A review of cemented carbides for rock drilling: An old but still tough challenge in geo-engineering

A review of cemented carbides for rock drilling: An old but still tough challenge in geo-engineering

Int. Journal of Refractory Metals and Hard Materials 39 (2013) 61–77 Contents lists available at SciVerse ScienceDirect Int. Journal of Refractory M...

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Int. Journal of Refractory Metals and Hard Materials 39 (2013) 61–77

Contents lists available at SciVerse ScienceDirect

Int. Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM

Review

A review of cemented carbides for rock drilling: An old but still tough challenge in geo-engineering Xiaoyong Ren a, Hezhuo Miao b, Zhijian Peng a,⁎ a b

Key Laboratory on Deep GeoDrilling Technology of the Ministry of Land and Resources, School of Engineering and Technology, China University of Geosciences, Beijing 100083, PR China State Key Laboratory of New Ceramics and Fine Processing, Tsinghua University, Beijing 100084, PR China

a r t i c l e

i n f o

Article history: Received 2 May 2012 Accepted 7 January 2013 Keywords: Cemented carbide Nanostructured Functionally graded materials Co-free Geo-engineering

a b s t r a c t Cemented carbide is an old and well-known WC-based hardmetal, which has been widely applied in geo-engineering as drill buttons and various wear-resistant parts. In order to extend the service life of cemented carbide components and enhance their efficiency for rock drilling under various conditions, the recent research efforts have focused on their failure mechanisms and developing nanostructured, functionally graded and Co-free cemented carbides. With the advance in synthesizing nanosized powders and advent of electric field assisted fast sintering techniques, the consolidation of nanostructured and Co-free cemented carbides and even pure WC materials has been possible; and because of their high hardness and wear resistance, they are much promising in geo-engineering drilling. Functionally graded cemented carbide provides a combination of high wear resistance and toughness in a single component, which is also much favorable for geo-engineering drillers. In addition, by replacing the binder phase Co with Ni or carbide binder, and even without binder phase, the corrosion and oxidation of the resultant materials can be significantly improved without considerable deterioration of fracture toughness. © 2013 Elsevier Ltd. All rights reserved.

Contents 1. 2.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Rock drilling and failure mechanisms of cemented carbides in geo-engineering 2.1. Rock drilling of cemented carbide in geo-engineering . . . . . . . . . . . . 2.2. The failure mechanisms of cemented carbide drill buttons . . . . . . . . . . 3. Developing nanostructured cemented carbide hardmetals . . . . . . . . . . . . . 3.1. Synthesis of nanosized WC/WC-Co powders . . . . . . . . . . . . . . . . 3.2. The sintering of nanostructured cemented carbide . . . . . . . . . . . . . 3.3. Mechanical properties of nanostructured cemented carbide . . . . . . . . . 4. Developing functionally graded cemented carbide . . . . . . . . . . . . . . . . . 4.1. The sintering of functionally graded cemented carbide . . . . . . . . . . . 4.2. Key factors affecting the formation of graded cemented carbide . . . . . . . 5. Developing Co-free cemented carbide . . . . . . . . . . . . . . . . . . . . . . . 5.1. WC-Ni cemented carbide . . . . . . . . . . . . . . . . . . . . . . . . . 5.2. Carbide/oxide-bonded cemented carbide . . . . . . . . . . . . . . . . . . 5.3. Pure WC hardmetals . . . . . . . . . . . . . . . . . . . . . . . . . . . 6. Conclusions and outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

⁎ Corresponding author at: School of Engineering and Technology, China University of Geoscience at Beijing, Beijing 100083, PR China. Tel.: +86 10 82320255; fax: +86 10 82322624. E-mail address: [email protected] (Z. Peng). 0263-4368/$ – see front matter © 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijrmhm.2013.01.003

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1. Introduction Cemented carbide (also called WC-based hardmetal) is an old material, since it was invented in 1923 [1]. However, almost after

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its invention, it has been used as tool materials in geo-engineering for rock drilling, mineral cutting, gas and oil drilling, tunneling projects and the like, because of its unique combination of excellent mechanical properties such as high hardness and high toughness [2–5]. Its relatively high wear resistance and low cost also protrude it among the other counterpart materials [5], which make it favorable for various wear-resistant parts in geo-engineering. For example, in the world-wide known China Continental Scientific Drilling Project, drilling bits equipped with cemented carbide buttons as shown in Fig. 1 were used to ream hole in a large amount. Rock drilling is an old but still tough challenge in geo-engineering, especially considering now we have to explore and/or mine fossil energy from more and more complicated stratum to satisfy the urgent and gluttonous needs of human beings. For quite a long time, the power of drilling machines had set the main drilling speed limitations. However, the situation has been changed with the development of drilling machines in recent decades [6]. Now, the failure of cemented carbide drill buttons (also called drill inserts) is the main life-limiting factors for the drilling tools in rock drilling, mineral cutting, gas and oil drilling, and tunneling [6–10]. And it has been even one of the factors restricting the efficiency of drilling machines. Aiming at improving the performances of drilling tools and drilling machines in geo-engineering, lots of investigations have been carried on the failure mechanisms of cemented carbide drill buttons, and the ways to improve the properties of cemented carbide [2–10]. Cemented carbide buttons undergo very complex working conditions during drilling, caused by the variation in properties between the types of rock or mineral, and in working patterns of various drilling machines [2–5,10–13]. Therefore, their failure mechanisms are much diversified. In general, the common failure mechanisms of drilling buttons [2–11] include microspalling, abrasion wear, WC grain pullout, extrusion of binder metal and reptile skin. But the primary failure mechanism varies with the different working conditions. In addition, the erosion of binder metal phase is also universal when the drills work in a corrosion or oxidation environment [3,5,9]. And through the observation on drill buttons selected from rotary/percussive drilling of a broad variety of rock types, Beste et al. [3,4] even indicated that during rock drilling, the rock may penetrate into cemented carbide drill buttons, which may have a significant influence on the button wear. On the other hand, the admirable properties of cemented carbide come from its special microstructure and constituent [14]. Generally, cemented carbides consist of two major phases, hard refractory phase

WC and soft metal binder phase Co/Ni. The two phases are consolidated through powder metallurgical techniques, which are deemed to be one of the oldest and most successful utilization of this technique [14–17]. The presence of WC grains imparts cemented carbide necessary hardness, strength and wear resistance, whereas binder phase Co/Ni contributes to the toughness and ductility of such hardmetal alloys [15,16]. In order to improve the properties of cemented carbide, much efforts have been directed towards developing nanostructured cemented carbide [15,17]. Such materials have been confirmed to possess considerably improved physical, mechanical and tribological properties at ambient as well as elevated temperatures, which make them more promising for geo-engineering drilling [15]. With the advance in synthesis methods of nanosized WC or WC-Co composite powders such as spray conversion process, and the advent of modern fast sintering techniques such as spark plasma sintering, dense nanostructured WC-Co bulk composites have been consolidated in several research laboratories [15,17–21]. However, even to date, the development of cemented carbide with WC grain size below 100 nm on a commercial scale has not been routinely achieved [15]. Functionally graded cemented carbides can provide a viable solution to the trade-off between wear resistance and toughness by varying cobalt content from the surfaces to the interior of a sintered piece [22–25]. The combination of high wear resistance and high toughness in a single component makes functionally graded cemented carbide much promising for the geo-engineering drillers [26,27]. Since most of the corrosion and oxidation foremost emerge in the binder phase of cemented carbide, finding a satisfactory alternative binder for WC-Co cemented carbide has also been a hot topic in this field [28–31]. The focus on cemented carbides with Ni or some carbide binders have been motivated by their improved erosion resistance [30,31]. However, the mechanical properties of such hardmetals are inferior to those of WC-Co cemented carbide. But as the initial powders applied became finer, the mechanical properties of such hardmetals can be greatly improved. And even pure WC can be prepared through electric field assisted fast sintering techniques [29]. They are very promising materials which can be applied in erosive environment in geoengineering. As for the contents of the present review, it will mainly focus on the recent advances in nanostructured, functionally graded, and alternatively-bonded other than Co or binderless cemented carbides, after a brief description of the application and failure mechanism of cemented carbide in geo-engineering. The synthesis methods of nanosized WC/WC-Co powders and some of the new sintering techniques are also included, because most of the newly-developed cemented carbides are derived from them. Along with a brief summary on the future research directions of cemented carbide, their potential applications in geo-engineering are also raised. 2. Rock drilling and failure mechanisms of cemented carbides in geo-engineering Cemented carbide has a long and successful history as wear-resistant parts in geo-engineering for their excellent combination of high hardness, high fracture toughness and relatively high wear resistance [3–5]. Usually, they serve as drill buttons in rock drilling, mineral cutting, gas and oil drilling, and even tunneling. The service environments of the buttons are very complex due to the different hardness of the drilling objectives and the various movement patterns of the drills [3–5,10]. Therefore, the failure mechanism of the cemented carbide buttons is quite different. 2.1. Rock drilling of cemented carbide in geo-engineering

Fig. 1. An example of cemented carbide drill bit (circled) to ream holes equipped with various types of buttons in China Continental Scientific Drilling Project.

Rock drill buttons are the main application of cemented carbide in geo-engineering, because of their excellent comprehensive

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wear-resistant properties, compared with their metal and other ceramic counterparts [3–5]. Commonly, a rock drill is a steel body in cylinder or other shapes equipped with cemented carbide (conventionally WC-Co hardmetal alloy) buttons, which is normally called as a drill bit or a drill crown. Figs. 1–4 show some examples of different types of drill bits with cemented carbide buttons. It can be seen that the cemented carbide buttons scatter on the mantles of the drill bodies. Generally, the buttons have two kinds of geometries, spherical buttons as seen in Fig. 2a, suitable for drilling hard rock such as quartzite, and ballistic ones as seen in Fig. 2b, suitable for drilling medium hard and moderately abrasive rock such as sand stone. According to the setting positions, the buttons are normally divided into peripheral and middle ones. During drilling, because of the higher velocity, the peripheral buttons always suffer from more abrasive wear, resulting in faster wearing out than the middle ones. Various rock types ask for different drilling methods and various application forms of cemented carbide buttons. Generally, rotarypercussive drilling method is suitable for drilling brittle rock and in this case, the drill bit as shown in Fig. 2 is always used. The drill buttons repeatedly impact the rocks from percussive movement, and the rotary movement lets the buttons impact on new positions each time. When the rock is of medium hard and plastic, the drill bit as shown in Fig. 3a is used to enhance the drilling speed. And the drill bit as shown in Fig. 3c is used for core drilling. When cemented carbide buttons are used to cut soft rock, such as coal, salt and concrete, they are often inlayed on steel roller or disc to form a roller bit, as shown in Fig. 4a–c. The different shapes of roller bits can be installed on tri-cone bit, reaming head for a raise drilling system and even the cutter head of a shield construction machine as seen in Fig. 4d–f, respectively. During drilling, the roller bit rotated and cut against the rock surface for mining coal and salt, reaming holes, tunneling and the like. For more complicated rock layer, composite drillers with different types of drill bits as shown in Fig. 1 may have to be applied.

