lattice on WC crystals. Pötschke et al. [19, 20] considered that the grain growth inhibition was mainly governed by the changed solution and precipitation processes during liquid-phase sintering, but during solid state sintering, it was also governed by the change in chemical processes which happen during heating-up to sintering temperatures. Pötschke et al. [19, 20] and Richter et al. [21] also reported an important information that grain growth inhibitors could perform obvious effects as long as their particle sizes were below 2.0 μm, since the inhibitors dissolve and disappear completely in the solid-phase sintering. Whether adding the grain growth inhibitor before WC carbonization or during the preparation of ready to press (RTP) powders, they both could be uniformly mixed with WC powders and worked in sintering [10]. For which method to be used in manufacturing, it mainly depends on the usage habits. The latest research on grain growth inhibitors was the effect of Cr3C2, VC, TaC, NbC, and TiC on the interface structure [22, 23].

Another process for preparing ultrafine-grained hardmetals is using WC–Co composite powders as starting materials. The WC–Co composite powders are two-phase powders prepared by the thermochemical synthesis, spray drying or sol–gel method. These powders contain ultrafine WC and cobalt grains with a uniform distribution on the micrometer scale. Figure 3 shows an SEM image of a WC–Co composite powder prepared in the laboratory, and its morphology is significantly different from that of tungsten carbide powders (Fig. 2). This technology can avoid the uneven mixing of raw materials during ball milling and is beneficial to manufacture hardmetals with uniform microstructures in theory.
Fig. 3

SEM image of a WC–Co composite powder prepared in the laboratory

McCandlish et al. [24] and Kear et al. [25] first reported composite powder technology in the early 1990s. Ouyang et al. [26] used tungsten salts and cobalt salts as starting materials to prepare WC–Co composite powders by tungsten–cobalt liquid-phase composite, spray drying, and fluidization thermochemical conversion process. The average grain size in the composite powder was less than 100 nm. The hardness of the sintered bulk material was greater than 92.0 HRA, and the TRS was greater than 3300 MPa [26]. Yi et al. [27] used a direct reduction and carbonization method to prepare WC–Co composite powders with a particle size of 100–300 nm, but the final powders contained W2C, Co3W3C, Co6W6C, and other carbon-deficient phases. Liu et al. [28, 29, 30] and Wei et al. [31] reported a rapid route for synthesizing WC–Co composite powders by in situ reduction and carbonization reactions of metal oxides and carbon. After sinter-hot isostatic pressing (HIP) or spark plasma sintering (SPS), the WC–Co hardmetals prepared with this WC–Co composite powders exhibited homogeneous and ultrafine microstructures [30, 31]. The WC grain size distribution in the bulk material was in a narrow range of 150–350 nm and its average grain size was 240 nm [30]. The sintered hardmetal also achieved good mechanical properties with the hardness of 92.7 HRA, the fracture toughness of 10.8 MPa m1/2 and the TRS of 3860 MPa [31].

Another key raw material for hardmetals is cobalt powders. With the development of ultrafine- and even nano-grained hardmetals, ultrafine spherical cobalt powders were desired for making a homogeneous and densified microstructure without cobalt-related flaws, such as pores, cobalt lakes, cobalt-deficient, or cobalt-free zones [32, 33].

In powder metallurgy industry, cobalt powders are mainly produced by a decomposition and hydrogen reduction technique using cobalt oxide (Co3O4) powders, cobalt carbonate (CoCO3·xH2O), or even cobalt oxalate (CoC2O4·2H2O) as raw materials in a push-type furnace at relatively low temperatures, generally below 800 °C [34]. However, it is not easy to industrially prepare ultrafine spherical cobalt powders with good performances by the conventional decomposition and hydrogen reduction technology. The cobalt nuclei usually nucleate and grow to form a short string or dendritic structure along the short fibrous particle of CoC2O4·2H2O, and thus, firm agglomerations are formed when several short-string cobalt particles grow or bond together if they are next to each other tightly in the reduction reaction. Similarly, using CoCO3·xH2O powders as the raw material, the Fisher sub-sieve size (FSSS) of cobalt powders obtained by the conventional decomposition and hydrogen reduction technique is usually larger than 0.8 µm and the main problem is that the coarse firm agglomerations still occurs in reaction.

