Complexions in WC-Co cemented carbides

Complexions in WC-Co cemented carbides

Accepted Manuscript Complexions in WC-Co cemented carbides Xingwei Liu, Xiaoyan Song, Haibin Wang, Xuemei Liu, Fawei Tang, Hao Lu PII: S1359-6454(18)...

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Accepted Manuscript Complexions in WC-Co cemented carbides Xingwei Liu, Xiaoyan Song, Haibin Wang, Xuemei Liu, Fawei Tang, Hao Lu PII:

S1359-6454(18)30119-8

DOI:

10.1016/j.actamat.2018.02.018

Reference:

AM 14374

To appear in:

Acta Materialia

Received Date: 25 September 2017 Revised Date:

6 February 2018

Accepted Date: 12 February 2018

Please cite this article as: X. Liu, X. Song, H. Wang, X. Liu, F. Tang, H. Lu, Complexions in WC-Co cemented carbides, Acta Materialia (2018), doi: 10.1016/j.actamat.2018.02.018. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

ACCEPTED MANUSCRIPT Graphical Abstract The existence of complexions is dominated by the formation energy and misfit at the interface between the complexion and WC. The stability of complexions can be tailored by matching inhibitor and carbon content, and both the high

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complexions in the cemented carbides.

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fracture toughness and high strength can be achieved by controlling

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Complexions in WC-Co cemented carbides Xingwei Liu, Xiaoyan Song*, Haibin Wang, Xuemei Liu, Fawei Tang, Hao Lu

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College of Materials Science and Engineering, Key Laboratory of Advanced Functional Materials, Education Ministry of China, Beijing University of Technology, Beijing 100124, P. R. China

Abstract

The studies are focused on the complexions in the WC-Co cemented carbides,

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which represent the phase-like interfacial state at the WC/Co interfaces. A series of experiments were designed to investigate the formation, growth and transformation behaviors of the complexions. It was found that the complexions with cubic structure

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can be stabilized by quenching or by dissolution of specific elements added in form of carbides as grain growth inhibitors. The characteristics and stabilization conditions of the complexions, as well as their influences on the interfacial structures and mechanical properties of the cemented carbides, were studied in detail for additions of various inhibitors such as VC, Cr3C2, TiC, TaC and NbC. The results disclose that the existence

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of complexions is dominated by their formation energy and misfit at the interface between the complexion and WC. The stabilization of complexions, on one hand, inhibits WC grain growth effectively, on the other hand separates the bonding of WC and Co, thus causes decrease of the resistance against intergranular fracture of the

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cemented carbides. It is proposed that tailoring the stability of complexions by matching inhibitor and carbon content can lead to both the high fracture toughness and high

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strength of the cemented carbides. Key words: complexion, cemented carbide, grain growth inhibitor, stability, interface coherency

1. Introduction

The WC-Co cemented carbides are generally considered as the most representative type of hard alloys that have been widely used in industries, such as cutting, machining, mining and drilling tools, owing to their high hardness, wear resistance and fracture strength [1-3]. To extend applications as tool materials for high-speed precision

ACCEPTED MANUSCRIPT machining, further increase of hardness and fracture strength is highly demanded for WC-Co cemented carbides. For this purpose, plenty of efforts have been paid to decrease WC grain size to ultrafine (submicron) or nanometer scale [4-6]. For example, by addition of VC and Cr3C2 powders as the grain growth inhibitors (GGI), the

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cemented carbides with an ultrafine microstructure (WC grain size on the submicron scale) were prepared by the conventional powder metallurgy methods [4,5]. The materials have higher hardness and wear resistance as compared to those with the

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micron grain size. Adding VC with an appropriate amount, the nanocrystalline cemented carbide was obtained by spark plasma sintering, which had greatly increased hardness and good fracture toughness [6]. It has been commonly recognized that to

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prepare the cemented carbides with fine grain size by means of powder metallurgy, the inhibitors are indispensable to be used in the powders for sintering to control the grain growth. Almost all the GGI reported in the literature are carbides of refractory metals, such as VC, Cr3C2, TiC, TaC and NbC [4,7-14], among which VC and Cr3C2 are the most frequently used ones. It was found that with the decrease of grain size owing to the

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presence of inhibitors, the hardness of the cemented carbides is generally increased, however, it is uncertain for the increase of the fracture toughness and strength [7,8]. The influences of GGI on the comprehensive properties of cemented carbides, as well as the

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mechanisms of the interactions between GGI and WC hard phase and Co binder on the microscale, have not been sufficiently understood yet.

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The effect of GGI additions on the phase constitution in the cemented carbides has

been seldom studied in the previous work. Weidow et al. reported the segregation phenomena at the WC/Co interfaces in the cemented carbides containing GGI [15,16]. To distinguish the interfacial layers from the bulk phases in the cemented carbides, Johansson and Wahnstroem introduced the term “complexion” in 2016 in their calculations of diagrams for phase boundaries in the V-doped cemented carbides [17]. The concept and category of “complexion” were described systematically in 2014 in the overview by Cantwell et al [18]. As proposed [18,19], “A complexion, concisely

ACCEPTED MANUSCRIPT defined, is interfacial material or strata that is in thermodynamic equilibrium with the abutting phase(s) and has a stable, finite thickness that is typically on the order of 0.2–2 nm”. With this definition, many complexion-related phenomena can be found in the literature [20]. The complexions exist in various materials and influence the thermal

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[21], electrochemical [22] electrical [23], magnetic [24] and mechanical properties [25]. Khalajhedayati et al. found that the formation of complexions could change the state of interfacial bonding, thus improved the combination of grains at the interface and the

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mechanical properties of the material [25].

Investigations on the complexions in the WC-Co cemented carbides are still very limited. Its formation mechanism, stability and effects on the microstructures and

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mechanical properties of cemented carbides have been remained unclear so far. In the present work, our attentions are focused on the complexions in the WC-Co cemented carbides containing GGI. The characteristics and stabilization conditions of the complexions, as well as their influences on the grain structure and mechanical properties of the cemented carbides with addition of various GGI, will be studied in

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detail. The paper is composed mainly of four parts. Following the section of materials design and experimental methods, the complexions are firstly investigated by freezing the state of formation at high temperatures by quenching (Section 3.1). The

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complexions are examined with regard to the crystal structure, combinative state with the matrix and formation capability with different GGI additions (e.g. VC, Cr3C2, TiC,

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TaC and NbC). Then the complexion stability and its correlation with the mechanical properties of the cemented carbide are studied and the related mechanisms are proposed (Section 3.2). Finally, the influence of the complexions on the crack-microstructure interactions and the effects of GGI on the interface structures and mechanical properties of the WC-Co cemented carbides are discussed (Section 4). It is considered that the results of our studies may provide a strategy to optimize the microstructures and achieve high comprehensive properties of the cemented carbides by tailoring the stability of complexions.

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2.