2.2. The failure mechanisms of cemented carbide drill buttons The wear and deterioration of cemented carbide have been well investigated for many years, but the failure mechanism appearing on cemented carbide drill buttons is still very difficult to be classified [2–10], because they undergo various and complex working environments during drilling. The rock type, drilling temperature and causticity of the service environment can all result in different failure mechanisms [6,7]. What's more, it is hard to observe their service

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environment underground, and even harder to simulate their real service environment. Different drilling patterns can also lead to different failure mechanisms. For instance, when percussion is the main drilling pattern, the main failure mechanism is fatigue impact wear; however, when rotary/ crushing is the main drilling pattern, the predominant failure mechanism becomes abrasion and crack [8,9,12]. Fig. 5 gives an example to the complex situation of deterioration and wear of cemented carbide button after drilling quartzite rock for 280 m [4]. Fractured, plastically deformed and oxidized WC grains can be seen on the severely worn button surface in Fig. 5. The present top surface suffers from cracked, partly removed and partly loosened WC grains. Moreover, rock fragments gradually cover large parts of the surface, which may influence the wear resistance of cemented carbide. Beste et al. have tried to give an overview on the surface damage of cemented carbide rock-drill buttons [6]. A brief extract is given here. Simply, the damage can be interpreted as a combination of the continuous wear of WC grains and Co-binder phase, and the formation of cracks. Generally, the main types of the damages of cemented carbide buttons involved in microspalling, abrasion wear, WC grain pullout, extrusion of binder metal and reptile skin (or snake skin) [2,5,7,9–13]. The microspalling can result in the formation of small cracks within WC grains [12]. The abrasion wear mainly refers to the abrasion of the Co-binder phase and WC grains, which depends on the temperature and the sizes of the abrasive materials [11]. The extruded binder can accumulate at surface defects and gradually cause the fragmentation of WC grains [9]. The surface impacts, surface impact fatigue and thermal fatigue can also result in cracks [5]. Reptile skin, a common surface damage in rock drilling may lead to more extensive, tool-destroying cracks [5]. Naturally, a mix between these mechanisms is generally present. In a more recent work, Beste et al. [3,4,32] presented a significant new view on the action of cemented carbide buttons during rock drilling. They indicated that the surface layer transformed from the original WC-Co structure into a new composite. The integration of rock and cemented carbide took place on three forms: rock cover formation, intermixed layers, and deep rock channels, as shown in Fig. 6a–c, respectively. It can be seen that rock fragments are pounded into the button surface, becoming an integral part of the outer layer of the material. In the intermixed layers, the binder is partly replaced or intermixed by rock materials, and deep rock channels are the result of local penetration of rock material into the cemented carbide structure. The Co-binder phase gradually became embrittlement and degradation as the rock penetrated into the cemented carbide, resulting in WC grain pullout and the composite-scale crack formation [4].

Peripheral buttons

Middle buttons

(a)

Steel crown

(b)

Fig. 2. Typical rock drill crowns: (a) with spherical buttons suitable for drilling hard rock; (b) with sharp ballistic buttons suitable for drilling medium hard and moderately abrasive rock.

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Cemented carbide button

(a)

(b)

(c)

Fig. 3. Two drillers with drill body containing cemented carbide buttons: (a) a spiral bit, (b) a drill body inserted with cemented carbide button and (c) a core bit, in which panel c was taken from the product specification of China Bainas Drilling Equipment Co., Ltd.

More details about the binder phase degradation and rock intermixing mechanisms can be checked in Ref [32]. The oxidation and corrosion of WC grains were also observed in Ref [4]. According to the failure mechanisms of cemented carbide buttons, in order to extend the service life of cemented carbide drill buttons, the main methods concentrate on enhancing the wear resistance of

the cemented carbide without considerable deterioration of fracture toughness, in which the development of several kinds of new conceptual cemented carbide hardmetals such as nanostructured cemented carbide, functionally graded cemented carbide and alloys without metallic Co-binder phase may give the clues. Thus the following text will focus on the progress in these aspects.

(c)

(b)

(a)

Cemented carbides buttons

(d)

(f)

(e)

Roller bits Fig. 4. Examples of roller bits equipped with cemented carbide and their applications: (a) a cone roller bit, (b) a cylinder roller bit, (c) a disc roller bit, (d) a tri-cone bit equipped with cone roller bit as shown in panel a, (e) a reaming head for a raising drilling system equipped with cylinder roller bit as shown in panel b, and (f) the cutter head of a shield construction machine equipped with disc roller bit as shown in panel c, in which panels (d–f) were taken from the product specifications of China Bainas Drilling Equipment Co., Ltd., China Juli High-Tech Co., Ltd., and Guangzhou Herrenknecht Tunneling Machinery Co., Ltd., respectively.

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3.1. Synthesis of nanosized WC/WC-Co powders

Fig. 5. Typical appearance of a worn drill button surface [4]. This example shows a peripheral button from drilling a quartzite rock for 280 m. (A) Fractured WC grains, (B) oxidized WC grains, (C) plastically deformed WC grains, (D) adhering rock fragments (dark) and (E) a crack in the rock layer stuck to the surface.

3. Developing nanostructured cemented carbide hardmetals Nanostructured cemented carbide has been of interest to science and engineering communities in hardmetal field in the past few decades, owing to their potential superior performance, which stems from the refinement of WC grain size down to nanometric scale, resulting in significant improvements in hardness, strength and even tribological properties. Usually, the hardness of cemented carbide is in inverse proportion to its grain size, but the fracture toughness is inversely proportional to the hardness [15]. However, there is a considerable debate over the effect of grain size refinement on the fracture toughness of cemented carbide, because the large volume fractions of grain boundaries may impede the motion of cracks in nanostructured cemented carbide. In any case, nanostructured cemented carbide has demonstrated its special characteristics and advantages when compared to its microstructured counterparts, largely due to its ultrafine grain size and high surface to volume ratio. So the focus of the present review will be paid on the advance in processing nanostructured cemented carbide in combination with its probable application in geo-engineering, including the synthesis of nanosized WC and WC-Co raw powders, sintering of hardmetals with advanced fast sintering techniques, and materials properties. However, it must be pointed out that, even to date, the truly ‘nanostructured’ cemented carbide (WC grain sizeb 100 nm) has not been achieved on a commercial scale. So, in present review, the nanostructured cemented carbide refers to ultra-fine grained cemented carbide with WC grain sizes below about 250 nm, which is known as ‘nanostructured’ cemented carbide in the community [15].

The successful synthesis of nanosized WC/WC-Co powders is the basis of fabricating nanostructured WC-based alloys, due to the requirement on grain size of such hardmetals. And it has been also possible to develop dense cemented carbide without considerable deterioration of fracture toughness in the absence of metallic binder phase Co but with nanosized WC powders [15] because such powders can also bind cemented carbide well. Nanosized WC/WC-Co powders can be produced by numerous techniques. In the present review, the emphasis is paid on the commonest synthetic methods of nanosized WC/WC-Co powders: high energy ball milling (HEBM), spray conversion process (SCP) and chemical vapor synthesis (CVS), which are the most promising techniques in commercial use. HEBM is a straightforward method to produce nanosized WC or WC-Co powders with high energy mechanical ball milling [18,19]. It can reduce the sizes of cemented carbide powders into nanosize grade. In Ref [18], Zhang et al. indicated that high energy ball milling can efficiently refine the microstructure of WC in a WC/Co composite. In their work, WC-10Co composite powder with WC of mean Fisher sub-sieve size particle size of 5.6 μm was mixed with ethanol and then milled in a high energy ball mill with a selected rotation velocity of 250 rpm. After 10 h milling, the grain size of WC could be reduced to 11 nm, but a severe internal strain was introduced, and face centered cubic (fcc) cobalt was transformed into hexagonal close packed (hcp) one by a mechanical induced allotropic transformation. Mahmoodan et al. [33] have investigated the effect of ball to powder weight ratio (BPR) on the WC particle size in WC-10 wt.% Co powder after high energy ball milling. The results indicated that the milling with BPR of 5 can not produce enough energy to reduce the WC particles to nanosize scale and the milling with BPR of 25 or 35 can increase the abrasion of milling media and the lattice strain of the samples. The milling with BPR of 15 at 10 h in high energy ball mill was suitable to achieve nanosized powder. And a very interesting phenomenon was reported by Liu et al. in Ref [34] that the grain size of rare earth doped WC-Co powder became two times smaller as compared with that of the undoped powder within ball milling times of 25–45 h. Furthermore, HEBM with a long time can also induce chemical reaction. El-Eskandarany et al. [35] have successfully fabricated nanocrystalline powders of WC by high-energy ball milling of elemental W and C powders. The same result has been also obtained by Wang et al. through ball-milling the mixture of tungsten powder and activated carbon in vacuum [36]. However, most conventional mechanical milling techniques suffer from the disadvantages of long processing time, contamination, and high energy expenditure [35,36]. But by optimizing the variables of milling process, these disadvantages can be minimized. Butler et al. [19], for instance, reported that by optimizing milling parameters of

Fig. 6. Typical examples of targeting micro-sections through used drill buttons [3,4]. The selected views show the outer surface and the fracture cross-section. (a) Magnetite rock cover layer on the side of a button. A part of the rock layer has intentionally been removed [4]. (b) A button drilled in quartzite, revealing a dark intermixed layer [3]. (c) Deep rock channels are clearly visible as bright patches. The arrows indicate the maximum rock channel depth [3].

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a unique high-energy dual-drive planetary mill, the grain size of WC powder could be reduced from 800 nm to about 10 nm in approximately only 10 h milling time. In Refs [37,38], another chemical reaction process was reported that, tungsten oxide (WO3), cobalt oxide (Co3O4) and other additives could be milled as the raw materials first to form nanosized precursor mixture powders, and then the mixture powders were reduced by carbon black powders or by hydrogen and methane mixture gas. The resultant mixture powder of WC-Co had an average size of ~ 100 nm. SCP was developed and patented by McCandlish and Kear at Rutgers University, which brought a breakthrough on nanocrystalline WC-Co materials [39–41]. The process involves three steps: first, preparation and mixing of an appropriate starting solution; then, spray drying to form a chemically homogeneous precursor powder; and the last, fluid bed thermochemical conversion of the precursor into a desired nanophase composite powder [39–42]. For example, Cha et al. [43] has prepared nanosized WC-10Co powders by spray conversion process from a solution containing (NH4)6(H2W12O40) · 4H2O and Co(NO3)2, followed by oxidation, reduction and carbonization. Such process has been demonstrated at industrial level to commercially produce nanocrystalline cemented carbide powders [39,44]. However, according to the result obtained by Cha et al. [45], the transverse rupture strengths of nanocrystalline WC-10Co fabricated by SCP were lower than those of conventional WC-10Co, due to the less hcp/fcc phase transformation of Co phases during powder preparation. In a US patent [46], Kim et al. proposed a method to produce nanosized powders with homogeneously distributed grain growth inhibitors, by adding a water-soluble salt of V, Ta or Cr at the time of mixing water-soluble salts of W and Co during the initial production process of WC-Co cemented carbides. Recently, SCP process was further improved by changing the second step spraying drying into spraying freeze-drying in liquid nitrogen to reduce the agglomeration of the prepared powder during drying. Nanopowder in the size range of 20–30 nm was obtained by the modified SCP process [47,48]. CVS of WC powder processing, in general, involves the reaction of a tungsten salt precursor gas with hydrocarbon gas and reducing gas to prepare nanosized WC powders [49–53]. It is an attractive method because it can produce high purity nanopowders with good controls over size, shape, and crystal structure as well as easy control of reaction rate [52,54]. Normally, the final characteristics of CVS powders are affected by the preparation parameters such as carrier gas, reaction atmosphere, and reaction temperature [50–52,55–57]. CVS process has been used to synthesize nanosized WC powders because various precursors are available [51–60]. For instance, tungsten hexachloride (WCl6), tungsten hexafluoride (WF6) or ammonium paratungstate are generally favored as the tungsten salt precursor to prepare nanosized WC powder, because of their relatively low volatilization temperatures as well as the ease of their reduction by hydrogen or thermal decomposition [51,53,58–60]. Generally, hydrogen was used as reducing gas to produce tungsten metal powder, and hydrocarbons were usually used as carburizing agent because they are not expensive and stable up to a high temperature [60–63]. Methane is most commonly used among them because it is easy to control the amount of carbon reacting in the reaction system [62,64]. 3.2. The sintering of nanostructured cemented carbide It has been regarded as the main challenge to prevent the extensive coarsening of WC grains during sintering dense nanostructured cemented carbides. Lots of literatures [65–70] indicated that the grain growth of WC proceeds in two ways: solution/re-precipitation and re-arrangement/ coalescence. Typically, cemented carbides are consolidated via liquid phase sintering (LPS) in the temperature range of 1400–1500 °C, in which the metallic binder is in molten