Ultrafine spherical Co3O4 powder is the best raw material for the preparation of ultrafine spherical cobalt powder. In theory, when the decomposition reaction is carried out in a fixed bed, the morphology of the Co3O4 can still inherit the short fiber structure of CoC2O4·2H2O. That is, the Co3O4 particles can easily grow separately to form a short fiber structure due to an exothermic reaction. The short fiber is composed of ultrafine spherical Co3O4 particles with a size of ~ 0.1 to 0.2 μm, and weak adhesion between Co3O4 particles makes them arrange short strings along to the fiber. This special structure further explains why the short strings or dendritic cobalt powders are easily obtained by decomposing and reducing CoC2O4·2H2O powders in a fixed bed. If the weak adhesion between Co3O4 particles could be broken in situ before strong bonds were formed, ultrafine Co3O4 powders with the size less than 0.2 μm might be obtained. Wu et al. [35, 36] invented a continuously dynamic-controlled combustion synthesis (CDCCS) technology, which is a new low energy-consumption industrial production technology for preparing spherical Co3O4 powders. In the CDCCS technology, the weak adhesion among Co3O4 particles can be broken by an airflow, and the heat can also be homogenized by the airflow to avoid local overheat and to impede the adhesion among spherical Co3O4 particles during the decomposition of CoC2O4·2H2O. Therefore, ultrafine spherical cobalt powders can be prepared by decomposition of CoC2O4·2H2O in a fluidized bed and followed by hydrogen reduction of spherical Co3O4 in a fixed bed. The average particle size of cobalt powders obtained is less than 0.6 μm, and the cobalt powders possess a narrow particle size distribution, good dispersity, and excellent fluidity, as shown in Fig. 4.
Fig. 4

Ultrafine spherical cobalt powders produced by hydrogen reduction of ultrafine spherical Co3O4 powders

3 Bulk densification techniques

Only by press forming and sintering can WC–Co powders become hardmetal products with certain shapes and excellent mechanical properties. Almost all carbide green products are formed using the die pressing technology or extrusion forming technology, which has been gradually developed in the last 30 years [37, 38, 39].

Most cross-sectional rods can be made with the extrusion process. The length of the rods is almost unlimited thereby and depends only on the type of the extruder used. In a German factory, the longest sintered extrusion was astonishingly 2.1 m long [39]. There are two types of extruders used, piston extruders and screw extruders with an obvious distinction is made between them. In screw extruders, lubricated powder is initially loaded into a horizontal screw feeder, which in turn feeds the near-vertical extrusion chamber. In this way, the screw extruders can work continuously. However, if piston extruders want to work uninterruptedly, there must be enough lubricated powder contained in the cylinder in which the piston works.

The extrusion presses can rotate the pieces as they emerge, producing extruded carbide drill blanks with preformed ducts of precisely defined helix angles. The number of spiral cooling ducts could be up to five [39]. Several smaller diameter-sized carbide rods could be produced with a more complicated machine, in which a single extrusion chamber was still employed, but with the same several diameter-sized dies in-line [39].

Before extruding, the powder mixes are tested with a capillary viscometer for obtaining information about the homogeneity of the mixes and their viscosity, which is important to successfully predict extrusion processes. Formisano et al. [40] reported that a rheological analysis was made for the case of cutting tools with inner helical holes, and a correct die was manufactured from additional finite element method (FEM) analyses. Pores, cracks and other defects, as well as otherwise invisible coolant holes in sintered carbide rods, can be visualized by ultrasonic.

Hardmetals with fine WC grains or the low metal binder content usually own high hardness, which is very suitable for making metal cutting tools for taking responsibility for most subtractive manufacturing. However, their hardness is very useful in service, making it difficult and expensive to adjust the size or shape of the sintered hardmetal component, resulting in increasingly complex dies for cutting inserts or other products. Therefore, people began to develop the additive manufacturing (AM) technology of hardmetal parts 30 years ago [41, 42]. However, AM of hardmetals is still in the early stages of development, whose major problem is the high melting point of the carbide phase. If an attempt is made to replicate a conventional composition of a WC–Co hardmetal, the time available is much short to achieve normal liquid-phase sintering. Cobalt can melt, but in a typical mixture, it does not have time to wet most of the carbide particles. However, if the power was sufficiently increased to melt the carbide grains, much of cobalt would probably evaporate [43, 44]. WC melts at 2770 °C and cobalt boils at 2870 °C, but it is difficult to control the temperature sufficiently to satisfy the one and avoid the other in a high-intensity melting beam.