Experimental The WC-Co composite was used as the base material of the cemented carbide bulk.

Firstly, the WC-12wt%Co composite powder was synthesized by the in situ reactions of

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metal oxides and carbon black powders as raw materials, and the processing details were described in our previous work [26]. A variety of inhibitors including VC, Cr3C2, TiC, TaC, NbC and VC+Cr3C2 (in a ratio of 1:1) with a mean particle size of ∼ 500 nm

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and content of 1.0wt% were added into the WC-Co composite powder separately, to prepare the cemented carbide bulk specimens containing different GGI. The powders were further mixed with polyethylene glycol and compressed into rectangular solids,

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which were then sintered in the furnace filled with argon gas at 1390°C for 60min. Subsequently, the sintered bulk specimens were quenched into cold water to examine the phase constitution in the cemented carbide at the high temperature. As a parallel processing, the VC and Cr3C2 powders were added in the WC-Co composite powder and then the powder mixture was sintered by hot isostatic pressing (sinter-HIP) at

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1390°C for 60min in the argon atmosphere with a pressure of 10 MPa. The bulk specimens prepared by sinter-HIP were cooled in furnace to the room temperature. The morphology of grain structures in the bulk specimens was observed by the

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scanning electron microscope (SEM, JSM 6500). The electron back scattered diffraction (EBSD) analysis was performed using the EDAX TSL Hikari EBSD detector with a

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scanning step of 10 nm. The microstructures were examined by the transmission electron microscopy (TEM) and high-resolution TEM (HRTEM) operated at 200 kV (JEM-2100F). The local composition was measured by the energy dispersive x-ray (EDX) spectroscopy performed on the microstructures. The hardness was tested by the Vickers hardness tester with a load of 30 kg. The fracture toughness (KIc) was determined by measuring the crack length then calculating with the equation KIc = 0.0028(HP/L)1/2 [27], where H is the indentation hardness, P is the load and L is the total crack length. The transverse rupture strength (TRS) was measured by the three-point

ACCEPTED MANUSCRIPT bending method with dimensions of the specimen as 20×6.5×5.25 mm according to the standard of ISO3327:2009. To guarantee the data reliability, at least five specimens

3.

Results

3.1 Formation of complexions in cemented carbides 3.1.1 Quenched cemented carbide without GGI

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were used to measure each property.

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The microstructure and phase constitution at the high temperature can be preserved till room temperature by quenching. The microstructure of the quenched cemented carbide bulk specimen is shown in Fig. 1. In the locally enlarged image of the

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microstructure, as shown in Fig. 1(b), the complexion with partially ordered structure is observed at the boundary of WC and Co, which is clearly different from the structures of the abutting two phases.

By indexing the Fast Fourier Transformation (FFT) patterns and combining with the HRTEM analysis, the crystallographic features of the complexion were obtained, as

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indicated in Figs. 1(b)-(d) at the WC/Co interfaces. The complexion has 2~4 atomic layers, as shown in Figs. 1(c) and (d), which mainly distributes at WC(10-10)/Co and WC(0001)/Co interfaces with coherent relationship with the WC grains. It has a cubic

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crystal structure with the lattice parameter as a=4.226Å. Referring to the literature [28], the complexion formed in the quenched cemented carbide is determined as fcc-WCx.

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It was previously reported that the metastable fcc-WCx, with an atomic ratio of

W:C in a range of 1.7:1 ~ 1:1 (i.e. x = 0.59 ~ 1.0), normally formed and existed at temperatures above 2800K [29]. Due to a big difference in solubility and diffusion rate of W and C in the liquid phase during sintering, the inhomogeneous distribution of composition occurs at the surface of WC grains, which may lead to the formation of WCx at WC/Co interfaces in the process of liquid-state sintering. Thus, the formed WCx at high temperatures can be kept till room temperature by quenching. From the analysis in Figs. 1(c) and (d), the misfits of WC and WCx at the WC(10-10)/WCx(200) and

ACCEPTED MANUSCRIPT WC(0001)/WCx(111) interfaces are 1.17% and 0.77%, respectively. The low misfit suggests low interface energy at these phase boundaries. The WC(0001)/WCx(111) interface has a lower lattice misfit than the WC(10-10)/WCx(200) interface, which can explain that the WCx

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layer is a little thicker on the WC(0001) basal plane.

Figure 1. Microstructures of the WC-Co cemented carbide (without GGI) quenched from 1390oC: (a) distribution of WC and Co phases; (b) complexion at WC/Co interface;

3.1.2 Quenched cemented carbide with VC

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(c) WCx complexion on WC(10-10) plane; (d) WCx complexion on WC(0001) plane.

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As one of the most popular GGI, VC plays an important role in controlling WC grain growth [10,30]. A representative theory proposed by Kawakami et al. indicated that the inhibition of VC on grain growth was realized through adsorbing and desorbing of VC on WC surfaces during sintering [30]. However, the details, particularly on the atomic scale, are lacking.

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As shown in Fig. 2, in the quenched cemented carbide with addition of VC, the complexions are observed distributing on the WC(0001) and WC(10-10) planes at WC/Co interfaces, which have a morphology of nanofilm and a fcc structure. The STEM-EDAX

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analyses in Figs. 2(b) and (c) indicate that there is a little higher V concentration at the WC/Co interface than that inside the phases, implying that the complexion has a

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composition of (W,V)Cx. From the details shown in Fig. 2(d), the misfit at the WC(0001)/(W,V)Cx(111) interface is less than 0.1%, which is much smaller than that in the cemented carbide without VC (see Fig. 1d). Again, it is found that (W,V)Cx is thicker on the WC(0001) plane than on the WC(10-10) plane (Figs. 2d and e). Due to the existence of complexions at interfaces, the migration of WC grain boundaries is effectively hindered in the sintering process. Moreover, the orientation relationship of WC and Co can be changed by the complexions, which will be discussed in detail in Section 3.2.4 and Section 4.

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Figure 2. Microstructures of the WC-Co cemented carbide with addition of VC quenched from 1390oC: (a) distribution of complexions; (b) STEM-EDX analysis at WC(0001)/Co interface; (c) STEM-EDX analysis at WC(10-10)/Co interface; (d) (W,V)Cx

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complexion on WC(0001) plane; (e) (W,V)Cx complexion on WC(10-10) plane.

3.1.3 Quenched cemented carbide with Cr3C2

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In the quenched cemented carbide with Cr3C2, the (W,Cr)Cx complexion was not detected at the phase boundaries of WC and Co. Instead, as seen in Fig. 3, the WCx complexion distributes discontinuously along the WC(10-10)/Co interface. The

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STEM-EDAX analysis shows that the Cr concentration in Co binder is higher than that in the WC grain, and no obvious Cr segregation is found at WC/Co interface in the quenched specimen. This indicates that the interfacial segregation of Cr does not occur at the liquid-state sintering temperature. It has been found that as compared with other metal elements, Cr has higher solubility in Co binder in the liquid-state sintering

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process [8,31], which is not favorable for formation and stabilization of the ordered complexion at the phase boundary [32].