state. Since the molten metallic binder (normally Co or Ni) can dissolve W and C to a certain extent, the solution/re-precipitation of these elements often results in significant growth of WC grains during sintering [65–67]. Moreover, according to the Gibbs–Thomson effect, the increase in WC/Co interfacial area of nanosized powders would result in a higher rate of dissolution of WC in Co, leading to solution/re-precipitation much more serious [15]. On the other hand, WC grains may re-arrange their orientation via rotation and slide. If their orientations get matched during this process, they can form a single coarser grain [15,57,69,70]. Due to the nanometer size and the surface energy anisotropy of the initial nanosized WC-Co powder, it can be consolidated in solid phase at a lower temperature, which is different from the case by using micron sized powder [71,72]. The results obtained by Fang et al. [72] indicated that a large fraction of densifications and WC grain coarsening might take place at solid-state temperatures during the sintering of nanosized WC-Co powders. Fig. 7 shows the changes of crystal shape and grain sizes of nanosized and micron-sized WC-Co powder compacts during heat-up [72]. It can be seen that during heat-up from 800 to 1200 °C, the WC grain growth and change from equi-axial to faceted platelet shape for nanosized powders were much more serious than those for micron-sized powders, demonstrating that nanosized WC-Co powders can be consolidated at a lower temperature but the grain growth is also serious at this stage. Wang et al. [73] have obtained similar results that a major fraction of grain growth occurs during the early stage of sintering nanosized WC-Co powders, before the LPS temperature is reached, because of the aggregation of nanosized WC grains. Therefore, in order to prepared nanostructured cemented carbides, it is a big challenge to retain the nanosized grains in dense bulk cemented carbide using conventional LPS route. In order to reduce the WC grain coarsening during sintering, many new sintering techniques have been investigated to consolidate nanostructured cemented carbides, including hot pressing sintering (HP) [74], hot isostatic pressing (HIP) [75,76], microwave sintering (MWS) [77,78], pulsed electric current sintering (PECS) [66,79], high frequency induction heating sintering (HFIHS) [80,81], pulse plasma sintering (PPS) [82], pressure assisted fast electric sintering (PAFES) [83] and spark plasma sintering (SPS) [20,48,76,84–89]. Table 1 presents a comprehensive survey of the microstructure and properties of cemented carbides, indicating that nearly full dense nanostructured cemented carbides at low sintering temperature and/or within a shorter holding time can be obtained by using these sintering techniques. Among these techniques, PECS, HFIHS, PPS and SPS are normally called electric field assisted sintering techniques, because they all involved the application of an external electric field along with uniaxial pressure. Such techniques are characteristic of rapid heating rate and shorter densification time even at lower temperature, which is very favorable to reduce the grain growth during sintering cemented carbides [90,91]. In particular, SPS may be the most well-known one among them, which is mainly characteristic of spark plasma created by a pulse direct current during heat treatment of powders in the graphite die [21,91]. A schematic of SPS apparatus is shown in Fig. 8 [92]. During sintering, the graphite die is heated directly by an electric current along with a uniaxial pressure, resulting in rapid heating and cooling of samples. Several remarkable features of the sintered ceramics via SPS process protrude it out over the commonly used sintering methods, including cleaner grain boundaries, noteworthy increase in superplasticity, improved bonding quality and so on [91]. More details can be checked in the review of Munir et al. [91]. More recently, Liu et al. [93] even proposed a novel rapid route to prepare WC-Co cemented carbide by in situ reactions and consolidation in SPS system with the start powders of WO3, Co3O4 and carbon black. Wei et al. [76] systematically compared the microstructures and properties of WC-Co cemented carbides prepared by SPS and HIP, respectively. The results indicate that the SPSed specimens have a finer grain structure, higher hardness but more massive

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(a)

(b)

(c)

(d)

67

Fig. 7. Changes of crystal shape and grain sizes of (a, b) nanocrystalline and (c, d) micron-sized WC-Co powder compacts during heat-up [72].

distribution of Co phase and lower fracture toughness than the specimen prepared by HIP sintering, because of the rapid heating rate and short sintering time of SPS technique. On the other hand, Shi et al. [94] consolidated WC-6.29Co nanocrystalline composite by combining SPS with HIP, and found that the combination of SPS with HIP would decrease the average grain size, thus improving the mechanical properties of cemented carbide. However, despite there have been a number of papers on electric field assisted sintering techniques in recent years, even to date there exist certain issues on them. The rapid heating rate and short holding time make them possible to sinter nanosized powders to near theoretical values with little grain growth [91]. However, because of its rapid heating rate and short holding time, temperature may not be uniform over the entire sample, which makes it difficult to confirm actual sintering temperature [79,95]. According to model calculations and experimental observations and measurements, Liu et al. [96] clarified that during SPS, the temperature of WC-Co particle was higher than the nominal temperature of the whole sample and increased sharply with extending the sintering time, due to the special heating mode of SPS and the heat transfer during SPS process. In Ref [79], Huang et al. obtained the similar result during PECS and concluded that the temperature gradient in the powder compact during PECS could be reduced by surrounding the graphite die with porous carbon felt insulation. In addition, the explanations to the superiority of the current activated sintering are also short of scientific adequacy. And since graphite dies are used in the sintering, inter-diffusion of carbon between the die and the solid powder compact may change the actual carbon content of the as-sintered cemented carbides, which is a critical factor for the phase and microstructure evolution and resulting in varying properties of such materials [85,95]. Another key method to prevent the extensive grain growth of WC is to apply suitable additives as grain growth inhibitors. Many transition metal carbides, such as VC, Cr3C2, Mo2C and TiC, are familiar grain growth inhibitors to the community, in which VC and Cr3C2 are regarded as the

most efficient ones due to their higher solubility and mobility in cobalt phase at low temperatures [86,97–103]. In the presence of metallic liquid phase, the inhibitors get dissolved and suppress the dissolution of W in the liquid phase, reducing the coarsening of WC grains by solution/ precipitation mechanism [97,98]. The segregation of these grain growth inhibitors, metal carbides, at WC/WC and WC/Co grain boundaries have been confirmed by transmission electron microscopy, high-resolution transmission electron microscopy, and even scanning electron microscopy (SEM) [99–101]. However, the inhibitory action by these additives has not been fully clarified yet. The grain growth inhibition mechanism was supposed to associate with the decrease of WC solubility in the binders, the increase of WC crystal edge energy, and the formation of a V or Cr rich thin interfacial layer on the surface of WC grains which acted as barrier to W diffusion [97,99,101]. But the use of some inhibitors may decrease the densification of the resultant cemented carbides, and the use of excessive amount of grain growth inhibitors may also result in the formation of other phases besides hard WC phase and Co-binder phase, thus decreasing in the mechanical properties of the resultant materials [104–106]. And the presence of solid solution phase (W,V)C and graphite phase can be seen in cemented carbide when the compositions were 8.1 wt.% V and 12 wt.% Co, as shown in Fig. 9 [104]. The existence of the graphite-phase and large sized (W,V)C cubic carbides may result in poor mechanical properties of the as-prepared cemented carbides [104]. And Hashe et al. [106] even used a pre-alloyed (W,V)C powder in the place of VC to reduce the driving force for the formation of (W,V)C during sintering. In addition, the grain growth of nanocrystalline WC-Co cemented carbides can also be controlled by doping rare earth or phosphorus [34,107,108]. Ou et al. found that with 0.5 or 1 wt.% La addition in WC-Co powder, the abnormal WC grains growth could be reduced and cobalt in the alloy would distribute more homogeneously around the WC grains [107]. The effect of phosphorus addition on the densification, grain growth and properties of nanocrystalline WC-Co composites by SPS was investigated by Zhang et al. in Ref [108]. The results indicate

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Table 1 The density, WC grain size and mechanical properties of WC-Co cemented carbides prepared using different sintering techniques. The properties of cemented carbide obtained by LPS are also mentioned for comparison. Compositions (wt.%)

Sintering technique

Sintering parameter

Relative density (%)

WC particle size in starting powder and in sintered composite (nm)

Vickers hardness (HV30)

Fracture toughness (KIC MPa· m1/2)

Refs

WC-12Co WC-Co

LPS HP

98.62 Nearly full

40–250 and 780 80 and 700

1381 –

17.3 –

[84] [74]

WC-10Co

HIP



150–250 and 500

1543

13.6

[76]

WC-12Co

HIP

1400 °C, 30 min 1300 °C, 25 MPa, 90 min 1420°C, 2 MPa, 60 min 1100 °C, 120 MPa, 30 min 1100 °C, 120 MPa, 30 min 1100 °C, 120 MPa, 30 min 1550 °C, 30 min

99.97

40–80 and 241–265

1800

10.7

[88]

99.85

40–80 and 204–224

1870

10.4

[88]

99.43

40–80 and 114–130

2070

9.9

[88]



38 and 650

15.1

[77]

1200 °C, 60 MPa, 1 min 1100 °C, 60 MPa, 5 min 1200 °C, 60 MPa, 5 min 1298 °C, 100 MPa, 5 min 1000 °C, 50 MPa, 10 min 1180 °C, 60 MPa, 10 min 1200 °C, 40 MPa, 5 min 1100 °C, 50 MPa, 10 min 1100 °C, 80 MPa, 5 min 1100 °C, 50 MPa, 10 min 1100 °C, 50 MPa, 10 min 1100 °C, 80 MPa, 5 min 1100 °C, 80 MPa, 5 min 1200 °C, 40 MPa, 5 min

99.5

b200 and 260

922 (periphery); 1466 (center) 1886

13

[81]

99

40–80 and 40–60

2250

15.3

[82]

99

40–80 and 85–135

2200

11.9

[82]

98.4

40–80 and ~1000

1748–1833

11.4–13.0

[83]

Nearly full

100 and 300

1580



[86]

99.72

~200 and 280

1836

12.87

[89]

99.5

~250 and 540

1828

12.17

[102]

99.89

40–250 and 800

1450

10.9

[84]

99.94

40–80 and 204–228

1840

10.6

[87]

97.66

40–250 and 550

1490

10.86

[84]

95.94

40–250 and 470

1570

11.42

[84]

98.95

40–80 and 144–164

2000

10

[87]

99.79

40–80 and 197–217

1880

10.3

[87]

98.7

~250 and 435

1922

14.14

[102]