Typical AM such as selective laser sintering or melting and electron beam melting has the advantage of direct forming and densification in one step. However, unfortunately, nearly 100% of dense WC–Co samples are not possible, due to local high-energy input, chemical imbalance phases such as W2C, carbon, or η phase (a substoichiometric carbide phase) and locally different cobalt contents occurred, resulting in cracks and internal residual stresses [42].

Another two-step (printing + sintering) AM technology, either powder bed-based techniques such as binder jetting or suspension-based techniques such as thermoplastic 3D printing, may be feasible. The WC–Co particles are not sintered, but selectively glued together by an organic binder using a special print head. Therefore, the 3D-printed components are green components, and their green density is mainly based on the powder density of the starting materials. After printing, the organic materials used during printing must be removed during a dewaxing step, and then, the components must be sintered according to a conventional sinter cycle [45, 46, 47]. In principle, this is near conventional hardmetal manufacturing. Companies with experience in metal injection molding (MIM) or extrusion forming should have advantages in this area, where they can put their skills and expertise to good use.

MIM or powder injection molding (PIM) is a kind of near-net-shape powder metallurgy forming process, which has great advantages for the production of difficult-to-machine hardmetal components with complex shapes. Early researches yielded positive results [48, 49, 50] and tooling products using the MIM technology emerged on exhibitions or market [51, 52]. However, frankly, until now the MIM technology is not widely used in hardmetal manufacturing. First, the product mold is too expensive. Second, the much high content of a polymer binder makes it hard to control the carbon content and porosity in sintered bulks, which possess poor strength [53, 54].

In the early days, most of the hardmetals were sintered in a reducing atmosphere, such as hydrogen or ammonia decomposition gas. However, the control of carbon in the hardmetal sintered in reducing environment was very difficult, which made the performance and dimensional control inconsistent. Therefore, from the early 1970s and the late 1980s, the hardmetal manufacturing gradually adopted the vacuum furnace and sinter-HIP furnace in the whole world, respectively. Other processes for sintering hardmetals are explored, such as microwave sintering or SPS, but not established in the industry.

The fast known sintering time of hardmetal was only several seconds with electro-sinter-forging [55]. However, WC–Co hardmetal sintering usually lasts for more than ten or even dozens of hours in the industrial vacuum furnace or sinter-HIP furnace, due to several key stages, such as pressing binder debinding (dewaxing), oxide reduction and removal, solid-phase sintering, liquid-phase sintering and cooling, which need to be completed [56], as shown in Fig. 5.
Fig. 5

Temperature ranges for densification stages, together with a corresponding schematic description of microstructure evolution and fracture SEM photos of 1.5 μm WC-6 wt% Co hardmetals, which were sintered at 1000–1400 °C.

Reproduced with permission from Ref. [57] Copyright 2016 Elsevier

Almost all of those lubricants’ adding for pressing need to be removed in a “dewaxing” process with the temperature of under 800 °C before sintering. Normally, dewaxing is accomplished by heating the compacts in a suitable atmosphere such as hydrogen or argon. By carefully controlling the temperature cycle and carrier gas velocity, lubricants can be almost completely removed at a suitable rate without unexpected pyrolysis or damaging the compact [58]. Before dewaxing, the paste extrusion rods, bars, or tubes often need drying for several days to volatilize organic solvents [59].

Freytag and Exner [60] studied the sintering behaviors of WC-12 wt% Co hardmetals and found that no shrinkage of samples was observed below 800 °C, but the elimination of the stress from ball milling and the reduction of cobalt oxide on the powder surface. The shrinkage between 800 and 1100 °C is limited, and WC itself is partially dissolved into the binder which is spread on the surface of WC. Continuing to heat up, but the sinter temperature is still below the eutectic temperature, which can cause a great proportion of overall shrinkage. It is emphasized that increasing the diffusion of cobalt-coated carbide grain boundaries should play a decisive role in this solid-phase sintering. When the cobalt binder phase melts, some of the carbides dissolve into the liquid phase and ultimately grain rearrangement and pore reduction result in full densification of hardmetals.