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Figure 3. Microstructure of the WC-Co cemented carbide with addition of Cr3C2 quenched from 1390oC: (a) WCx complexion at WC(10-10)/Co interface; (b) STEM-EDX

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analysis at WC(10-10)/Co interface.

3.1.4 Quenched cemented carbide with both VC and Cr3C2 In the quenched specimen with co-addition of VC and Cr3C2, as shown in Fig. 4,

the complexion on the WC(10-10) prismatic plane seems a little thicker than that in the specimen with single addition of VC or Cr3C2 (see Fig. 2e and Fig. 3a). As indicated by the STEM-EDAX analysis, there are higher concentrations of V and Cr at WC/Co interface, which implies that the concurrent dissolution of VC and Cr3C2 promotes Cr

ACCEPTED MANUSCRIPT segregation at the phase boundary of WC and Co during liquid-state sintering, as compared with the single addition of Cr3C2 (see Fig. 3b). Consequently, the stability of the complexion is enhanced by both V and Cr segregation at the phase boundary. Therefore, the complexion on the WC(10-10) prismatic plane in the specimen with

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co-addition of VC and Cr3C2 is thicker than that with single addition of VC or Cr3C2.

Figure 4. Microstructures of the WC-Co cemented carbide with co-addition of VC and

at WC(10-10)/Co interface.

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3.1.5 Quenched cemented carbide with TiC

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Cr3C2 quenched from 1390oC: (a) complexion at WC(0001)/Co interface; (b) complexion

It was reported by Lay et al. [33], Weidow et al. [16] and Meingast et al. [34] separately that in the cemented carbides containing TiC, the (W,Ti)Cx layer formed and distributed on the WC basal plane at the WC/Co phase boundary. This is confirmed in our experiments. As shown in Fig. 5(a), the plate-like WC grains are observed in the

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quenched cemented carbide specimen with addition of TiC. The plate shape results from the anisotropic growth of WC grains, which is due to the migration of (0001) basal plane along [0001] direction is hindered by the existence of (W,Ti)Cx complexion on the

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WC basal plane, as indicated by the HRTEM analysis in Figs. 5(b) and (c). From the measurements in Fig. 5(c), it can be evaluated that the corresponding

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atomic planes of the complexion and WC have good matching, e.g. the WC(10-10)/(W,Ti)Cx(200) interface has a very small misfit (< 0.1%). Combining the relatively lowest formation energy (as compared in Table 1) and the full interface coherency (Fig. 5c), it is considered that the (W,Ti)Cx complexion can exist stably on the WC(0001) basal plane. However, on the WC prismatic planes, there is almost no complexion found. The results indicate that due to the inhibiting effect of (W,Ti)Cx that exists stably on the basal plane, the WC grain prefers to grow along the [10-10]

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Figure 5. Microstructures of the quenched WC-Co cemented carbide with addition of

3.1.6 Quenched cemented carbide with TaC

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TiC: (a) plate-like WC grains; (b,c) (W,Ti)Cx complexions on WC(0001) basal plane.

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The microstructure of the quenched WC-Co cemented carbide with addition of TaC is shown in Fig. 6. It can be seen that the morphology of WC grains is distinctly different from that in the specimen with TiC addition. The round borders of WC grains

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are observed at the WC/Co interfaces where Ta is detected within the WC grains (see the local composition analysis in Fig. 6a). From the STEM-EDX analysis across the WC grain, WC/Co interface and Co binder, as shown in Fig. 6(b), it is obvious that Ta dissolves in WC and (W,Ta)C is formed in the vicinity of the round border of the WC grain. In contrast, there is little Ta detected at the WC/Co interface or in Co, indicating

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that the complexion does not exist at the interface. This is confirmed by the HRTEM analysis in Fig. 6(c), where the interface between the (W,Ta)C-hcp and Co-fcc is observed. Due to the dissolution of Ta in WC, the spacing of the (W,Ta)C(10-10) plane is

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decreased to 2.308Å from the standard value 2.517Å for the WC(10-10) plane. Consequently, the misfit between the (W,Ta)C solid solution and complexion is

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increased, thus the complexion is hard to form on the surface of (W,Ta)C due to the high interface energy and formation energy (Table 1).

Figure 6. Microstructures of the quenched WC-Co cemented carbide with addition of TaC: (a) morphology of WC grains and local composition analysis; (b) STEM-EDX analysis along the white line from WC to WC/Co interface then to Co; (c) interface between (W,Ta)C and Co.

ACCEPTED MANUSCRIPT 3.1.7 Quenched cemented carbide with NbC The characteristics of the microstructure in the quenched cemented carbide with addition of NbC can be observed in Fig. 7. It is found that the WC grains are equiaxed with straight edges. Moreover, the (W,Nb)Cx complexion is very thin with 2∼3 atomic

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layers. This is consistent with the report from Yamamoto et al. [35] that there was a little segregation of Nb at the WC/Co interfaces in the cemented carbide containing NbC. In the present work, it is further found that the (W,Nb)Cx complexions distribute on both

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the WC basal and prismatic planes, as the examples shown in Figs. 7(b) and (c), respectively. Referring to the calculation results in Table 1, (W,Nb)Cx has a relatively lower formation energy. However, the misfit between the corresponding planes of the

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complexion and WC is increased due to the dissolution of Nb in WC in the vicinity of WC/Co interface. Therefore, it limits the growth of (W,Nb)Cx complexion on the surface of WC. Since the complexions can exist on both the WC basal and prismatic planes, the WC grains grow isotropically and have an equiaxed shape in the

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microstructure.

Figure 7. Microstructures of the quenched WC-Co cemented carbide with addition of NbC: (a) morphology of WC grains; (b) (W,Nb)Cx complexion on WC(0001) basal

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plane; (c) (W,Nb)Cx complexion on WC(10-10) prismatic plane.

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3.1.8 Evaluation of formation capability of complexions To understand the formation and stability of complexions in the viewpoint of energy,

the formation energies of the refractory metal carbides with cubic structure were calculated by the first principles method. Moreover, the misfits between the cubic carbides and WC at the interface were evaluated. The first principles calculation implementation of the Cambridge serial total energy package (CASTEP) [36] was applied. The generalized gradient approximation (GGA) [37] with the Perdew-Burke -Eruzerh (PBE) formula [38] were used for the exchange-correlation potential. The

ACCEPTED MANUSCRIPT cubic lattice cell of WCx was estabilished in advance, then the W atoms were substituted by the M (M=V, Cr, Ti, Ta, Nb and V+Cr) atoms with a ratio of 1/4. The formation energy, ∆Ef, is described as: ∆Ef = E(W,M)Cx-(nWµW+nCµC+nMµM)

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where E(W,M)Cx is the total energy of the (W,M)Cx cell, nW, nC and nM are the numbers of W, C and M atoms in the doped WCx cell, µW, µC and µM are the chemical potentials, respectively, each of which is equal to the energy per atom correspondingly in the unit

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cell of (W,M)Cx.