WC-12Co-1.0Cr3C2 HIP WC-12 Co-1.0VC

HIP

WC-12Co

MWS

WC-10Co

HFIHS

WC-12Co

PPS

WC-12Co

PPS

WC-12Co

PAFES

WC-10Co

SPS

WC-10Co

SPS

WC-11Co

SPS

WC-12Co

SPS

WC-12Co

SPS

WC-12Co-0.5VC

SPS

WC-12Co-1.0VC

SPS

WC-12Co-1.0VC

SPS

WC-12Co-1.0Cr3C2 SPS WC-11Co-0.6 Cr3C2

SPS

that with small amount of phosphorus doped, the sintering activity of nanostructured WC-Co can be enhanced and even full density WC-7Co cemented carbide can be obtained at temperature as low as 1000 °C by SPS. 3.3. Mechanical properties of nanostructured cemented carbide In the various mechanical properties, the scientists and engineers in geo-engineering are more concerned with two properties of cemented carbides: hardness and fracture toughness. Generally, the hardness of cemented carbide at room temperature is in inverse proportion to its grain size. And it has been observed that the hardness can reach a significantly high value for nanostructured cemented carbides. In fact, a type of Hall–Petch relationship between WC grain size and hardness of WC-Co cemented carbides has been frequently and regularly observed [84,86,87]. On the other hand, a finer grain size results in a smaller mean free path at a given constant volume fraction of cobalt and a smaller plastic zone, which can also improve the hardness of the cemented carbides [15,109]. The results summarized in Table 1 also confirm that the hardness of cemented carbides increased with the decrease in WC grain size. However, there is a considerable debate over the effect of refinement in WC grain size on the fracture toughness of cemented carbides [82,84,86]. Usually, the fracture toughness increases with increasing WC grain size or the binder mean free path. For a conventional cemented

carbide hardmetal with microsized WC grains, a finer WC grain size can result in smaller mean free path, leading to a lower fracture toughness of cemented carbides [86]. However, Jia et al. [109] indicated that the increase in hardness of nanostructured composites does not imply the decrease in their bulk fracture toughness. For nanostructured cemented carbides, their flaw sizes, e.g. pore sizes, would be drastically reduced when compared to those of their conventional bulk hardmetal counterparts. And, there are large volume fractions of grain boundaries in the nanosized cemented carbides, which may impede the motion of cracks, thus resulting higher fracture toughness, because the deformation mechanisms of cemented carbides depend on grain boundary sliding and diffusion-controlled processes [57,110]. Because tribology is a system dependent property, the reported observations may have special meaning to the investigated tribological system with nanostructured cemented carbides. It has been observed that the wear resistance of cemented carbides increased with the increase in hardness, and in most cases, the wear process is primarily initiated by the removal and extrusion of the ductile binder phase, following plastic deformation and micro-abrasion [111,112]. Thus, cemented carbides with finer grain size and concomitantly finer binder mean free path, are expected to show improved wear performance due to the enhanced resistance to deformation of the binder phase and also the increased hardness of the finer grain sized composites. Therefore, nanostructured cemented carbides are very promising for the applications in geo-engineering.

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Fig. 8. Schematic of SPS apparatus [92].

4. Developing functionally graded cemented carbide The main failure mechanisms of a drill bit are of fragmentation and wear. Generally, it is accepted that the surface of a drill bit should have high hardness and thus good wear resistance to endure the abrasive wear, and the core of it have high strength to bear percussive, preventing cracks from propagation. Functionally graded cemented carbides provide a viable solution to the trade-off between their wear resistance and toughness by varying cobalt content from the surfaces to the interior of a sintered piece [22–25]. For instance, the near-surface area, with relatively low cobalt content, would have high wear resistance, while the interior, with relatively high cobalt content, would have high toughness. The functionally graded structure with a cobalt gradient in cemented carbide provides advantages in terms of the combinations of fracture toughness and

wear resistance in comparison with the conventional homogeneous WC-Co cemented carbide [113]. Such combinations in a single component would make functionally graded cemented carbides much promising in modern geo-engineering applications as tools for rock drilling, mineral cutting, gas and oil drilling, and concrete and asphalt milling [26,27,113–115]. It should be also noted that the difference in Co content in WC-Co composites may cause residual stresses within the material. And in most cases, when the cobalt content at the surface is lower than that in the interior, the residual stress on or near the surface is compressive, which is beneficial for enhancing the surface fatigue properties of components [23,116,117], and because the content of Co presents a gradual variation, the residual stress also changes gradually, which is also helpful in improving the performance of functionally graded cemented carbides. There are two existing industrial technologies producing functionally graded cemented tungsten carbides [118–121]. One has been used in producing cutting tools, preparing cemented carbide with a cobalt-enriched surface layer which is used to provide improved toughness to the cutting edge of a coating thereon and an improved surface on which one can deposit a coating such as TiN and TiC [122]. Another is the so-called dual-phase carbide which is normally characterized by three-zone structure, i.e. a cobalt-deficient zone at the surface to provide high wear resistance required to penetrate rocks, a cobalt-rich zone in the intermediate course to give the cemented carbide high toughness, and a carbon-deficient zone in the core to provide a balance of properties [113–115,120,123–125]. Fig. 10 shows a sketch diagram of the cross-section of a button insert [120]. The sintering techniques and the key factors affecting the formation of functionally graded cemented carbides are discussed below. 4.1. The sintering of functionally graded cemented carbide

Fig. 9. Typical backscattered SEM image showing the microstructure of WC-VC-Co with WC grains (light), binder phase (dark), and (W,V) C grains (gray). Some of the (W,V)C grains exhibit a core–rim structure [104].

Due to the potential advantages of functionally graded cemented carbide, there have been considerable reports over the past decades

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on the methods to manufacture it [126–129]. The manufacturing process of functionally graded cemented carbide can, usually, be divided in building spatially inhomogeneous structure (gradation) and transformation of this structure into a bulk material (consolidation). Normally, the gradation is built via die compaction of layers, continuous dry deposition of layers, sheet lamination and so on, which is predetermined during the deposition of the powders. More details about the processing techniques for functionally graded cemented carbide can be consulted in the review of Kieback et al. [130]. Typical functionally graded cemented carbide was manufactured via LPS [22,24,131]. However, during LPS, the liquid phase Co can easily migrate, and any initial gradient of cobalt content, that might have been built in a powder compact, might be eliminated [23,24,132,133]. Therefore, the main challenge is to develop a method that maintains or creates cobalt gradient during LPS. Liu et al. [125] have developed graded cemented carbides by presintering carbon-deficient compacts and subsequently carburizing. The formation mechanism of their prepared graded structure can be attributed to the Ostwald ripening induced by carbon diffusion. The rate difference of the Ostwald ripening in areas of different carbon contents leads to the variation in WC grain packing structure, hence driving the flow of the liquid Co from the surface to interior. In Ref [126], Lin et al. fabricated axisymmetric WC/Co functionally graded hardmetals via microwave sintering a WC-Co green compacts, which were molded by cold isostatic pressure with a discontinuously axisymmetrical laying distribution of Co (3-25 wt.%). Heat activation of WC grains and selective heating of microwave have enhanced the metallurgical and mechanical properties of functionally graded WC-Co hardmetals. A square-shaped large-size functionally graded WC-Co cemented carbide with 4 layers has been fabricated by Tokita et al. through a fully automated production-type SPS system [134]. The hardness of the top and bottom layers, containing 6 and 20 wt.% Co, were 20 and 9 GPa, respectively. Zhang et al. [129] prepared dense functionally graded cemented carbide via die compaction of WC-10Co and WC-20Co two layers using SPS. The analytical results of cross-section about the samples show that the WC-10Co layer with high hardness had a well combination with the WC-20Co layer with high toughness. Put et al. [128] indicated that continuous hardmetal deposits with a cobalt gradient can be prepared using electrophoretic deposition from a suspension of hardmetal powder in acetone. The deposits

can be sintered to nearly full density at the temperatures from 1290 to 1340 °C with a Co gradient ranging from 4 up to 17 wt.%. However, when the sintering temperature surpassed 1400 °C, the cobalt gradient of the deposits would disappear after sintering. 4.2. Key factors affecting the formation of graded cemented carbide Typical cemented carbide manufacturing processes relies on LPS for densification. However, during LPS, the liquid phase tends to distribute between WC grains to form homogeneous structure, resulting in the disappearance of any initial gradient of binder phase content in the powder compact. This phenomenon makes the processing of functionally graded WC-Co difficult. A solution to this problem is to use pressure assisted sintering or fast sintering techniques to consolidate the powders at solid state. However, high pressure processes or fast sintering techniques are usually costly, and there are obvious interface between the compact layers in the graded cemented carbide produced at solid state, which may downgrade the mechanical properties of functionally graded cemented carbide. Therefore, LPS, up to now, is the most viable and economical method for producing functionally graded WC-Co if the cobalt gradient can be maintained or created during LPS. Fang et al. [22,24,25,131–133] systematically investigated the key factors that affect the migration of liquid phase in WC-Co alloys: initial Co gradient, WC grain size, and carbon content. These studies discussed the impacts of each factors on the migration of liquid phase and explained the formation of cobalt gradient during LPS. Fang et al. [25,131] indicated that when the initial WC grain sizes were identical and C contents were stoichiometric in each layer in initial bi-layer compacts, the initial differences in Co contents affected a little to the Co gradient sintering at 1400 °C for 1 h, at which the Co-binder was liquid phase. At the LPS temperature, the layer with higher cobalt content will have wider liquid channels in comparison with the layer with lower cobalt content. Due to the capillary force, liquid will flow from higher cobalt content layer into lower cobalt content layer. Therefore, the cobalt content has completely homogenized across the two layers after sintering at 1400 °C for 1 h. A difference in WC grain size will produce a difference liquid migration pressure, resulting in the migration of liquid phase [22,25,131,133–135]. From Fig. 11a, it can be seen that finer grains form smaller liquid channels while coarser grains form larger liquid

Fig. 10. Sketch diagram of the cross-section of typical WC-Co dual phase cemented carbide button [120].

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Fig. 11. (a) Typical SEM micrograph of WC-10Co bi-layer and (b) distribution of cobalt in a WC-10Co bi-layer specimen, both with difference in grain size sintered at 1400 °C for 1 h [131].

channels. Thus, liquid cobalt will flow into the layer with finer WC grain size by migration pressure during sintering. Fig. 11b shows the cobalt distribution profile of the bi-layer WC-Co specimen as shown in Fig. 11a measured using EDS [131]. A step wise profile of cobalt concentration can be observed between the two layers after sintering due to the difference of particle sizes. The liquid phase migration was also investigated by heat treating the assembles, which were fabricated by stacking on the top of each other of two cylinders of WC-Co hardmetals, which were different in mean WC particle size (relative size ratio of 3.5) and cobalt concentration (6 or 15 wt.%), at 1400 °C under protective atmosphere for a dwell time of 7 h by Colin et al. [136]. The results also indicated that the Co liquid always tends to migrate from the hardmetals with larger grain size and higher content cobalt concentration to the ones with smaller grain size and lower content cobalt concentration. The carbon content affects the volume fraction of liquid phase a lot [22,24,25,127,131,135,137]. Fan et al. [137] found that during liquid-phase sintering of cemented carbide, the liquid phase migrated from the region with higher C content toward that with lower C content when the initial WC grain sizes and liquid phase content were identical. Due to the phase reaction between carbon and CoxWxC (η) phase, the volume fraction of Co-binder phase can vary depending on the total carbon content in the alloy. Significant deviations below or above the stoichiometric carbon content will result in the occurrence of three-phase equilibrium structure involving WC + Co + η or WC + Co + C (graphite), respectively. Normally, η phase is a complex phase formed in WC-Co alloy during sintering due to a significant carbon deficiency. Based on carbon diffusion and phase reaction processes, carbon was identified as a key factor

affecting the cobalt gradient within a sintered WC-Co sample. Fig. 12 shows the EDS analysis of the cobalt distribution of WC-10Co bi-layer with identical initial cobalt contents and grain sizes but different initial total carbon content in the layers sintered at 1400 °C [131]. It can be seen that the cobalt content varied gradually from 8 wt.% in the layer initially with excess total carbon content to about 13 wt.% in the layer initially with a deficiency in total carbon content. It can also demonstrate that as the sintering time is increased, carbon diffuses further into carbon deficient layer and reacts with η phase to produce WC-Co resulting in the migration of cobalt in the same direction. When the initial Co and carbon contents were both different, the bi-layer WC-Co sample shows a continuous gradient in cobalt after sintering at 1400 °C, which can be illustrated in Fig. 13 [25]. The results indicated that cobalt migrated in the direction of carbon diffusion during LPS. However, the existence of graphite phase or η phase would be detrimental to the mechanical properties of WC-Co materials. Therefore, suitable carbon content gradient is very momentous in producing cobalt gradient. In Refs [23,127,138], cemented carbide with a cobalt content gradient was fabricated by the heat treatment of fully sintered WC-Co materials in carburizing atmospheres. The results demonstrated that the heating temperature, carbon potential of the atmosphere and heating time had significant effects on the formation of Co gradients. The Co gradient formed only when the proposed process was conducted at the temperature within the WC(s)-Co(l)-Co(s) coexisting range (1275-1325 °C), the temperature range of the shaded area in Fig. 14 [23]. A higher carbon potential of the atmosphere would generate higher surface carbon composition, resulting in a larger driving force for carbon diffusing. Consequently, Co will

Fig. 12. Cobalt distribution of WC-10Co bi-layer with identical initial cobalt contents and grain sizes but different initial total carbon contents in the layers sintered at 1400 °C [131].