Surface oxides are usually present in industrial RTP powders. These oxides are reduced in a temperature range of 650–850 °C with a maximum reduction temperature of ~ 750 °C [61, 62]. In addition, for doped WC materials, there is an additional peak of CO/CO2 released at 950 °C [61]. It is pointed out that the most significant carbon loss caused by the reaction with the surface oxygen of the carbide during sintering occurs in a temperature range of 700–1000 °C [60]. Wang et al. [62] reported that in the sintering process of ultrafine-grained WC–Co hardmetal containing VC and Cr3C2, the reduction reaction of tungsten oxides exists in a range of 700–900 °C, and the reduction reaction temperature ranges of vanadium oxides and chromium oxides are 900–1100 °C and 1200–1400 °C, respectively.

The results of the thermogravimetry, differential thermal analysis, mass spectrometry, and thermal expansion further indicated that the onset of shrinkage occurred after desorption of oxygen and oxide reacts with carbon in the mix to form CO and CO2, that was, when all of the powder surface oxides were reduced, the sintering shrinkage began [63]. Göthelid et al. [64] observed the spreading behavior of cobalt on WC used scanning tunneling microscope (STM) and also pointed out that there was a strong correlation between an oxide film and poor wetting conditions. However, these views were questioned by Macedo et al. [65], whose research results showed that the wetting of cobalt, nickel and iron could also occur on the surface of tungsten oxides. Therefore, there is no correlation between the initial point of sinter shrinkage and oxide reduction, although they appear in order.

Further densification during hardmetal sintering occurs through grain rearrangement followed by dissolution–precipitation. When the amount of liquid phase is high and the carbide grain size is small, full densification can be achieved quickly. A high density could also be obtained at the initial stage or before liquid-phase formation in hardmetal sintering, which was an important difference of liquid-phase sintering between WC–Co hardmetals and other materials [66].

Froschauer et al. [67] in situ observed the evolution of hardmetal microstructures and demonstrated that the rearrangement was the most important process of densification. Haglund et al. [68] reported that for WC-10 wt% Co hardmetals, densification in the solid-phase sintering could be seen as a WC rearrangement process, such as the first stage of liquid-phase sintering. When the liquid phase occurs, there was no significant change in the densification rate. However, the above conclusions are correct for hardmetals with conventional WC grain sizes, but not applicable to ultra-coarse hardmetals.

Carbide particles are partially encapsulated by cobalt, and fine carbide debris may be embedded in the cobalt during ball milling and this combination is critical for initial sintering. Ball milling and cold pressing also produce large deformed contact areas with high inner stresses, which can be released by creep deformation or other transport mechanisms during heating. Cobalt-rich zones with aggregated tungsten carbide particles are formed during ball milling, and local densification begins in these zones.

When the sinter temperature was higher than 650–710 °C, the surface of the tungsten carbide were wetted by the dispersed metal binder particles which were similar to the spreading of viscous material [69, 70]. The cobalt immediately spreading on the interfacial surface of tungsten carbide after oxide reduction could be seen as the beginning of densification. The diffusion of cobalt reduced the surface free energy, and the WC-gas and Co-gas interfaces were changed to become the WC–Co interface by the spreading of cobalt, which could reduce the interface energy by ~ 1/5 [69]. Cobalt has a high spreading rate on tungsten carbide during heating, and cobalt spread onto the surface of the tungsten carbide was only a metal film rather than a block. The increase of WC dissolution due to cobalt spread on the WC surface would further reduce the viscosity of the tungsten carbide-metal contact surface.

The smaller the particle size is, the larger the powder surface area in per unit volume of a compact is, and the higher the surface free energy is. Therefore, the earlier the beginning of sintering of fine-grained hardmetals is, the higher the shrinkage proportion during solid-phase sintering is. In addition, for finer tungsten carbides, ball milling and mixing introduce more structure defects, resulting in higher internal stresses, and shorter distances from the diffusion source to the diffusion well, and, therefore, lower viscosity. When the maximum sinter shrinkage occurred at a lower temperature, the sinter shrinkage curves become wider and flatter. For ultrafine and nano-grained hardmetals, about 90% of densification happened during solid state sintering [70]. However, for ultra-coarse-grained hardmetal sintering, since the powder surface area in the unit volume of the compact was small and the total surface energy was low, it was necessary to add fine-grained tungsten carbide powder to promote the sintering [71].