To simplify the construction of the lattice cell, x=1 was used as a reference for comparison. Table 1 shows the calculation results. It is seen that by doping of Ti, Nb, V

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and V+Cr in the cubic cell, the formation energies of the cubic carbides are reduced. Among the cubic carbides for calculation, the (W,Ti)Cx carbide has the lowest formation energy. Moreover, the lattice parameter of the cubic crystal structure changes with the doping element. It implies that the complexion can be stabilized by doping due to the decrease of the formation energy. In addition, the change in the lattice parameter due to

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doping can influence the misfit between the cubic carbide (W,M)Cx and WC at the interface. From the calculated lattice parameters, as compared with the misfits at the WCx(111)/WC(0001) and WCx(200)/WC(10-10) interfaces, due to V doping the misfits at the

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(W,V)Cx(111)/WC(0001) and (W,V)Cx(200)/WC(10-10) interfaces are reduced by 15% and 10%, respectively. As for Cr doping, although the misfits between (W,Cr)Cx and WC at

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the basal and prismatic planes are smaller than those of (W,V)Cx and WC, the formation energy of the (W,Cr)Cx carbide is much larger than that of (W,V)Cx, thus the formation capability of the (W,Cr)Cx complexion is confined by the high formation energy. It is interesting to find that the formation energy of the (W,V,Cr)Cx carbide is lower

than that of both (W,V)Cx and (W,Cr)Cx. Particularly, from the calculated lattice parameters, the misfits at the (W,V,Cr)Cx(111)/WC(0001) and (W,V,Cr)Cx(200)/WC(10-10) interfaces are further reduced as compared with those of (W,V)Cx/WC and (W,Cr)Cx/WC correspondingly. Therefore, it can be considered that by specific doping,

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As compared in Table 1, among all the cubic carbides, (W,V,Cr)Cx has the lowest formation energy and the smallest misfit with WC at the basal and prismatic planes, implying that it has the highest formation capability. By contrast, the (W,Cr)Cx carbide

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has the highest formation energy, indicating that the complexion is hard to form in the cemented carbide containing Cr3C2. The calculations agree well with the experimental results in Sections 3.1.2−3.1.4 on the characteristics of the complexions in the cemented

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carbides with single addition of VC or Cr3C2 and mixed addition of VC and Cr3C2. The behavior of complexions in the cemented carbides with additions of TiC, TaC and NbC will be discussed in Section 4.

(W,V)Cx

(W,Cr)Cx

(W,V,Cr)Cx

(W,Ti)Cx

(W,Ta)Cx

(W,Nb)Cx

2.167

1.576

2.778

1.354

0.282

2.426

1.481

4.323

4.244

4.363

4.432

4.403

4.375

4.338

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Formation energy (eV) Lattice parameter (Å)

WCx

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Carbide

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Table 1. The calculated formation energy and lattice parameter of the cubic carbides

3.2 Relation of carbon content, complexion and mechanical properties It is known from Section 3.1 that addition of certain GGIs in the cemented carbides

can increase the stability of complexions. This enlightens that with proper composition design of GGI, the complexions may form in the cemented carbides prepared by the conventional slow cooling instead of quenching treatment. In this section, take the WC-12Co bulk material with co-addition of VC and Cr3C2 as an example, the correlations of complexion with composition and mechanical properties of the cemented

ACCEPTED MANUSCRIPT carbide will be studied. The specimens were prepared by sinter-HIP which were cooled in the furnace instead of quenching. Since the content of carbon is significant for the composition of complexion and the transformation from complexion to WC, in this section the attention is firstly paid to study the effects of carbon content on the

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formation and evolution of the complexion, and then its influences on the mechanical properties of the cemented carbide. A series of carbon contents in a range of 16.71-16.85wt% are designed close to the theoretical stoichiometric ratio in the

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synthesis of the initial WC-12wt%Co composite powder. In the following, the changes of mechanical properties of the cemented carbides with carbon content are characterized. Then two representative specimens with relatively low (16.71wt%) and high (16.82wt%)

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carbon contents are selected to examine the related complexions.

3.2.1 Mechanical properties at different carbon contents

Fig. 8(a) shows the mean grain sizes and mechanical properties of the cemented carbides prepared by sinter-HIP with co-addition of VC and Cr3C2 in a ratio of 1:1 and a

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total amount of 1.0wt% and different carbon contents. It can be seen that the specimen with lower carbon content has a little higher hardness, which is consistent with the smaller grain size. As marked by the shadows, under the same condition, the specimen

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with lower carbon content has lower fracture toughness and strength. In contrast, the combination properties of high toughness and strength are obtained in the specimen

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with relatively higher carbon content. However, the properties are decreased with further increase of carbon content due to the excessive free carbon. The effects of carbon on mechanical properties of the WC-Co cemented carbides containing GGI and without any addition of GGI [26] are distinctly different, which implies that the properties are dependent on the integration of GGI and carbon content. It should be noted that the mechanical properties of the cemented carbides are also influenced by the nature, size and distribution of the flaws. However, in the present experiments, all the specimens whose properties are compared were prepared with the same processing.

ACCEPTED MANUSCRIPT From the SEM observations, the flaws in the specimens can be considered on the similar level. Therefore, the present study is focused on the effects of microstructures on the mechanical properties of the cemented carbides. Figs. 8(b-e) show the representative microstructures with the same GGI but low and

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high carbon contents, respectively, corresponding to the specimens whose properties are marked with blue and yellow shadows in Fig. 8(a). In the specimen with low carbon content, as shown in Figs. 8(b) and (d), the WC grains have morphology with sharp

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corners and facets on the surface. However, with high carbon content, as seen in Figs. 8(c) and (e), the WC grains have nearly straight edges. As will be studied in Sections 3.2.2 and 3.2.3, the different WC morphology results from the effect of carbon content

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on the formation and evolution of complexions at the WC/Co interfaces.

Figure 8. The mean grain size, mechanical properties and microstructures of the cemented carbides prepared by sinter-HIP with co-addition of VC and Cr3C2 and different carbon contents: (a) changes of the mean grain size, hardness, fracture

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toughness and strength with carbon content; (b,d) microstructures with low carbon content (16.71wt%), corresponding to the properties marked with blue shadow in (a); (c,e) microstructures with high carbon content (16.82wt%), corresponding to the

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properties marked with yellow shadow in (a).