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Fig. 13. Comparative plot of the cobalt distribution of WC-Co with identical initial grain sizes but different initial carbon and cobalt contents (solid line) and WC-Co with identical initial grain sizes and stoichiometric carbon contents but different initial cobalt contents (dotted line) sintered at 1400 °C for 1 h [25].

migrate further into the specimen, resulting in greater thickness and amplitude of cobalt gradient. The increase of the thickness of Co gradient obeys the parabolic law of diffusion with time; whereas when the time is extended to 360 min, the linear relationship is no longer followed.

5. Developing Co-free cemented carbide Since the working conditions of the cemented carbide tools in geo-engineering are complex during mineral cutting, gas and oil drilling, and so on, enhancing the corrosion and oxidation resistance also becomes a hot research direction of cemented carbides. As mentioned in Section 2, the main corrosion and oxidation of cemented carbides foremost emerge in the binder phase. Therefore, considerable effort to find a satisfactory alternative binder has been prompted. The focus on Ni or carbide such as TiC and Mo2C as binders has been motivated by their results showing a higher corrosion and oxidation resistance [28–31]. However, the mechanical properties (hardness and strength) of WC-Ni or WC-TiC hardmetals are relatively inferior to those of WC-Co counterparts. With the recent development of synthesizing nanosized powders and the electric field assisted fast sintering techniques such as HFIHS and SPS, it has been possible to develop ultrafine WC-Ni cemented carbide with high hardness and high toughness. And the sintering of Co-free WC-TiC, WC-Mo2C, WC-Al2O3, and even pure WC has been also possible. At the same time, through adding some materials with high hardness/strength

Fig. 14. Vertical section of the ternary phase diagram of W-Co-C with constant 10 wt.% Co [23].

into the alloy matrix, the hardness/strength of WC-Ni cemented carbide can be directly increased. Table 2 presents a comprehensive survey of the microstructure and properties of Co-free cemented carbides, indicating that nearly full dense microstructure, high hardness and relatively acceptable fracture toughness can be obtained. Moreover, Engqvist et al. [139] have studied the abrasive wear of two Co-free carbides (WC-Mo2C and WC-TiC-TaC) and then indicated that compared with alumina ceramics and WC-6Co cemented carbide, the WC-Mo2C hardmetal showed the lowest abrasion rate and WC-TiC-TaC an intermediate. The wear mechanism was surprisingly ductile for WC-Mo2C, whereas WC-TiC-TaC suffered from grain pullout of TiC.

5.1. WC-Ni cemented carbide To date, there are several metals which have been proved possible substitutes for cobalt as binder metal phase in WC-based cemented carbides [140,141]. Nickel is an exciting and promising candidate among all those binder metals because of its good wet ability to WC and much better performance of WC-Ni cemented carbides in erosion resistance to chemical attack than that of WC-Co cemented carbides [140–142]. Moreover, nickel appears to be more ductile than cobalt, so that it can reduce loss of WC grains by extrusion of nickel outwards to replace the eroded binder between WC grains, which can prevent early exposure of the WC grains with consequent loss by tearing out from the binder. And the last, nickel has a relatively lower price than that of cobalt, which is potential to decrease the cost of the drilling tools [140]. However, in practice, there are some challenges in replacing cobalt with nickel. Nickel, unlike cobalt, has no fcc to hcp transformation. Normally, the fcc structure can be expected to be more ductile and with a lower work-hardening coefficient, to absorb less energy in the powder milling process [140]. Therefore, nickel would absorb lower energy in the powder milling process than cobalt, which may make WC-Ni cemented carbide more difficult to achieve complete densification than WC-Co cemented carbide. Another challenge that nickel has not been more widely substituted for cobalt in industry is that the hardness and strength of WC-Ni hardmetals are lower than those of WC-Co alloys of similar composition. Although so, the recent advances in synthesizing nanosized powders and electric field assisted fast sintering techniques, as mentioned in Sections 3.1 and 3.2, make it possible to consolidate ultrafine WC-Ni cemented carbide with almost full density. Both the hardness and strength can be greatly improved. Kim et al. [141] have consolidated cemented carbides of WC-xNi (x = 8–12 wt.%) via HFIHS. The density was about 98% with WC grain size of 300 nm. The hardness

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Table 2 The density, WC grain size (for different initial WC particle size) and mechanical properties of cemented carbide with other binder phase other than Co, carbide/oxide bonded cemented carbide and pure WC materials. Compositions (wt.%)

Initial WC particle size (μm)

Sintering technique

Sintering temperature and pressure

Sintered density (%)

WC grain size in sintered composite (nm)

Vickers hardness (HV)

Fracture toughness (KIC, MPa · m1/2)

Fracture strength (MPa)

Refs

WC-10Ni-0.7VC-0.3 TaC WC-10Ni-0.7VC-0.3 TaC-15.5 cBN WC-10Ni WC-20at% TiC WC-6Mo2C WC-2VC WC-16VC WC-6.8Al2O3 WC-5ZrO2 WC-6ZrO2 Pure WC Pure WC Pure WC Pure WC Pure WC Pure WC-0.1Carbon

0–1

SPS

1350 °C, 50 MPa

97.1–97.5

330

1840



2023–2259

[30]

0–1

SPS

1350 °C, 50 MPa

98

200–400

3200



1250

[143]

0.4a 0.4a 0.4a 0.2 0.2 0.1 – 0.2 0.57 0.04–0.07 0.2 0.4 0.4 0.2

HFIHS HFIHS HFIHS PECS PECS SPS PECS SPS SPS SPS SPS SPS HFIHS SPS

1400 1900 1700 1900 1800 1770 1700 1300 1700 1750 1500 1600 1500 1460

97.5 98.5 99 96 100 99.2 99.6 97.1 98 100 99.6 97.6 98.5 –

300 200 450 260 – – – 200–300 600 305 570 360 380 350

1810 2300 2460 2650 2130 2450 2400 2100 – 2730 2650 2480 2850 (HRA 94)

13.5 6.3 4.8 4.2 3.8 6 5.6 10 – – 9 6.6 7.1 8.376

– – – – – – – 1300 – – – – – –

[141] [31] [145] [146] [146] [151] [150] [148] [29] [92] [152] [153] [154] [155]

a

°C,60 MPa °C, 60 MPa °C, 60 MPa °C, 60 MPa °C, 60 MPa °C, 70 MPa °C, 60 MPa °C, 50 MPa °C, 50 MPa °C, 126 MPa °C, 60 MPa °C, 60 MPa °C, 60 MPa °C, 30 MPa

As measured by Fisher sub-sieve sizer.

of WC-10Ni in this work can reach 1810 HV and the toughness was 13.5 MPa · m 1/2. Wittmann et al. [142] investigated the WC grain growth and effects of grain growth inhibition in nickel binder cemented carbide. The results indicated that WC grains grew more significantly in high carbon content WC-Ni cemented carbide, compared with that in low carbon content WC-Ni cemented carbide. In addition, VC has been proved to be by far the most effective grain growth inhibitor in WC-Ni cemented carbide, followed by TaC, Cr3C2, TiC and ZrC, as shown in Fig. 15 [142]. For WC-10Ni alloys, the suitable addition of VC is about 1 wt.% and the hardness can reach about 1840 HV. With VC and TaC as grain growth inhibitors, ultrafine WC-Ni cemented carbides with different fractions (6–10 wt.%) of binder metal nickel were fabricated by SPS in Ref [30]. The average WC grain size of the prepared WC-Ni cemented carbides was about 330 nm, and there were lots of micro-pores in the samples. The relative density of the samples was all higher than 92%. The obtained ultrafine WC-Ni cemented carbides showed a very short binder mean free path, resulting in excellent performance in mechanical strength. Afterwards, different amounts of cubic boron nitride (cBN) were added into WC-Ni-VC-TaC cemented carbides [143]. The relative density of the cemented carbides increased with the addition fraction of cBN, which could almost reach 98% as shown in Fig. 16, when the addition fraction of cBN increased from 0 up to 50 vol.%, and meanwhile, the hardness of the samples increased from 2100 to 3200 HV, but the flexural strength decreased from 1950 to 1250 MPa. It was worth to note that the cBN particles were coated with nickel thin

Fig. 15. Two-phase WC-10Ni alloys with additions of VC, Cr3C2 and TiC, sintered at 1480 °C for 1 h [142].

film of about 500 nm thickness using a physical vapor deposition method before use and there was strong bonding force between the cBN particles and cemented carbide matrix as shown in Fig. 17. Correa et al. [144] investigated the influence of silicon as an additive on the microstructure and mechanical properties of WC-Ni cemented carbide processed by conventional powder metallurgy. The results revealed that silicon can strengthen the WC-Ni cemented carbide via solid solution in nickel. The WC-Ni-Si with 10 wt.% of binder showed bulk hardness similar to the conventional WC-Co cemented carbides and superior flexure strength and fracture toughness. 5.2. Carbide/oxide-bonded cemented carbide Carbide-bonded cemented carbides are WC-based hardmetals consolidated by carbides of TiC, Mo2C, VC and so on, which act as binder phase without any metallic binder addition. Generally, the binder phase, such as Co, Fe and Ni, not only facilitate densification during sintering but also lead to an improved fracture toughness and strength. However, these binder phases are inferior to the carbide phase in chemical characteristics, in which corrosion and oxidation always start with the binder phase [12,15,17]. In addition, the rocks are easily penetrated into cemented carbide drill buttons during rock drilling with the existence of binder metals due to its relatively

Fig. 16. Apparent density and relative density of WC-Ni-VC-TaC-cBN cemented carbide samples as a function of the addition fraction of cBN [143].