In addition to ultrafine-grained hardmetals, solid-phase shrinkage of other carbides can be divided into two stages by kinetics. The temperature of the first stage of densification is typically less than 1100–1200 °C, which is considered to be uniform shrinkage caused by the spread of the binder phase and creep. While in the second one, although densification is still dominated by the metal binder creep, asymmetric forces become strong enough to cause rotation, distortion, and deflection between particles. In the first stage, the intrinsic stress of the binder phase provides the force to drive densification. While in the second phase, the binder phase is heated, and the intrinsic stress is negligible compared to the capillary force. In addition, the binder phase is almost free to creep in the first stage, while in the second stage, the creep is limited by the huge WC–Co-phase interfaces. The first stage of densification is closely related to the process before sintering, and the second stage is weakly related to the early process and is strongly influenced by the WC grain size. For ultrafine-grained materials, solid-phase sintering was more complicated and could not simply be divided into two stages [70, 72].

Densification in liquid-phase sintering is achieved by shape adjustment of grains. For hardmetals, it is specifical due to increased contact between tungsten carbide grains, grain boundary flattening because of dissolution of small grains and precipitation on large grains, and grain coarsening because of grain boundary migration and dissolution–precipitation (Ostwald ripening). A WC grain coarsens and forms a trigonal prism crystal with truncated corners. However, such a trigonal prism crystal shape does not necessarily coincide with the aggregate structure formed at the beginning, and if the amount of the binder phase is sufficiently large, it may cause rearrangement again. In practice, most WC–Co materials can be fully densified by rearrangement and grain shape adjustment during liquid-phase sintering.

Grain growth inhibitors, such as VC and Cr2C3 shifted the shrinkage onset temperature to high temperatures, reduced the rate of shrinkage in solid-phase sintering, and lowered the liquidus point of the sintered system [62, 73]. Whether carbon or tungsten is added to the fine-grained hardmetals, the shrinkage rate is reduced and the effect of tungsten is stronger than that of carbon. However, there is little effect on the coarse-grained tungsten carbides.

Carbon has a great influence on the sintering of WC–Co hardmetals. The change of the carbon content could change the eutectic temperature of the WC–Co system from 1298 to 1368 °C [73], which significantly changed the progress of solid-phase sintering and liquid-phase sintering (Fig. 6). In modern production lines, total carbon of WC-10 wt% Co hardmetals should be controlled in a range of ± 0.015% by weight. The carbon content could also change wetting angles of liquid cobalt on the surface of tungsten carbide, with a high carbon content of 0 and a low carbon content of 15° [74]. Therefore, the well-known phenomenon of cobalt migration from carbide regions with high carbon contents into regions with low carbon contents was most probably caused by the variance in capillary forces related to the different wetting angles of WC by liquid metal binders with various carbon contents [74, 75].
Fig. 6

Sintering densification curves of WC–Co hardmetals with different carbon contents. The FSSS particle sizes of starting WC powders were 8.0 µm

Cobalt capping was a term that refers to local thin layers of cobalt which were occasionally observed on the surface of sintered WC–Co parts when they came out of the furnace. Cobalt capping, which is usually believed undesirable and needs to be removed, was also considered to be related to the carbon atmosphere in the sintering furnace during cooling [76, 77, 78, 79, 80].

With the change of parts size during sintering, such as shrinkage and creep deformation, the microstructure is also changed (Fig. 5). The grain size distribution of hardmetals is substantially dependent on the particle size distribution of the starting powders and the corresponding ball milling conditions. For discontinuous grain growth, large particles may act as nucleation, which broadens the grain distribution.

The shrinkage during sintering is ~ 20%, making it impossible to suppress all movements and thus to prevent even the slightest warpage. Then, additional vacuum furnaces need be installed for straightening sintered parts. In the sintering process, the local deformation behaviors of hardmetal products at specific temperatures is still not clear, which makes it difficult to control or reduce the sintering deformation of hardmetals. Raether et al. [81] and Baber et al. [82] reported a thermooptical measurement system for components or samples with a maximum size of ~ 40 mm, which is operated in a controlled atmosphere or vacuum. Dimensional changes are detected optically with a resolution of 0.3 µm and the maximum temperature can be up to 2400 °C. This system can be used to study the deformation of hardmetals during sintering in future.