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3.2.2 Complexion in cemented carbide with low carbon Fig. 9 shows the microstructure details in the cemented carbide with co-addition of

VC and Cr3C2 and low carbon content (16.71wt%). Viewed in the WC[0001] direction, the complexion is observed existing at the WC/Co interface with a thickness of 1~2 nm (Fig. 9b). It is determined as fcc-(W,V,Cr)Cx by HRTEM and STEM-EDAX analyses. This implies that different from the stabilization of WCx by quenching, the complexion can be stabilized by the dissolved V and Cr in the WCx lattice. When observed in the WC[1-210] direction, as shown in Figs. 9(c-e), there exists coherent relationship between

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and

(W,V,Cr)Cx,

where

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corresponding

atomic

planes

are

WC(0001)/(W,V,Cr)Cx(111) (Fig. 9d) and WC(10-10)/(W,V,Cr)Cx(200) (Fig. 9e), respectively, with a very small misfit of less than 0.3%. The coherency between (W,V,Cr)Cx and WC leads to a low energy state at interfaces of complexion and WC. Based on the HRTEM

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analysis, the lattice parameter of (W,V,Cr)Cx is estimated as a=4.102Å for the formation on the WC(0001) plane and a=3.916Å for the formation on the WC(10-10) plane, respectively. The difference is attributed to the anisotropic distribution of the dissolved

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elements in the WCx lattice. As compared with the lattice parameter of WCx (a=4.226Å, measured in the WC-Co cemented carbide without addition of GGI), the decrease of the lattice parameter of the (W,V,Cr)Cx complexion with co-addition of VC and Cr3C2

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confirms the calculation results in Section 3.1.8 (Table 1).

During the cooling process of sinter-HIP, the V and Cr atoms segregate at the border of Co due to the decrease of solubility of V and Cr in the Co binder, which can facilitate the formation of new (W,V,Cr)Cx film or growth of the existed (W,V,Cr)Cx layer, as shown in Fig. 9(f). In the thickened complexion, as marked by the arrow in Fig. 9(f), the

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twinning structure is observed, which is revealed more clearly in Fig. 9(g). The twins are considered to be formed with the growth of (W,V,Cr)Cx. Assisted by twinning, part of the crystal in the (W,V,Cr)Cx complexion changes the orientation from close to

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WC(10-10) plane to WC(0001) plane. This implies that the formation and growth of the complexion are facilitated by the low WC/(W,V,Cr)Cx interface energy, which is

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influenced by the local concentrations of V, Cr and C.

Figure 9. Microstructures of the WC-Co cemented carbide with co-addition of VC and Cr3C2 and low carbon content: (a) WC grains with facets; (b) complexion viewed in WC[0001] direction, with insert of EDS analysis; (c) complexion viewed in WC[11-20] direction; (d,e) (W,V,Cr)Cx on WC(0001) and WC(10-10) planes, corresponding to (c); (f) growth of complexion, with FFT patterns of WC and (W,V,Cr)Cx; (g) enlargement of region indicated by the arrow in (f).

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Fig. 10 provides the schematic diagrams for the formation and growth of the (W,V,Cr)Cx complexion at the surface of WC grain. At the initial formation stage, there is very good atomic matching at the WC/(W,V,Cr)Cx interface. The growth of WC grain

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along WC[0001] and WC[10-10] directions can be stagnated due to the existence of (W,V,Cr)Cx films at the grain boundaries. With the thickening of (W,V,Cr)Cx, crystal defects (e.g. point defect) may be introduced at the WC surface, which can act as

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resource of stress concentration and nuclei of microcracks in the cemented carbide. Moreover, with further growth of the (W,V,Cr)Cx layer, the twinning may occur within (W,V,Cr)Cx to decrease the interface energy, as shown in Fig. 10(b). The twinning

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planes are (W,V,Cr)Cx(111) and (W,V,Cr)Cx(112). This appearance of complexion (referring to Fig. 9f) is a consequence of formation and growth on the perpendicular WC surfaces, which is likely to occur at the WC/WC/Co interjunctions. Based on the analysis of the microstructures and its formation mechanisms, it can be considered that the complexions in the cemented carbide with GGI and low carbon, on one hand, hinder

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the migration of WC grain boundaries and hence inhibit grain growth effectively; on the other hand, separate the bonding of WC and Co due to the low-energy coherent interfaces with WC grains, hence cause decrease of the resistance against intergranular

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fracture of the cemented carbide.

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Figure 10. Schematic diagrams for the formation and growth of (W,V,Cr)Cx complexion at surface of WC grain: (a) formation stage; (b) growth with occurrence of twinning.

3.2.3 Complexion in cemented carbide with high carbon In the cemented carbide with co-addition of VC and Cr3C2 and high carbon content (16.82wt%), as shown in Fig. 11, it is found that a part of complexion can transform from (W,V,Cr)Cx to WC, leading to the migration of the WC/(W,V,Cr)Cx interface towards Co binder. This indicates that the (W,V,Cr)Cx complexion is unstable in the

ACCEPTED MANUSCRIPT condition of high carbon content. As a result, with the increase of carbon content, the complexions may appear in the form of discontinuous films along WC/Co interfaces. Thus, WC and Co cannot be completely separated by the complexion, and some WC/Co interface coherency may exist in the region free of the complexion, as shown in Fig.

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11(b). In this case, the bonding strength of the WC/Co phase boundary is not affected detrimentally by the complexion due to its discontinuous distribution. Moreover, as measured from the HRTEM microstructures, the interplanar distances of (111) and (200)

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planes of fcc-Co are 2.09 Å and 1.84 Å, respectively, which are larger than the values of 2.03 Å and 1.77 Å in the cemented carbides without GGI. This implies that during the destabilization of (W,V,Cr)Cx and transformation to WC, the dissolution of V, Cr and C

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in Co results in the expansion of Co lattice. The increase of Co lattice parameter leads to a decrease in the misfit at the WC/Co interface, which is beneficial for the occurrence of interfacial coherency with good atomic matching between WC and Co.

Figure 11. Microstructures of the WC-Co cemented carbide with co-addition of VC and

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Cr3C2 and high carbon content: (a) transformation from (W,V,Cr)Cx to WC, with the arrow indicating interface migration and the EDX analysis along the solid white line; (b)

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coherent WC/Co interface in the region free of (W,V,Cr)Cx complexion.

The schematic diagrams in Fig. 12 demonstrate the transformation from

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(W,V,Cr)Cx to WC, as well as the resultant change in the configuration of WC/(W,V,Cr)Cx/Co and WC/Co interfaces, in the cemented carbide with co-addition of VC and Cr3C2 and high carbon content. At the initial sintering stage (Fig. 12a), the (W,V,Cr)Cx complexion forms on the surface of WC in the dissolution-precipitation process, with incoherent interface between (W,V,Cr)Cx and Co, as indicated by the solid line in Fig. 12(a). During the subsequent liquid-state sintering, the high carbon content in the liquid Co facilitates the transformation of (W,V,Cr)Cx to WC at the (W,V,Cr)Cx/Co interface. Thus the former incoherent (W,V,Cr)Cx/Co interface changes

ACCEPTED MANUSCRIPT to coherent WC/Co interface, as indicated by the solid line in Fig. 12(b). The local composition is critical to the formation and transformation of complexion. It was proposed in the literature that the grain/phase boundaries may undergo complexion transitions independently of the bulk phase transformations, and thus the

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interface-related properties may change unexpectedly and unpredictably [18,39,40].

