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Microstructure and mechanical properties of WC–Ni multiphase ceramic materials with NiCl2·6H2O as a binder

  • Xingxing Lyu , Xiaosong Jiang EMAIL logo , Hongliang Sun and Zhenyi Shao
Published/Copyright: June 13, 2020
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Nanotechnology Reviews
From the journal Nanotechnology Reviews Volume 9 Issue 1

Abstract

High-density WC–Ni composite ceramics were prepared by cold isostatic pressing–vacuum pressureless sintering–hot isostatic pressing with tungsten carbide (WC) powder and NiCl2·6H2O as a binder. Results show that with an increase in the contents of Ni in the metal binder phase, the relative density of WC–Ni composite ceramics is improved, and the formation of the carbon-deficient W2C phase is reduced. There is no W2C generated in the WC–1 wt% Ni material. At high temperatures, the Ni phase changes into the liquid phase and enters between the WC particles, thereby promoting the close alignment of the WC particles. Moreover, the WC particles will be more closely aligned under their own surface tension and capillary action, thereby promoting the densification of WC–Ni composite ceramics. The WC–0.5 wt% Ni composite ceramics are fully dense and show the best comprehensive performance with a microhardness of 23.0 GPa, a fracture toughness of 5.28 MPa m1/2, and a flexural strength of 1,396.58 MPa. WC–Ni composite ceramics are mainly composed of elongated triangular prism WC particles and Ni phase. Transgranular fracture was the main fracture mode of WC–Ni multiphase ceramic materials with a small amount of intergranular fracture due to the existence of the Ni phase. Such a fracture mode can increase the flexural strength of the composite material.

Graphical abstract

Keywords: WC–Ni composite ceramics; hot isostatic pressing; densification

1 Introduction

Cemented carbides are obtained by the powder metallurgy preparation process, using a hard compound of refractory metal as the matrix and adding a transition metal (such as cobalt, nickel, and iron) as a sintering aid. WC-based cemented carbide is the most widely used cemented carbide [1,2,3]. Due to its excellent combination of high strength, high hardness, and high wear resistance, it is widely used in cutting tools, mining, wear parts, drilling, metal cutting, molds, and other fields [1,2,3,4,5,6,7,8]. In particular, it has excellent wear resistance and is often used in some areas with high wear resistance [9,10]. However, WC-based cemented carbides have low toughness and sinterability [11]. This is because the WC material has a hexagonal crystal structure, and carbon atoms exist in the gap of the tungsten metal lattice to form gap solid solution. The crystal structure has a strong bonding force between the WC particles, resulting in high hardness, low toughness, and low sinterability. Therefore, it is difficult to sinter densified WC materials by conventional sintering methods (such as hot press sintering and vacuum pressureless sintering (VPS)) [12].

To overcome the low sinterability and high brittleness of the WC materials, sintering aid such as Co or Fe is added into the WC particles to improve the density and toughness of the WC-cemented carbide [13]. Co is the most widely used metal binder in WC-based cemented carbides [14]. Because metal Co has excellent wettability to WC, it can reduce the sintering temperature of WC ceramics and can improve the sintering density and toughness of the material [15,16]. However, the addition of Co could decrease the hardness, wear resistance, and oxidation resistance of the WC-cemented carbides. Thermal stress, to a large extent, limits the application of WC–Co-cemented carbides in the industrial field, which is caused by the different thermal expansion coefficients of Co and WC substrates [17,18]. In addition, Co has low reserves and costs high so that it cannot be used in a wide range [19,20,21]. Therefore, a growing number of researchers are striving to find a new binder to substitute Co. Studies have shown that metal Ni can be used as a binder in the sintering of cemented carbides to improve its corrosion resistance and oxidation resistance [19,20].

Results have shown that WC–Ni-cemented carbide has higher corrosion resistance and oxidation resistance than WC–Co-cemented carbide [8,22,23]. As a binder of WC-based cemented carbide materials, Ni can reduce the sintering difficulty of WC-based cemented carbide materials effectively and enables to lower the metal binder content [24,25]. This is because during the sintering process, Ni can form a liquid phase easily, which provides better wettability of WC surfaces. With the generation of the liquid-phase Ni adsorption process, the surrounding WC particles are also easily aggregated, thereby accelerating the densification of WC-based cemented carbides [26]. In addition, under vacuum sintering conditions, the liquid-phase Ni in WC–Ni cemented carbide increases as the temperature increases. An increase in the liquid-phase Ni accelerates the movement of WC particles and, on the other hand, promotes its homogeneous distribution around WC particles, expediting the densification process. Finally, an almost completely densified WC–Ni cemented carbide material is obtained [3]. In the actual industrial production, Ni is lower in costs than Co and the storage amount of Ni is much larger; hence, it is a better choice to use Ni instead of Co as the bonding phase of the cemented carbide [27].