4 Performance characterization and testing

When the hardmetal products are taken out of the sintering furnace, it is necessary to measure the properties such as metallography, density, hardness and TRS in time, according to the corresponding International Organization for Standardization (ISO) documents. However, an important aspect to consider when manufacturing hardmetals is the control of carbon content, which is related to many performances of hardmetals, such as magnetism, density, and hardness, as shown in Fig. 7.
Fig. 7

Properties of hardmetals varying with carbon content: a magnetic saturation and coercivity; b density; c hardness

The variation of the carbon content in the cobalt binder has a relatively narrow range, the deviation of which can result in the formation of a third phase. When at low carbon contents, a η phase, which is complex carbide (WxCox)C, can be formed. The η phase tends to increase the brittleness of hardmetals, especially if it precipitates as large dendrites. Conversely, if a hardmetal with a high carbon content was produced, free carbon may precipitate in the form of graphite and also reduce the mechanical properties of the material. Figure 8 shows two metallographic photos of WC–Co hardmetals, containing η phase and free carbon, respectively. Therefore, the carbon balance, which actually depends on the overall composition of the materials, is always strictly monitored and maintained during manufacturing hardmetals to avoid the occurrence of the third phase.
Fig. 8

Metallographic photos of WC–Co hardmetal, containing a η phase and b free carbon, respectively

Measurement of the magnetic properties is a well-established method for determining whether the final product is suitable for an acceptable carbon window, which ensures optimum performances of the hardmetal. This non-destructive quality control is possible thanks to the ferromagnetism of cobalt. Magnetic saturation is used as an indirect, fast, and reliable method of measuring the carbon content of sintered hardmetals. An advantage of this method is the linear relationship between the carbon content and the magnetic saturation values in the region of interest. The carbon content accuracy can be estimated to be 0.01 wt% by measuring the magnetic saturation of the samples prepared under closely comparable conditions. As the amount of solid solution tungsten in cobalt decreased, the magnetic saturation of the sintered hardmetal increased, and vice versa [83]. Therefore, the magnetic saturation in the two-phase region indicates the amount of tungsten dissolved in the binder phase. In general, hardmetals with low magnetic saturation are selected in metal cutting, while the ones with high magnetic saturation are selected in construction and stone working.

Another crucial magnetic property that controls the quality of sintered hardmetals is the coercive force (Hc). Coercivity is a non-destructive evaluation of the microstructure of hardmetals, such as the degree of sintering, cobalt distribution, and WC grain size. The coercive force is inversely proportional to the WC grain size of hardmetals, which means that high coercive force values indicate a fine-grained microstructure and vice versa. It is considered that Hc increases linearly as the mean WC grain size decreases. However, in fact, coercivity depended on the fine fraction of WC grains [84]. The significances for coercivity as a quality measurement, used to approximate the mean WC grain size in hardmetals, were decided by the ability to predict the WC grain size distribution, for example, the stability of the raw material and the production process. Therefore, coercive force values can also be used to optimize milling conditions. In addition, the coercive force is also affected by the cobalt content and carbon content in hardmetals. As the cobalt content or carbon content increased, the coercive force value decreased [85].

During the last years, significant progress has been made in micromechanical testing such as nanoindentation, micropillars, and micro-cantilever testing. WC–Co hardmetals with outstanding mechanical properties are geometrically complex composites consisting of two interpenetrating networks of the constitutive ceramic and metal phases. To improve the performance of these materials, a deeper understanding of the mechanisms controlling hardness, strength and toughness is required. These mechanisms are primarily dependent on their microstructural characteristics. Micromechanical testing is a quantitative method for measuring local mechanical properties of hardmetals and seems to be a promising technique to study these mechanisms.

Liu et al. [86] reported that the nano-grained hardmetal has a low indentation modulus, which could be attributed to the large interface area and high fraction ratio of the hexagonal close-packed (hcp) cobalt phase caused by the rapid heating and cooling process during SPS. Roa et al. [87] attempted to combine the massive nanoindentation, statistical analyses, and implementation of a thin film model for deconvolution of the intrinsic hardness and flow stress of the metallic phase. In addition, the yield stress values as a function of the cobalt mean free path resulting in a Hall–Petch strengthening relationship with a slope of 0.98 MPa m1/2 were plotted. This resulting value pointed out the effectiveness of WC–Co-phase boundaries as powerful obstacles to slip propagation. Duszováet al. [88] revealed that the nanohardness of the cobalt binder was ~ 10 GPa and the average nanohardness of the WC basal planes and prismatic planes were 40.4 ± 1.6 GPa and 32.8 ± 2.0 GPa, respectively. The corresponding average values of the indentation moduli for WC basal and prismatic planes were 674 ± 14 GPa and 542 ± 34 GPa, respectively. Csanádi et al. [89] reported that measured nanohardness values were ~ 1.6 times higher on basal planes (H = 43 ± 0.8 GPa) than that of on prismatic planes (H = 28 ± 1.0 GPa), and observed sink-in and pile-up effects in the case of basal and prismatic orientations, respectively. Obviously, the results of Refs. [88, 89] are different. Roa et al. [90] also assessed local hardness and the elastic modulus through high-speed massive nanoindentation and subsequent statistical analyses, reporting the strong correlation between mechanical properties and compositional phases. The test results of micropillars testing [91, 92, 93] and micro-cantilever testing [94, 95] were scattered and more micromechanical testing still needs researches in the future.