Figure 12. Schematic diagrams for the transformation of complexion and change of

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interfaces in the cemented carbide with co-doping of V and Cr and high carbon content: (a) WC/(W,V,Cr)Cx/Co interfaces; (b) WC/Co bonding after transformation of

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(W,V,Cr)Cx to WC.

3.2.4 Influence of complexion on distribution of WC/Co coherency In the experiments we found that as compared with the specimen with GGI and low carbon, there is a larger probability to observe the coherent or semi-coherent WC/Co interfaces in the specimen with GGI and high carbon content. As reported in recent

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publications [6,25,41,42], a relatively higher fraction of WC/Co interfacial coherency is favorable for increase of both the fracture toughness and strength of the cemented carbides. Accordingly, the Influence of complexion on the coherency of WC/Co

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interfaces is analyzed here.

Fig. 13 shows the possible coherent and semi-coherent interfaces of WC with

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fcc-Co and hcp-Co. Figs. 13(a, b) show some examples of crystal orientations at typical coherent (indicated by the green line in Fig. 13a) and semi-coherent (indicated by the yellow lines in Fig. 13b) WC/Co interfaces, and a variety of crystallographic orientation relationship of WC/Co interface coherency are given in Fig. 13(d). The fractions of the WC/Co interface coherency, including coherent and semi-coherent interfaces which were defined as misfit < 5% and 5%~25% [43], respectively, were counted based on the EBSD measurements for the specimens with low and high carbon contents, as shown in Fig. 13(c). It’s seen that in the specimen with high carbon content, the fraction of

ACCEPTED MANUSCRIPT WC/Co interface coherency is obviously higher than that with low carbon content. The statistics is consistent with the analyses of microstructures with different carbon contents (see Fig. 9 and Fig. 11), where the effect of complexions on separating WC and Co hence reducing the probability of WC/Co interface coherency is demonstrated. In

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the cemented carbide with co-addition of VC and Cr3C2 and low carbon content, due to the stable existence of complexions with low-energy coherent (W,V,Cr)Cx/WC interfaces, the bonding of WC and Co is interfered at the phase boundary. In contrast, in

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the condition of high carbon content, higher fraction of WC/Co interface coherency can be obtained due to the discontinuous distribution of (W,V,Cr)Cx on the surface of WC grains, which is caused by the formation of complexion in a small amount and

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transformation to WC at the later stage.

Figure 13. WC/Co coherency at interfaces free of complexion: (a) example of crystal orientations at coherent interface between WC and fcc-Co, indicated by the green line; (b) examples of crystal orientations at semi-coherent interfaces between WC and

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hcp-Co, indicated by the yellow lines; (c) comparison of fractions of WC/Co interface coherency in specimens with low and high carbon contents; (d) orientation relationship of WC/Co interface coherency. In the statistics, the coherent and semi-coherent

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interfaces are classified by misfit <5% and in a range of 5% ~ 25%, respectively.

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From the studies above, it can be considered that high stability of complexions is

favorable to inhibit WC grain growth in the cemented carbides, however, it is disadvantageous to the mechanical properties due to weakening of the interfacial bonding between WC and Co binder. As described in Section 3.2, the formation and stabilization of complexions are sensitive to the carbon content. In this case, optimal design of the carbon content is an effective access to adjust complexions in the cemented carbides. An increase of carbon content can cause destabilization of complexions and thus enhance WC/Co interface bonding. Consequently, excellent

ACCEPTED MANUSCRIPT comprehensive mechanical properties can be achieved in the cemented carbide (Fig. 8a) by suitable combination of GGI and carbon content. Generally, the ultrafine WC grain structure can be obtained by use of VC and Cr3C2 inhibitors, and the transgranular fracture strength of the cemented carbide can be increased. Further, the composition

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design with appropriate carbon content is indispensable to reduce the negative effect of complexions generated by addition of VC and Cr3C2 on the WC/Co bonding, thus to

4.

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improve the toughness and intergranular fracture strength of the cemented carbide.

Discussions

In Section 3, the effects of a variety of GGI on the formation and stabilization of

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complexions, as well as the mechanical properties of the cemented carbides, were studied. Among these inhibitors, VC and Cr3C2 are the most popular in applications in WC-Co cemented carbides. Other refractory metal carbides, such as TiC, TaC and NbC, are also used in the cemented carbides, especially in the condition to improve the mechanical properties at high temperatures. In this section, the effects of various GGI

will be discussed.

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on the WC/Co interface coherency and mechanical properties of the cemented carbides

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4.1 Influence of complexion on crack propagation The influence of complexion on the crack-microstructure interactions was

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examined, in a way of detecting cracks emanating from the corners of the indentation. In the experiments, a film specimen was prepared using focused ion beam (FIB), which contains the indented cracks from the bulk material after indention test, as shown in Fig. 14(a). The specimens were sampled from the cemented carbides with co-addition of VC and Cr3C2 and different carbon contents. It is observed from the magnified microstructure in Fig. 14(b) that the crack path is along the interfaces and within Co, as marked by the dashed lines. In Fig. 14(c), as an enlargement of the region indicated by the arrow in Fig. 14(b), the crack is observed to propagate along the interface where the

ACCEPTED MANUSCRIPT complexion exists. This implies that the resistance against the intergranular fracture is decreased due to the separation of the bonding between WC and Co by the complexion. In the region indicated by the square in Fig. 14(b), the crack path is within the Co binder. The corresponding enlargement is shown in Fig. 14(d) and (e), it is observed that in this

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region WC and Co are not separated due to the discontinuous distribution of the complexion. In this case, the crack propagated in the Co binder instead of along the interface.

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To quantify, the distribution proportions of the cracks in Co binder and along the interfaces with complexions and free of complexion, respectively, were counted for the specimens with low and high carbon contents, and the results are shown in Fig. 15. The

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statistic data indicate that in the specimen with low carbon content (corresponding to a lower fraction of the WC/Co interface coherency (Fig. 13c) and a higher probability of the complexions existing at the interfaces (Fig. 9)), there is a larger proportion of the cracks along the interfaces (∼42%). In contrast, in the specimen with high carbon content, which has a larger fraction of the WC/Co interface coherency and a lower

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probability of the complexions at the interfaces, the proportion of the cracks along the interfaces is much smaller (∼28%). The results confirm that the complexions decrease

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the resistance against the intergranular fracture of the cemented carbide.