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c-BN

WC

Fig. 17. Typical secondary electron SEM images on the fracture surface of WC-Ni cemented carbide samples, showing strong bonding force between the cBN particles and cemented carbide matrix [143].

low hardness. Therefore, carbide-bonded cemented carbides were proposed [28,29] for their potential relatively high hardness. TiC is a typical carbide binder for WC because they can form WC-TiC solid solution phase at lower sintering temperature [28,31]. WC-TiC cemented carbide exhibits excellent corrosion and oxidation resistance, which has been applied in mechanical seals and sliding parts. However, both the hardness and fracture toughness of WC-TiC decreased when compared with those of WC-Co cemented carbide. Due to C segregation near the WC/TiC interface, surface fractures were easily occurred during friction or under abrasive conditions, resulting in lower wear resistance than that of WC-Co cemented carbides. Consolidation of cemented carbide with nanosized WC grains can result in significant improvements in mechanical properties. However, conventional sintering methods often result in grain growth during bulk densification. The recent development of fast sintering techniques can effectively prevent grain growth, resulting in considerable improvement of the mechanical properties of carbide-bonded cemented carbide. In Ref [31], Kim et al. consolidated WC-TiC cemented carbide using HFIHS with ultrafine powders of WC and TiC. The relative density of the composite was about 98.5% for an applied pressure of 60 MPa and an induced current for 90% output of total capacity. The average WC grain size was about 200 nm. At the sintering temperatures of carbides, around 1450 °C, WC-TiC may form the mixed carbide (W,Ti)C plus free carbon, as shown in Fig. 18a [31]. They also used this fast sintering technique to fabricate WC-Mo2C cemented carbide [145]. An almost complete densification (about 99%) of the materials was achieved within

1 min for an applied pressure of 60 MPa and an induced current for 90% output of total capacity. The hardness decreased and the fracture toughness increased with increasing Mo2C content. VC, which shows the best inhibitor effect in WC-Co, can also act as a binder phase as what TiC or Mo2C act [146]. The sintering temperature for full densification can be lowered with increasing VC addition to 1600 °C for a 16 wt.% VC grade. The WC dissolved into VC to form mixed carbide (V,W)C, as shown in Fig. 18b, a bright WC grains and a darker contrast (V,W)C phase [146]. The consolidation of WC with the additions of VC and/or Cr3C2 using SPS was also investigated in Ref [147]. The results indicated that both VC and/or Cr3C2 can reduce the grain growth of WC. With the same added content, the addition of VC resulted in smaller WC grains but lower hardness and fracture toughness. In recent studies, metallic oxides such as Al2O3 and ZrO2 were also added as a binder phase in Co-free cemented carbide [148–150]. Malek et al. fabricated WC-ZrO2 via SPS within a few minutes at 1700 °C [150]. According to the results, they concluded that the addition of small amounts of ZrO2 can lead to an acceptable fracture toughness and significantly higher flexural strength without sacrificing hardness. Basu B [149] obtained WC-6 wt.% ZrO2 nanocomposites with full density using SPS at relatively lower sintering temperature of 1300 °C. More recently, Mukhopadhyay et al. [148] proposed that it is possible to develop dense WC-based ceramic nanocomposites, possessing mechanical properties comparable or even superior to those of WC-6 wt.% Co cemented carbides, by adding ZrO2 and using SPS processing. WC-Al2O3 materials without metallic binder addition were also consolidated via SPS at a temperature range of 1800– 1900 °C by Zheng et al. [151]. The results indicated that a proper content of Al2O3 additive helps to limit the formation of W2C phase in sintered WC materials. 5.3. Pure WC hardmetals In order to produce hard and erosion resistance tool materials, pure WC was also favorable. However, due to the high melting point of WC, it is very difficult to sinter pure WC via conventional LPS. Thanks to the development of nanosized WC synthesis process and the advent of electric field assisted fast sintering techniques, the preparation of pure WC cemented carbides has been possible. Omori et al. [21] indicated that pure WC powders can be consolidated by SPS above 1900 °C, resulting in samples with Vickers hardness about 24 GPa. However, the transverse rupture strength of the resultant samples, only 1 GPa, was much lower than that of WC-Co cemented carbide due to the large size of WC grains. Therefore, fine WC grain size powder, which may result in higher hardness and strength, is also necessary to prepare pure WC hardmetals. To date, electric field assisted fast sintering techniques such as SPS and HFIHS are favorable methods to produce dense pure WC materials

C WC

(Ti, W)C WC (V, W)C

(a)

(b)

Fig. 18. Typical SEM micrographs of (a) WC-40at% TiC composites sintered by HFIHS [31]; (b) the polished WC-12VC grades sintered by PECS for 1.5 min at 1800 °C [146].

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Fig. 19. (a) Dimensional changes of sintered WC at 1700 °C without holding time from various initial powder sizes; (b) variation of density in sintered WC from 4.06 μm powders at a temperature range from 1550 to 1800 °C [29].

at relative lower sintering temperatures. During the sintering, the key factors that influenced the densification and mechanical properties of pure WC materials were the sintering temperature, initial WC powder size and holding time [29,92,152–154]. However, it must be noted that these features are in turn related to each other. For instance, Cha et al. [29] indicated that when the initial WC powder size becomes finer, the shrinkage of WC powder compact starts at a lower temperature and the materials can be sintered with full densification at lower sintering temperature as shown in Fig. 19a, and the holding time decreased with increasing sintering temperature. For example, when the sintering temperature is over 1700 °C, the full density of WC can be obtained even without holding as seen in Fig. 19b. Moreover, during SPS sintering, the decarburization of WC would always happen. For instance, Zhang et al. [155] found that the carbon-lacking phases of WC1-x and W2C occupied a large volume fraction of the phase configuration in the samples prepared by direct SPS sintering 0.2 μm sized WC powder, which decreased the fracture toughness of the materials. At the temperature of 1460 °C and with 0.1 wt.% carbon addition, highly dense WC bulks with fracture toughness of 8.376 MN· m −3/2 were obtained. 6. Conclusions and outlook The research results over the past decades have unambiguously established the potential application in geo-engineering of bulk nanostructured cemented carbide, functionally graded cemented carbide, and Co-free cemented carbide. Nanostructured cemented carbide can provide high hardness and high wear resistance which may improve the service life of the drill bits. The drill buttons industry has already seen the benefits of functionally graded cemented carbide for its combination of hardness and high toughness. WC-Ni cemented carbide, cemented carbide with carbide binder or even pure WC materials have been fabricated via electric field assisted fast sintering techniques, especially SPS, performing high erosion resistance without considerable deterioration of fracture toughness, which are also of great potential to be used as wear part in gas or oil drilling at erosive environment. Although such materials can be fabricated via electric field assisted fast sintering techniques, they are not yet used commercially. One of the reasons is the expenditures involved and limitations on the shapes/sizes of the components, which can be sintered via such sintering techniques. Considering that the fracture toughness is an important property considered in the geo-engineering application, further investigation is necessary to understand the fracture behavior of nanostructured cemented carbide. Further improvement of the fracture toughness of the cemented carbide without metallic metal binder phase is also necessary. The design of functionally graded cemented carbide with optimized gradient is also of interest in further.

Acknowledgments The work was supported by Grand Survey on Land and Nature Sources of China sponsored by China Geological Survey (grant no. 1212010916026), Ph.D. Programs Foundation by Ministry of Education of China (grant no. 20100022110002), and National Key Technology R & D Program of China (grant no. 2011BAB03B08). References [1] Spriggs GE. A history of fine grained hardmetal. Int J Refract Met Hard Mater 1995;13:241–55. [2] Montgomery RS. The mechanism of percussive wear of tungsten carbide composites. Wear 1968;12:309–29. [3] Beste U, Jacobson S, Hogmark S. Rock penetration into cemented carbide drill buttons during rock drilling. Wear 2008;264:1142–51. [4] Beste U, Jacobson S. A new view of the deterioration and wear of WC/Co cemented carbide rock drill buttons. Wear 2008;264:1129–41. [5] Larsen-Basse J. Wear of hard-metals in rock drilling: a survey of the literature. Powder Metall 1973;16:1–32. [6] Beste U, Hartzell T, Engqvist H, Axén N. Surface damage on cemented carbide rock-drill buttons. Wear 2001;249:324–9. [7] Beste U, Jacobson S. Micro scale hardness distribution of rock types related to rock drill wear. Wear 2003;254:1147–54. [8] Duan LC, Liu XY, Mao BS, Yang KH, Tang FL. Research on diamond-enhanced tungsten carbide composite button bits. J Mater Process Technol 2002;129: 395–8. [9] Larsen-Basse J. Binder extrusion in sliding wear of WC-Co alloys. Wear 1985;105: 247–56. [10] Stjernberg KG, Fisher U, Hugoson NI. Wear mechanisms due to different rock drilling conditions. Powder Metall 1975;18:89–106. [11] Larsen-Basse J, Perrott CM, Robinson PM. Abrasive wear of tungsten carbide-cobalt composites. I. Rotary drilling tests. Mater Sci Eng 1974;13:83–91. [12] Swick KJ, Stachowiak GW, Batchelor AW. Mechanism of wear of rotary-percussive drilling bits and the effect of rock type on wear. Tribol Int 1992;25:83–8. [13] Larsen-Basse J. Effect of composition, microstructure, and service conditions on the wear of cemented carbides. J Met 1983;35:35–42. [14] Exner HE. Physical and chemical nature of cemented carbides. Int Mater Rev 1979;24:149–73. [15] Mukhopadhyay A, Basu B. Recent developments on WC-based bulk composites. J Mater Sci 2011;46:571–89. [16] Upadhyaya GS. Materials science of cemented carbides – an overview. Mater Des 2001;22:483–9. [17] Prakash LJ. Application of fine grained tungsten carbide based cemented carbides. Int J Refract Met Hard Mater 1995;13:257–64. [18] Zhang FL, Wang CY, Zhu M. Nanostructured WC/Co composite powder prepared by high energy ball milling. Scr Mater 2003;49:1123–8. [19] Butler BG, Lu J, Fang ZZ, Rajamani RK. Production of nanometric tungsten carbide powders by planetary milling. Int J Powder Metall 2007;43:35–43. [20] Jia CC, Tang H, Mei XZ, Yin FZ, Qu XH. Spark plasma sintering on nanometer scale WC-Co powder. Mater Lett 2005;59:2566–9. [21] Omori M. Sintering, consolidation, reaction and crystal growth by the spark plasma system (SPS). Mater Sci Eng A 2000;287:183–8. [22] Eso OO, Fan P, Fang ZZ. A kinetic model for cobalt gradient formation during liquid phase sintering of functionally graded WC-Co. Int J Refract Met Hard Mater 2008;26:91–7. [23] Guo J, Fang ZZ, Fan P, Wang X. Kinetics of the formation of metal binder gradient in WC-Co by carbon diffusion induced liquid migration. Acta Mater 2011;59: 4719–31.