5 Cobalt-enrichment zone hardmetals and binderless hardmetals

By developing a cobalt-enriched and cubic-carbide-free surface layer, the surface-region composition within hardmetals could be modified to provide added toughness for coated cutting tool substrates. These hardmetals, which can increase crack propagation resistance in cutting tools, are named after cobalt-enrichment zone (CEZ) hardmetals. It was recognized that nitrogen diffused outwards and titanium diffused inwards, so that a CEZ was formed. Therefore, typically, such modifications were prepared by establishing nitrogen gradients in the hardmetal during sintering through the additions of cubic nitrides or nitrification of the cubic carbides.

CEZ hardmetals were invented in early 1980s which are still widely used in steel turning insert substrates, and one of them is certainly the KC850 grade of Kennametal [96]. CEZ hardmetals have two important structural properties: the thickness of the cobalt-rich zone and the cobalt-phase fraction distribution in the cobalt-rich zone. In general, CEZ hardmetals with excellent properties possess a thickness of more than 15 μm of the cobalt-rich zone and the maximum content of cobalt in the zone is more than 160% of the nominal content. The above two structural properties are affected by many factors, such as the Ti (C, N) content, Co content, WC particle size, [N]/[C] ratio, sintering temperature, nitrogen pressure, and even the conversion timing of the nitrogen pressure [97, 98, 99, 100]. Flexibly adjusting these process parameters could prepare the desired cutting tool gradient substrate. In addition, a CEZ hardmetal was also prepared by employing TiN- or Ti (C, N)-free initial powder mixture and nitridation–denitridation sintering [101]. Those research results offered guidance for forward adjusting the composition of the starting mixture and sintering procedure in a positive direction.

The gradient zone formation of hardmetals were simulated and the simulation results showed a good agreement with the experimental results [102, 103, 104, 105]. Comparisons between the calculated and measured results showed that most of the experimental data could be well improved by the currently available databases.

Binderless hardmetals, another typical hardmetals, are WC-matrix ceramic materials with WC as a main hard phase and < 0.5 wt% Co, Ni, or other metals as binders. Although WC-matrix ceramic materials were researched as early as the 1980s [106], it was not until 1992 [107] that the commercial products were used. The detailed research on the binderless hardmetal was mostly concentrated in the past 20 years, so it is a relatively new material. Compared with conventional hardmetal materials, binderless hardmetals have a good abrasion resistance, high temperature deformation resistance, corrosion resistance, low thermal expansion, and good thermal conductivity. Moreover, it can be widely used as an ultra-high pressure waterjet nozzle, high precision die, high wear resistance seal ring, electronic packaging materials, and heavy-load sliding seal wear-resistant parts, etc. Binderless hardmetals also have potential applications in drawing die and cutting tools.

There are three main difficulties in preparing binderless hardmetals with ultrafine grains:
  1. 1.
    Precise carbon and oxygen control. The limit of the two-phase region (WC and Co) in WC-0.2 wt% Co is nearly 50 times smaller than that in WC-10 wt% Co (Fig. 9). To prevent the formation of η phases or free carbon, the following three measures are the key points: (1) calibrating accurately the oxygen and carbon contents in starting powders; (2) controlling strictly the change of the oxygen content in the preparation process; and (3) the stable and controllable carbon atmosphere in the sintering furnace at high temperatures.
    Fig. 9

    Phase diagrams of WC-10 wt% Co and WC-0.2 wt% Co, respectively.

    Reproduced with permission from Ref. [108] Copyright 2018 China Academic Journal Electronic Publishing House

  2. 2.