Figure 14. Influence of complexion on crack propagation in the cemented carbide with

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co-addition of VC and Cr3C2: (a) a specimen containing cracks sampled using FIB from the bulk material after indention test; (b) microstructure of the specimen with cracks; (c) enlargement of the region marked by the arrow in (b), showing crack propagation along the interface with the existence of complexion; (d) enlargement of the region marked by the square in (b), showing crack propagation in Co with the interface free of complexion; (e) enlargement of the WC/Co interface indicated by the dashed line in (d).

Figure 15. The statistic distribution of the cracks in Co binder and along the interfaces with complexions and free of complexion, respectively, in the cemented carbides with

ACCEPTED MANUSCRIPT co-addition of VC and Cr3C2 and different carbon contents.

4.2 Influence of GGI on WC/Co interface coherency From the results in Figs. 2, 3 and 5-7, it can be seen that the GGI addition can affect

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the morphology of WC grains as well as the formation of complexions. The reasons are considered to lie in the formation energy and distributing orientation of the complexion and the misfit between the complexion and WC. Due to the existence of complexions,

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the bonding of WC and Co at the interfaces is influenced. Moreover, the orientation relationship of the crystals on both sides of the complexion can be changed [20]. Here the fractions of the WC/Co interface coherency in the cemented carbides with addition

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of various GGI, i.e. VC, Cr3C2, TiC, TaC and NbC, were counted and compared, as the results shown in Fig. 16. For comparison, in these specimens the carbon content was kept constant as 16.75wt.% (an intermediate value). It is seen that the probability of the WC/Co interface coherency is decreased due to the addition of GGI. Among the inhibitors studied in the present work, the Cr3C2 has the weakest influence on

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decreasing the WC/Co interface coherency in the cemented carbide. As can be proposed, the more stably the complexion exists, the more strongly the WC grain growth is inhibited, and the lower the probability of the WC/Co interface coherency becomes in

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the cemented carbide.

The mechanical properties of the cemented carbides prepared by sinter-HIP without

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GGI and with various GGIs were measured and compared, as shown in Fig. 17. It is seen that the specimen without GGI has the highest fracture toughness, this corresponds to the largest grain size and fraction of the WC/Co interface coherency of the specimen (see Fig. 16), which implies that the combined effects of grain size and interface coherency are significant to the fracture toughness of the cemented carbide. The specimen with the addition of Cr3C2 has the highest fracture strength, which is a comprehensive outcome of the high hardness (with smaller grain size) and large fraction of the WC/Co interface coherency (Fig. 16). For comparison, the specimen with VC has

ACCEPTED MANUSCRIPT the smallest grain size, highest hardness, lowest fracture toughness and strength, which is consistent with the high stability of the complexions on both basal and prismatic planes in this specimen (Fig. 2) and the lowest fraction of the WC/Co interface

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coherency (Fig. 16).

Figure 16. Comparison of fractions of the WC/Co interface coherency in the cemented

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carbides with addition of various inhibitors.

Figure 17. Comparisons of the mean grain size and mechanical properties of the cemented carbides prepared by sinter-HIP without GGI and with GGI of VC, Cr3C2, TiC,

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TaC and NbC, respectively: (a) mean grain size; (b) hardness; (c) fracture toughness; (d) fracture strength.

From the discussions above, as far as the role of the complexions in the WC-Co cemented carbides is concerned, the positive aspect is that the existence of complexions

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can effectively inhibit growth of the WC grains, leading to an ultrafine or even nanocrystalline grain structure of the cemented carbides. On the other side, the negative aspect of the complexions lies in its layer-like distribution on the surfaces of WC grains

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due to the low-energy interfacial state. Consequently, the bonding of WC and Co is separated by the complexions, and the WC/Co combination and particularly the

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interface coherency are interfered, which causes the decrease in the strength of the cemented carbide due to the weakening of the resistance against intergranular fracture.

5. Conclusions

In the present work, based on a series of experiments and related calculations, the formation behavior and stabilization mechanisms of the complexions in the WC-Co cemented carbides were demonstrated. The effects of GGIs on the occurrence of complexions as well as on their stabilities were examined in detail, and the relationship

ACCEPTED MANUSCRIPT of the complexions with microstructures and mechanical properties of the cemented carbides were disclosed. The main conclusions are summarized as follows. (1) The WCx complexion with partially ordered cubic structure can be stabilized on the WC surface by quenching the WC-Co cemented carbide without any addition. By

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contrast, the complexions can form and exist stably at the WC/Co interfaces by addition of GGI such as VC, TiC, NbC and VC+Cr3C2 in the process of conventional liquid-state sintering. The complexions mainly have coherent or semi-coherent relationship with

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WC on the (0001) basal and (10-10) prismatic planes.

(2) The formation capability of complexions was evaluated by the first principles calculation. It shows that the formation energy and lattice parameter of the complexions

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are reduced with the doping elements of V, Ti, Nb and V+Cr, but increased with doping Cr and Ta. The (W,V,Cr)Cx complexion has the highest stability due to the lowest formation energy and the smallest misfit with WC on the basal and prismatic planes. This agrees well with the experimental finding that the complexions grow continuously in the cemented carbide with the co-addition of VC and Cr3C2.

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(3) The interfacial structure of the cemented carbides is influenced by the stability and distributing orientation of the complexions. The fraction of the WC/Co interface coherency decreases with the addition of GGI from which the complexions are caused.

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The characterization indicates that the more stably the complexion exists, the more strongly the WC grain growth is inhibited, and the lower the probability of the WC/Co

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interface coherency is in the cemented carbide. (4) The resistance against the intergranular fracture of the cemented carbides is

decreased due to the existence of complexions, which separate the bonding of WC and Co. However, the stability of complexions can be tailored by selecting suitable GGI and optimizing the carbon content correspondingly. As a result, the comprehensive properties of high fracture toughness and high strength can be obtained in the cemented carbides by controlling the complexions based on balanced optimization of inhibiting grain growth and holding WC/Co interface bonding.

ACCEPTED MANUSCRIPT Acknowledgements This work was supported by the Key Program of National Natural Science Foundation of China (51631002) and the National Science Fund for Distinguished

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41. C.B. Wei, X.Y. Song, J. Fu, X.M. Liu, H.B. Wang, Y. Gao, Y. Wang Simultaneously high fracture toughness and transverse rupture strength in ultrafine cemented carbide, CrystEngComm 15 (2013) 3305–3307. 42. H. Xie, X.Y. Song, F.X. Yin, Y.G. Zhang, Effect of WC/Co coherency phase boundaries on fracture toughness of the nanocrystalline cemented carbides, Sci. Rep. 6 (2016) 31047. 43. H. Luth, Solid surfaces, interfaces and thin films. fifth, in: W.T. Rhodes, H.E. Stanley, R.

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AC C

EP

TE D

M AN U

SC

RI PT

Needs (Eds.) Graduate Texts in Physics, Springer, Berlin, Heidelberg, 2010, pp. 80-81.