76

X. Ren et al. / Int. Journal of Refractory Metals and Hard Materials 39 (2013) 61–77

[24] Eso OO, Fang ZZ, Griffo A. Kinetics of cobalt gradient formation during the liquid phase sintering of functionally graded WC-Co. Int J Refract Met Hard Mater 2007;25:286–92. [25] Fang ZZ, Eso OO. Liquid phase sintering of functionally graded WC-Co composites. Scr Mater 2005;52:785–91. [26] Fischer UKR, Hartzell ET, Akerman JGH. Cemented carbide body used preferably for rock drilling and mineral cutting. US Patent No. 4,743,515, 1988. [27] Fischer UKR, Waldenstrom M, Hartzell ET. Cemented carbide body with increased wear resistance. US Patent No. 5,856,626, 1999. [28] Imasato S, Tokumoto K, Kitada T, Sakaguchi S. Properties of ultra-fine grain binderless cemented carbide ‘RCCFN’. Int J Refract Met Hard Mater 1995;13: 305–12. [29] Cha SI, Hong SH. Microstructures of binderless tungsten carbides sintered by spark plasma sintering by spark plasma sintering process. Mater Sci Eng A 2003;356: 381–9. [30] Rong HY, Peng ZJ, Ren XY, Peng Y, Wang CB, Fu ZQ, et al. Ultrafine WC-Ni cemented carbides fabricated by spark plasma sintering. Mater Sci Eng A 2012;532:543–7. [31] Kim HC, Kim DK, Woo KD, Ko IY, Shon IJ. Consolidation of binderless WC-TiC by high frequency induction heating sintering. Int J Refract Met Hard Mater 2008;26:48–54. [32] Beste U, Coronel E, Jacobson S. Wear induced material modifications of cemented carbide rock drill buttons. Int J Refract Met Hard Mater 2006;24:168–76. [33] Mahmoodan M, Aliakbarzadeh H, Gholamipour R. Microstructural and mechanical characterization of high energy ball milled and sintered WC-10 wt% Co-χTaC nano powders. Int J Refract Met Hard Mater 2009;27:801–5. [34] Liu S, Huang ZL, Liu G, Yang GB. Preparing nano-crystalline rare earth doped WC/Co powder by high energy ball milling. Int J Refract Met Hard Mater 2006;24:461–4. [35] El-Eskandarany MS, Mahday AA, Ahmed HA, Amer AH. Synthesis and characterizations of ball-milled nanocrystalline WC and nanocomposite WC-Co powders and subsequent consolidations. J Alloys Compd 2000;312:315–25. [36] Wang GM, Campbell SJ, Calka A, Kaczmarek A. Synthesis and structural evolution of tungsten carbide prepared by ball milling. J Mater Sci 1997;32:1461–7. [37] Lee GH, Kang S. Sintering of nano-sized WC-Co powders produced by a gas reduction–carburization process. J Alloy Compd 2006;419:281–9. [38] Wei CB, Song XY, Zhao SX, Zhang L, Liu WB. In-situ synthesis of WC-Co composite powder and densification by sinter-HIP. Int J Refract Met Hard Mater 2010;28:567–71. [39] McCandlish LE, Kear BH, Bhatia SJ. Spray conversion process for the production of nanophase composite powders. US Patent No. 5,352,269, 1994. [40] McCandlish LE, Kear BH, Kim BK. Chemical processing of nanophase WC-Co composite powders. Mater Sci Technol 1990;6:953–7. [41] McCandlish LE, Kear BH, Kim BK. Processing and properties of nanostrutured WC-Co. Nanostruct Mater 1992;1:119–24. [42] Zhu YT, Manthiram A. A new route for the synthesis of tungsten carbide-cobalt nanocomposites. J Am Ceram Soc 1994;77:2777–8. [43] Cha SI, Hong SH, Ha GH, Kim BK. Mechanical properties of WC-10Co cemented carbides sintered from nanocrystalline spray conversion processed powders. Int J Refract Met Hard Mater 2001;19:397–403. [44] Fang Z, Eason JW. Study of nanostructured WC-Co composites. Int J Refract Met Hard Mater 1995;13:297–303. [45] Cha SI, Hong SH, Ha GH. Microstructure and mechanical properties of nanocrystalline WC-10Co cemented carbides. Scr Mater 2001;44:1535–9. [46] Kim BK, Ha GH, Woo Y. Method of production WC/Co cemented carbide using grain growth inhibitor. US Patent No. 6,511,551, 2003. [47] Bonache V, Salvador MD, Busquets D, Burguete P, Martínez E, Sapiña F, et al. Synthesis and processing of nanocrystalline tungsten carbide: towards cemented carbides with optimal mechanical properties. Int J Refract Met Hard Mater 2011;29:78–84. [48] Xi XL, Pi X, Nie ZR, Song SL, Xu XY, Zuo TY. Synthesis and characterization of ultrafine WC-Co by freeze-drying and spark plasma sintering. Int J Refract Met Hard Mater 2009;27:101–4. [49] Kear BH, Strutt PR. Chemical processing and applications for nanostructured materials. Nanostruct Mater 1995;6:227–36. [50] Hojo J, Oku T, Kato A. Tungsten carbide powders produced by the vapor phase reaction of the WCl6-CH4-H2 system. J Less Common Met 1978;59:85–95. [51] Kim JC, Kim BK. Synthesis of nanosized tungsten carbide powder by the chemical vapor condensation process. Scr Mater 2004;50:969–72. [52] Ryu T, Sohn HY, Han G, Kim YU, Hwang KS, Mena M, et al. Nanograined WC-Co composite powders by chemical vapor synthesis. Metallogr Mater Trans B 2008;39:1–6. [53] Ryu T, Sohn HY, Hwang KS, Fang ZZ. Tungsten carbide nanopowder by plasma-assisted chemical vapor synthesis from WCl6-CH4-H2 mixtures. J Mater Sci 2008;43:5185–92. [54] Sohn HY, Ryu T, Choi JW, Hwang KS, Han G, Choi YJ, et al. The chemical vapor synthesis of inorganic nanopowders. JOM 2007;59:44–9. [55] Yu JH, Kim SY, Lee JS, Ahn KH. In-situ observation of formation of nanosized TiO2 powder in chemical vapor condensation. Nanostruct Mater 1999;12:199–202. [56] Dong XL, Choi CJ, Kim BK. Chemical synthesis of Co nanoparticles by chemical vapor condensation. Scr Mater 2002;47:857–61. [57] Fang ZZ, Wang X, Ryu T, Hwang KS, Sohn HY. Synthesis, sintering, and mechanical properties of nanocrystalline cemented tungsten carbide. Int J Refract Met Hard Mater 2009;27:288–99. [58] Fitzsimmons M, Sarin VK. Comparison of WCl6-CH4-H2 and WF6-CH4-H2 systems for growth of WC coatings. Surf CoatTechnol 1995;76–77:250–5. [59] Ryu T, Sohn HY, Hwang KS, Fang ZZ. Chemical vapor synthesis (CVS) of tungsten nanopowder in a thermal plasma reactor. Int J Refract Met Hard Mater 2009;27: 149–54.

[60] Tang X, Haubner R, Lux B, Kieffer B. Preparation of ultrafine CVD WC powders deposited from WCl6 gas mixtures. J Phys IV 1995;5:1013–20. [61] Won CW, Chun BS, Sohn HY. Preparation of ultrafine tungsten carbide powder by CVD method from WCl6–H2–H2 mixtures. J Mater Res 1993;8:2702–8. [62] Leclercq G, Kamal M, Giraudon JM, Devassine P, Feigenbaum L, Leclercq L, et al. Study of the preparation of bulk powder tungsten carbides by temperature programmed reaction with CH4 + H2 mixtures. J Catal 1996;158:142–69. [63] Medeiros FFP, De Oliveira SA, De Souza CP, Da Silva AGP, Gomes UU, De Souza JF. Synthesis of tungsten carbide through gas–solid reaction at low temperatures. Mater Sci Eng A 2001;315:58–62. [64] Gao L, Kear BH. Low temperature carburization of high surface area tungsten powders. Nanostruct Mater 1995;5:555–69. [65] Kim BK, Ha GH, Lee DW. Sintering and microstructure of nanophase WC/Co hardmetals. J Mater Process Technol 1997;63:317–21. [66] Kim HC, Shon IJ, Yoon JK, Doh JM. Consolidation of ultra fine WC and WC-Co hard materials by pulsed current activated sintering and its mechanical properties. Int J Refract Met Hard Mater 2007;25:46–52. [67] Kim HC, Jeong IK, Shon IJ, Ko IY, Doh JM. Fabrication of WC-8wt.%Co hard materials by two rapid sintering processes. Int J Refract Met Hard Mater 2007;25:336–40. [68] Sommer M, Schubert WD, Zobetz E, Warbichler P. On the formation of very large WC crystals during sintering of ultrafine WC-Co alloys. Int J Refract Met Hard Mater 2002;20:41–50. [69] Herber RP, Schubert WD, Lux B. Hardmetals with “rounded” WC grains. Int J Refract Met Hard Mater 2006;24:360–4. [70] Delanoë A, Lay S. Evolution of the WC grain shape in WC-Co alloys during sintering: effect of C content. Int J Refract Met Hard Mater 2009;27:140–8. [71] Porat R, Berger S, Rosen A. Dilatometric study of the sintering mechanism of nanocrystalline cemented carbides. Nanostruct Mater 1996;7:429–36. [72] Fang Z, Maheshwari P, Wang X, Sohn HY, Griffo A, Riley R. An experimental study of the sintering of nanocrystalline WC-Co powders. Int J Refract Met Hard Mater 2005;23:249–57. [73] Wang X, Fang ZZ, Shon HY. Grain growth during the early stage of sintering of nanosized WC-Co powder. Int J Refract Met Hard Mater 2008;26:232–41. [74] Jia C, Sun L, Tang H, Qu X. Hot pressing of nanometer WC-Co powder. Int J Refract Met Hard Mater 2007;25:53–6. [75] Azcona I, Ordóñez A, Sánchez JM, Castro F. Hot isostatic pressing of ultrafine tungsten carbide–cobalt hardmetals. J Mater Sci 2002;37:4189–95. [76] Wei CB, Song XY, Fu J, Liu XM, Gao Y, Wang HB, et al. Microstructure and properties of ultrafine cemented carbides-differences in spark plasma sintering and sinter-HIP. Mater Sci Eng A 2012;552:427–33. [77] Sunil BR, Sivaprahasam D, Subasri R. Microwave sintering of nanocrystalline WC-12Co: challenges and perspectives. Int J Refract Met Hard Mater 2010;28: 180–6. [78] Agrawal D, Cheng J, Seegopual P, Gao L. Grain growth control in microwave sintering of ultrafine WC-Co composite powder compacts. Powder Metall 2000;43:15–6. [79] Huang SG, Vanmeensel K, Li L, Van der Biest O, Vleugels J. Tailored sintering of VC-doped WC-Co cemented carbides by pulsed electric current sintering. Int J Refract Met Hard Mater 2008;26:256–62. [80] Kim HC, Oh DY, Guojian J, Shon IJ. Synthesis of WC and dense WC-5 vol.% Co hard materials by high-frequency induction heated combustion. Mater Sci Eng A 2004;368:10–7. [81] Kim HC, Shon IJ, Munir ZA. Rapid sintering of ultra-fine WC-10 wt.% Co by high-frequency induction heating. J Mater Sci 2005;40:2849–54. [82] Michalski A, Siemiaszko D. Nanocrystalline cemented carbides sintered by the pulse plasma method. Int J Refract Met Hard Mater 2007;25:153–8. [83] Deorsola FA, Vallauri D, Ortigoza Villalba GA, De Benedetti B. Densification of ultrafine WC-12Co cermets by pressure assisted fast electric sintering. Int J Refract Met Hard Mater 2010;28:254–9. [84] Sivaprahasam D, Chandrasekar SB, Sundaresan R. Microstructure and mechanical properties of nanocrystalline WC-12Co consolidated by spark plasma sintering. Int J Refract Met Hard Mater 2007;25:144–52. [85] Nygren M, Shen Z. On the preparation of bio-, nano- and structural ceramics and composites by spark plasma sintering. Solid State Sci 2003;5:125–31. [86] Cha SI, Hong SH, Kim BK. Spark plasma sintering behavior of nanocrystalline WC-10Co cemented carbide powders. Mater Sci Eng A 2003;351:31–8. [87] Bonache V, Salvador MD, Rocha VG, Borrell A. Microstructural control of ultrafine and nanocrystalline WC-12Co-VC/Cr3C2 mixture by spark plasma sintering. Ceram Int 2011;37:1139–42. [88] Bonache V, Salvador MD, Fernández A, Borrell A. Fabrication of full density near-nanostructured cemented carbides by combination of VC/Cr3C2 addition and consolidation by SPS and HIP technologies. Int J Refract Met Hard Mater 2011;29:202–8. [89] Zhao SX, Song XY, Wei CB, Zhang L, Liu XM, Zhang JX. Effects of WC particle size on densification and properties of spark plasma sintered WC-Co cermet. Int J Refract Met Hard Mater 2009;27:1014–8. [90] Groza JR, Zavaliangos A. Sintering activation by external electrical field. Mater Sci Eng A 2000;287:171–7. [91] Munir ZA, Anselmi-Tamburini U, Ohyanagi M. The effect of electric field and pressure on the synthesis and consolidation of materials: a review of the spark plasma sintering method. J Mater Sci 2006;41:763–77. [92] Zhao JF, Holland T, Unuvar C, Munir ZA. Sparking plasma sintering of nanometric tungsten carbide. Int J Refract Met Hard Mater 2009;27:130–9. [93] Liu WB, Song XY, Wang K, Zhang JX, Zhang GZ, Liu XM. A novel rapid route for synthesizing WC-Co bulk by in situ reactions in spark plasma sintering. Mater Sci Eng A 2009;499:476–81.