    Sinter densification to get a pore-free microstructure. To obtain near-net-shaped products, first, the net-shaping green compact can be prepared by the pressing/extrusion/injection/cold isostatic pressure/3D additive manufacturing technology, and then the sintered bulk with a small amount of nano-pores can be obtained by vacuum sintering at temperatures of above 1700 °C. Finally, the nano-pores in the above sintered bulk can be eliminated completely by HIP at temperatures of above 1400 °C.

  3. 3.

    WC grain refinement. Based on the sintering temperature and time, sintering densification and inhibiting grain growth are contradictory, so the optimal sintering process and adding grain inhibitor are necessary [109].

Figure 10 summarizes the relationship between hardness and toughness of reported binderless hardmetals [108]. The carbide additives MexCy (Me = Ti, Ta, Nb, Mo, Zr, Cr, V, Si, etc.) could react with WC in solid solution, promote sintering densification, improve oxidation resistance of binderless hardmetals, and also inhibit the growth of WC grains. The oxide additives MexOy (Me = Zr, Al, Mg, etc.) could improve toughness of binderless hardmetals, so as to promote its potential application in cutting tools. The boride additives MexBy (Me = Ti, Zr, Cr, V, etc.) could react with W, Co, Ni, and Fe to form ternary borides with high hardness, and promote sintering densification and the mechanical performance of binderless hardmetals, which is one of the future research directions. In addition, the work mechanisms of rare earth on microstructures and properties of binderless hardmetals are not clear so far.
Fig. 10

Relationship of hardness and toughness of reported binderless hardmetals. 

Reproduced with permission from Ref. [108] Copyright 2018 China Academic Journal Electronic Publishing House

6 Summary and outlooks

Hardmetals are still a growing market. Consumption continues to expand, from an annual world total of 1067 kg in 1926 to ~ 85000 t in 2018 [110, 111]. As the world economy grows, hardmetals continuously play a crucial role in technological advancement.

In particular, the introduction of coated hardmetal grades in 1969 strengthened their position, as the corresponding coatings can now be tailored to the respective processing problems (materials, machining conditions), which leads to a significant increase in manufacturing productivity [112, 113, 114, 115, 116, 117].

Powder metallurgy (including MIM and AM) also provides a simple and flexible approach to mass production of components and tools with complex geometries, improved designs, and modern technologies. For example, pressure-aided sintering significantly promotes the high reliability of materials during applications, which is a prerequisite in automated manufacturing. The formation of graduated materials, modern joining techniques and laser honing further broadens the application range of hardmetal tools and the development of polycrystalline diamond which is integrally sintered into a tough hardmetal substrate under a high pressure, as a good example of recent innovation.

New technologies and new materials can require new tooling solutions, and hardmetal tools can provide a cost-effective option due to their attractive properties. Modern recycling technologies and the efficient collection system of scrap materials will contribute in this regard, driven by both the price and the need to maintain and conserve natural resources. In the long run, recycling will inevitably become a key strategic factor for sustained economic growth, and the respective recycling strategies of hardmetals are on the line [118, 119].

After developing for more than 90 years, hardmetals are produced automatically on an industrial scale and are almost the irreplaceable material for modern manufacturing. So, you think we know all about WC-based hardmetals? The answer is NO. It was advised that advanced high-resolution characterization methods, such as baseline structures [3D, optical and electron methods, atomic probe microscopy (APM)], deformation and damage development, length scale dependence simulation methods of properties, heterogeneity testing method, should be developed to understand the hardmetals more in-depth [120]. Both tungsten and cobalt are scarce, and more and more cobalt are used in production of batteries, which has a potential impact on the supply of raw materials of hardmetals. Therefore, development of new hard materials, such as alternative metallic binder (beyond Co) and carbides (beyond WC), is a continuing research hotspot.



The work was financially supported by the Major Special Projects of Fujian Science and Technology Plan (Grant No. 2017HZ0001-1). The authors are grateful to Prof. Yang Mingchuan of SLU for his WC–Co composite powder.


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Copyright information

© The Nonferrous Metals Society of China 2019

Authors and Affiliations

  1. 1.School of Materials Science and EngineeringBaise UniversityBaiseChina
  2. 2.China National R&D Center for Tungsten Technology, Xiamen Tungsten Co., LTD.XiamenChina

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