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Figure 1 

b

WC

Co

Co

RI PT

    a                                  200nm   c                                1nm    

WC

WC

Complexion

WC

SC

WC

M AN U

Co

Structural disorder

WC Co

WC

1nm

d

EP

TE D

Co

71.6°

Co

55.6°

AC C

2.328Å

WC

0111

2.402Å

1101

Co-hcp 0001

2.384Å

1010

WC-hcp

1nm

WC

ACCEPTED MANUSCRIPT

a

b

Co

Co e Complexion

M AN U

WC

d

2.512Å

2.362Å

EP

1nm

1nm

V [1010]WC

W

V

2 layers

AC C

70.2°

Co

TE D

5 layers

e

W

20nm

c

5nm

[0001]WC

SC

2.816Å

d

RI PT

Figure 2

20nm

ACCEPTED MANUSCRIPT

111

Co

b

111

1010

EP

WC-hcp

AC C

1nm

0001

TE D

WC

W

M AN U

Co-fcc

Co

SC

a

RI PT

Figure 3

Cr WC

Co

5nm

ACCEPTED MANUSCRIPT

RI PT

Figure 4

TE D EP AC C

2nm

WC

Co

M AN U

Co

b

SC

a

WC 2nm

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a

b

c

b

Co

RI PT

Figure 5

WC

WC

SC

WC

c Co WC

M AN U

200nm

2.108Å 2.552Å

(W,Ti)CX

AC C

EP

TE D

1nm

WC

1nm

ACCEPTED MANUSCRIPT

Figure 6

Co WC

Co

M AN U

SC

WC

(W,Ta)C

200nm

Co

c

Ta C 10nm

(W,Ta)C

Co

TE D

W

EP

b

WC

AC C

                                                   

RI PT

a

2.308Å

1.826Å

1nm

Co

ACCEPTED MANUSCRIPT

a

WC

Co

M AN U

WC

SC

WC

RI PT

Figure 7

WC Co

200nm

c

TE D

b

AC C

WC

EP

Co

1nm

Co WC

1nm

ACCEPTED MANUSCRIPT

a

       

 

0.36 0.30 0.24

13

1800

4000

12

1750 1700

3200

1650

Mean grain size Hardness TRS Fracture toughness

1600 1550

16.71

16.74 16.77 16.80 16.83 Carbon content (wt.%)

16.86

9 2400

c

1μm

AC C

d

EP

TE D

b

Co

1μm

e

WC

WC

WC

Facets

Co WC 50nm

50nm

10

2800

M AN U

 

11

1/2

 

3600

TRS (MPa)

 

0.42

4400  

SC

 

Mean grain size (μm)

 

1850

0.48

 

Fracture toughness (MPam )

 

Hardness (kg·mm-2)

 

RI PT

Figure 8

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Figure 9

a

b Co

[0001] 300nm

2nm

c

WC

[0001]

d

M AN U

[1210]

71°

SC

At.% W 27.7 C 35.9 V 29.9 Cr 6.5

RI PT

54.5°

WC

WC

4 layers

0.5nm

d

e

TE D

Complexion

e

Co

2nm

AC C

f

EP

3~5 layers

WC 0.5nm

g

WC

WC WC

(W,V,Cr)CX 2nm

(W,V,Cr)CX (W,V,Cr)CX

1nm

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Figure 10

[1010]

[0001]WC

[1210]WC

[0001]

[1010]

[111] [110]

[110](W,V,Cr)Cx

b [1010]

[200]

[1210]WC

[111]

AC C

EP

[111]

TE D

[0001]

[200]

[111]

M AN U

[200]

[002]

SC

[1210]

RI PT

a

WC (W,V,Cr)Cx with different orientations Defect Twinning plane in (W,V,Cr)Cx

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Figure 11

a

RI PT

b [1210]

M AN U

Co C W Cr V

Co

EP

TE D

WC

AC C

2nm

SC

WC

Co

Co-fcc

2nm

WC-hcp

ACCEPTED MANUSCRIPT

a

[1210]WC

b

[1010]

SC

[0001]

M AN U

[111] [200]

[200] [111]

EP

TE D

 

AC C

RI PT

Figure 12

WC (W,V,Cr)Cx V, Cr Co-fcc

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c

b

[1210]WC

[002]

[1210]WC

[0001]

[110]

[1010]

AC C

[1010]

[110]Co

18 16

[1210]WC

14 12 10

16.82 16.71 Carbon content (wt.%)

[1210]WC

[1010]

[0001]

[0001]

[111]

[1210]WC

[1010]

TE D

[1010]

EP

d

30nm

M AN U

(1010)WC/(0001)Co

(0001)WC/(1010)Co

20

SC

(0001)WC/(002)Co 30nm

22

RI PT

a

Fraction of interface coherency (%)

Figure 13

WC Co-fcc

[0001]

Co-hcp Interface with misfit < 5%

[111]

Interface with misfit 5% ~ 25%

[111] [200]

[112]Co [1210]

[011]Co

[0001]WC

[0001]

[1010]

[0001]

[1210]WC

[1010]

[0001]

[0001]

[1010]

[1010] [1010]

[1101]

[1010] [1210]Co

[1213]Co

[1101] [1210]Co

[1213]Co

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a

c WC

   

WC

   

Indented crack

 

1nm 2μm

   

b WC

   

WC

WC

Co

     

AC C

100nm

EP

WC

 

 

d

WC

1nm

TE D

 

 

e

Co

 

 

Crack

M AN U

 

SC

 

 

RI PT

Figure 14

Crack tip

Co 2nm

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Figure 15

 

RI PT

Total crack length (μm) Proportion of cracks in Co binder (%) Proportion of cracks along interfaces (%)

 

 

375

45

 

350

40

325

35

 

300

30

 

275

25

 

250

20

20

 

225

15

15

10

10

   

200

16.71 16.82

 

16.71 16.82 16.71 16.82 Carbon content (wt.%)

AC C

EP

TE D

   

SC

 

40

M AN U

 

45 35 30 25

ACCEPTED MANUSCRIPT

Without GGI

20

Cr3C2

NbC

VC

10

5

M AN U

TiC

SC

TaC

15

Addition of GGI

EP AC C

RI PT

25

TE D

Fraction of interface coherency (%)

Figure 16

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Fracture toughness (MPam1/2)

 

0.2

c

Without GGI

18 16

TaC Cr3C2

14

TiC

12 VC

10 8

1700

1500 Without GGI

1400 1300

4000

NbC

3500

VC

2000 1500

TiC

TaC

NbC

d

Cr3C2

Without GGI

3000 2500

b

Cr3C2

1600

1200

20

VC

SC

VC 0.4

0.0

   

NbC

TRS (MPa)

 

0.6

TaC

1800

M AN U

 

TiC Cr3C2

TE D

 

0.8

EP

 

1900

a

Without GGI

AC C

 

1.0

Mean grain size (μm)

 

Hardness (kg·mm-2)

 

RI PT

Figure 17

NbC TiC

TaC