X. Ren et al. / Int. Journal of Refractory Metals and Hard Materials 39 (2013) 61–77 [94] Shi XL, Shao GQ, Duan XL, Yuan RZ, Lin HH. Mechanical properties, phases and microstructure of ultrafine hardmetals prepared by WC-6.29Co nanocrystalline composite powder. Mater Sci Eng A 2005;392:335–9. [95] Räthel J, Herrmann M, Beckert W. Temperature distribution for electrically conductive and non-conductive materials during field assisted sintering (FAST). J Eur Ceram Soc 2009;29:1419–25. [96] Liu XM, Song XY, Zhang JX, Zhao SX. Temperature distribution and neck formation of WC-Co combined particles during spark plasma sintering. Mater Sci Eng A 2008;488:1–7. [97] Elfwing M, Norgren S. Study of solid-state sintered fine-grained cemented carbides. Int J Refract Met Hard Mater 2005;23:242–8. [98] Hashe NG, Norgren S, Andrén HO, Neethling JH, Berndt PR. Reduction of carbide grain growth in WC-VC-Co by sintering in a nitrogen atmosphere. Int J Refract Met Hard Mater 2009;27:20–5. [99] Yamamoto T, Ikuhara Y, Watanabe T, Sakuma T, Taniuchi Y, Okada K, et al. High resolution microscopy study in Cr3C2-doped WC-Co. J Mater Sci 2001;36: 3885–90. [100] Jaroenworaluck A, Yamamoto T, Ikuhara Y, Sakuma T, Taniuchi T, Okada K, et al. Segregation of vanadium at the WC/Co interface in VC-doped WC-Co. J Mater Res 1998;13:2450–2. [101] Lay S, Hamar-Thibault S, Lackner A. Location of VC in VC, Cr3C2 codoped WC-Co cermets by HREM and EELS. Int J Refract Met Hard Mater 2002;20:61–9. [102] Sun L, Jia CC, Cao RJ, Lin CG. Effects of Cr3C2 additions on the densification, grain growth and properties of ultrafine WC-11Co composites by spark plasma sintering. Int J Refract Met Hard Mater 2008;26:357–61. [103] Weidow J, Norgren S, Andrén HO. Effect of V, Cr and Mn additions on the microstructure of WC-Co. Int J Refract Met Hard Mater 2009;27:817–22. [104] Hashe NG, Neethling JH, Berndt PR, Andrén HO, Norgren S. A comparison of the microstructures of WC-VC-TiC-Co and WC-VC-Co cemented carbides. Int J Refract Met Hard Mater 2007;25:207–13. [105] Andrén HO. Microstructures of cemented carbides. Mater Des 2001;22:491–8. [106] Hashe NG, Norgren SM, Andrén HO, Neethling JH. Characterization of WC-(W, V) C-Co made from pre-alloyed (W, V)C. Int J Refract Met Hard Mater 2009;27: 229–33. [107] Ou XQ, Xiao DH, Shen TT, Song M, He YH. Characterization and preparation of ultra-fine grained WC-Co alloys with minor La-additions. Int J Refract Met Hard Mater 2012;31:266–73. [108] Zhang FM, Shen J, Sun JF. The effect of phosphorus additions on densification, grain growth and properties of nanocrystalline WC-Co composites. J Alloy Compd 2004;385:96–103. [109] Jia K, Fischer TE, Gallois B. Microstructure, hardness and toughness of nanostructured and conventional WC-Co composites. Nanostruct Mater 1998;10:875–91. [110] Gleiter H. Materials with ultrafine microstructures: retrospective and perspectives. Nanostruct Mater 1992;1:1–19. [111] Gant AJ, Gee MG, Roebuck B. Rotating wheel abrasion of WC/Co hardmetals. Wear 2005;258:178–88. [112] Pirso J, Viljus M, Letunovitš S. Friction and dry sliding wear behavior of cermets. Wear 2006;260:815–24. [113] Fischer UKR, Hartzell ET, Akerman JGH. Cemented carbide body with a binder phase gradient and method of making the same. US Patent No. 4,820,482, 1989. [114] Hartzell ET, Akerman JGH, Fischer UKR. Cemented carbide body used preferably for abrasive rock drilling and mineral cutting. US Patent No. 5,401,461, 1995. [115] Akerman JGH, Fischer UKR, Hartzell ET. Cemented carbide body with extra tough behavior. US Patent No. 5,453,241, 1995. [116] Withers PJ, Bhadeshia HKDH. Residual stress. Part 1—Measurement techniques. Mater Sci Technol 2001;17:355–65. [117] Withers PJ, Bhadeshia HKDH. Residual stress. Part 2—Nature and origins. Mater Sci Technol 2001;17:366–75. [118] Fang Z, Lockwood G, Griffo A. A dual composite of WC-Co. Metallogr Mater Trans A 1999;30:3231–8. [119] Nemeth BJ, Grab GP. Preferentially binder enriched cemented carbide bodies and method of manufacture. US Patent No. 4,610,931, 1986. [120] Zhang L, Wang YJ, Yu XW, Chen S, Xiong XJ. Crack propagation characteristic and toughness of functionally graded WC-Co cemented carbide. Int J Refract Met Hard Mater 2008;26:295–300. [121] Li T, Li Q, Fuh JYH, Yu PC, Lu L. Two-material powder injection molding of functionally graded WC-Co components. Int J Refract Met Hard Mater 2009;27: 95–100. [122] Yohe WC. Coated carbide cutting tool insert. US Patent No. 4,548,786, 1985. [123] Huang ZQ, He YH, Li PL. Finite element analysis of thermal stresses in functionally graded cemented carbides. Powder Metall 2010;53:278–84. [124] Wang X, Hwang KS, Koopman M, Fang ZZ, Zhang L. Mechanical properties and wear resistance of functionally graded WC-Co. Int J Refract Met Hard Mater 2013;36:46–51. [125] Liu Y, Wang H, Long Z, Liaw P, Yang J, Huang B. Microstructural evolution and mechanical behaviors of graded cemented carbides. Mater Sci Eng A 2006;426: 346–54.

77

[126] Lin W, Bai XD, Ling YH, Jiang ZH, Xie ZP. Fabrication and properties of axisymmetric WC/Co functionally graded hard metal via microwave sintering. Mater Sci Forum 2003;423–425:55–8. [127] Konyashin I, Hlawatschek S, Ries B, Lachmann F, Sologubenko A, Weirich T. A new approach to fabrication of gradient WC-Co hardmetals. Int J Refract Met Hard Mater 2010;28:228–37. [128] Put S, Vleugels J, Van der Biest O. Functionally graded WC-Co materials produced by electrophoretic deposition. Scr Mater 2001;45:1139–45. [129] Zhang L, Zhang G, Zhang J, Zhou M. Study of graded cemented carbide prepared by stack moulding and spark plasma sintering (SPS). Rare Metal Mater Eng 2006;35:70–3. [130] Kieback B, Neubrand A, Riedel H. Processing techniques for functionally graded materials. Mater Sci Eng A 2003;362:81–106. [131] Eso OO, Fang ZZ, Griffo A. Liquid phase sintering of functionally graded WC-Co composites. Int J Refract Met Hard Mater 2005;23:233–41. [132] Fang ZZ. Functionally graded cemented tungsten carbide. US Patent No. 7,699,904, 2010. [133] Fan P, Eso OO, Fang ZZ, Sohn HY. Effect of WC particle size on Co distribution in liquid-phase-sintered functionally graded WC-Co composite. Int J Refract Met Hard Mater 2008;26:98–105. [134] Tokita M. Large-size WC/Co functionally graded materials fabricated by spark plasma sintering (SPS) method. Mater Sci Forum 2003;423–425:39–44. [135] Fan P, Guo J, Fang ZZ, Prichard P. Design of cobalt gradient via controlling carbon content and WC grain size in liquid-phase-sintered WC-Co composite. Int J Refract Met Hard Mater 2009;27:256–60. [136] Colin C, Guipont V, Delannay F. Equilibrium distribution of liquid during sintering of assemblies of WC/Co cermets. Metallogr Mater Trans A 2007;38:150–8. [137] Fan P, Guo J, Fang ZZ, Prichard P. Effects of liquid-phase composition on its migration during liquid-phase sintering of cemented carbide. Metallogr Mater Trans A 2009;40:1995–2006. [138] Fan P, Fang ZZ, Guo J. A review of liquid phase migration and methods for fabrication of functionally graded cemented tungsten carbide. Int J Refract Met Hard Mater 2013;36:2–9. [139] Engqvist H, Botton GA, Axén N, Hogmark S. Microstructure and abrasive wear of binderless carbides. J Am Ceram Soc 2000;83:2491–6. [140] Tracey VA. Nickel in hardmetals. Int J Refract Met Hard Mater 1992;11:137–49. [141] Kim HC, Shon IJ, Yoon JK, Doh JM, Munir ZA. Rapid sintering of ultrafine WC-Ni cermets. Int J Refract Met Hard Mater 2006;24:427–31. [142] Wittmann B, Schubert WD, Lux B. WC grain growth and grain growth inhibition in nickel and iron binder hardmetals. Int J Refract Met Hard Mater 2002;20: 51–60. [143] Rong HY, Peng ZJ, Ren XY, Wang CB, Fu ZQ, Qi LH, et al. Microstructure and mechanical properties of ultrafine WC-Ni-VC-TaC-cBN cemented carbides fabricated by spark plasma sintering. Int J Refract Met Hard Mater 2011;29:733–8. [144] Correa EO, Santos JN, Klein AN. Microstructure and mechanical properties of WC-Ni-Si based cemented carbides developed by powder metallurgy. Int J Refract Met Hard Mater 2010;28:572–5. [145] Kim HC, Park HK, Jeong IK, Ko IY, Shon IJ. Sintering of binderless WC-Mo2C hard materials by rapid sintering process. Ceram Int 2008;34:1419–23. [146] Huang SG, Vanmeensel K, Van der Biest O, Vleugels J. Binderless WC and WC-VC materials obtained by pulsed electric current sintering. Int J Refract Met Hard Mater 2008;26:41–7. [147] Poetschke J, Richter V, Holke R. Influence and effectivity of VC and Cr3C2 grain growth inhibitors on sintering of binderless tungsten carbide. Int J Refract Met Hard Mater 2012;31:218–23. [148] Mukhopadhyay A, Chakravarty D, Basu B. Spark plasma-sintered WC-ZrO2-Co nanocomposites with high fracture toughness and strength. J Am Ceram Soc 2010;93:1754–63. [149] Basu B. Development of WC-ZrO2 nanocomposites by spark plasma sintering. J Am Ceram Soc 2004;87:317–9. [150] Malek O, Lauwers B, Perez Y, Baets PD, Vleugels J. Processing of ultrafine ZrO2 toughened WC composites. J Eur Ceram Soc 2009;29:3371–8. [151] Zheng D, Li X, Ai X, Yang C, Li Y. Bulk WC-Al2O3 composites prepared by spark plasma sintering. Int J Refract Met Hard Mater 2012;30:51–6. [152] Huang B, Chen LD, Bai SQ. Bulk ultrafine binderless WC prepared by spark plasma sintering. Scr Mater 2006;54:441–5. [153] Kim HC, Shon IJ, Garay JE, Munir ZA. Consolidation and properties of binderless sub-micron tungsten carbide by field-activated sintering. Int J Refract Met Hard Mater 2004;22:257–64. [154] Kim HC, Yoon JK, Doh JM, Ko IY, Shon IJ. Rapid sintering process and mechanical properties of binderless ultra fine tungsten carbide. Mater Sci Eng A 2006;435–436: 717–24. [155] Zhang JX, Zhang GZ, Zhao SX, Song XY. Binder-free WC bulk synthesized by spark plasma sintering. J Alloy Compd 2009;479:427–31.