However, the commonly used nickel binders for preparing WC–Ni multiphase ceramics are nickel powder and nickel oxide powder. After the addition of nickel oxide powder, the appearance of nickel pool is easy after the mechanical alloying treatment [28,29]. Therefore, this article hopes to find a source of Ni that can be dissolved in water or alcohol solvents and then attempts to use NiCl2·6H2O as the source of Ni. This is because NiCl2·6H2O has the advantage of being easily soluble in water; hence, it is not easy to agglomerate in the ball mill to avoid the nickel pool phenomenon. Therefore, in this article, NiCl2·6H2O was used as the sintering aid, and after ball milling with WC powder, WC–Ni composite ceramics were prepared by cold isostatic pressing–vacuum pressureless sintering (VPS)–hot isostatic pressing (HIP) densification technology. The three different Ni addition amounts (0.25, 0.5, and 1 wt%) of WC–Ni composite ceramics were vacuum sintered at 1,530 and 1,560°C, then subjected to HIP. Microstructure properties were characterized and analyzed. And the strengthening and toughening mechanisms of experimental WC–Ni composites were discussed.

2 Experimental procedure

2.1 Material preparation

In this experiment, ultrafine WC powder (Xiamen Jinlu Special Alloy Co., Ltd, 0.53 µm, containing a small amount of VC and Cr3C2) and NiCl2·6H2O powder (Chengdu Kelong Chemical Reagent Factory, Analytical Grade) were used as raw materials. NiCl2·6H2O was used as the Ni source to substitute Co as a metal binder, and the composition of the design specimens having different Ni contents is given in Table 1.

Table 1

WC–Ni composite ceramics raw material powder content

Numbering Ingredient WC (wt%) Ni (wt%) (NiCl2·6H2O)
WCN-1# WC–0.25 wt% Ni 99.75 0.25
WCN-2# 99.75 0.25
WCN-3# WC–0.5 wt% Ni 99.50 0.50
WCN-4# 99.50 0.50
WCN-5# WC–1 wt% Ni 99.00 1.00
WCN-6# 99.00 1.00

The main experimental steps are shown in Figure 1. The raw material powder was placed in a planetary ball mill (DECO-PBM-AD-2; Changsha Deco Instrument Equipment Co., Ltd) for 2 h to make the mixture uniform. The ball-milled raw material was then dried in a freeze dryer (FD-1B-50; Beijing Bo Medical Laboratory Instrument Co., Ltd) for 24 h to obtain a dried mixed powder. The mixed powder was pressed into a cylindrical piece with a diameter of 30 mm, and a total of six specimens were designated as WCN-1#, WCN-2#, WCN-3#, WCN-4#, WCN-5#, and WCN-6#. The press-formed sheet specimen was placed in a plastic wrap to vacuum and then placed in a cold isostatic pressing tester (LDJ600-200-300; Chuanxi Machinery Factory) for 10 min at the pressure of 200 MPa. The cold isostatically pressed specimen was sintered in a vacuum hot press sintering furnace (VTP-1; Shenyang Weitai technology) according to the process given in Table 2. After that, the specimens were placed in an HIP furnace (QIH-15, ABB, USA) and sintered without encapsulation in an Ar gas medium; the temperature, pressure, and time were set according to the data given in Table 2.

Figure 1 Flow chart of the experimental specimen preparation.
Figure 1

Flow chart of the experimental specimen preparation.

Table 2

WC–Ni composite ceramics sintering process

Numbering Ingredient Sintering process HIP
WCN-1# WC–0.25 wt% Ni VPS 1,530°C × 1.5 h 1,600°C × 80 MPa × 1 h
WCN-2# VPS 1,560°C × 1.5 h 1,600°C × 80 MPa × 1 h
WCN-3# WC–0.5 wt% Ni VPS 1,530°C × 1.5 h 1,600°C × 80 MPa × 1 h
WCN-4# VPS 1,560°C × 1.5 h 1,600°C × 80 MPa × 1 h
WCN-5# WC–1 wt% Ni VPS 1,530°C × 1.5 h 1,600°C × 80 MPa × 1 h
WCN-6# VPS 1,560°C × 1.5 h 1,600°C × 80 MPa × 1 h

2.2 Material characterization

The densities of WC–Ni composite ceramic materials were measured by the Archimedes method, and the phase composition of the composites was analyzed using X-ray diffraction (XRD; Philips PM1700). The hardness tester (HV-50 Vickers) was used to test its Vickers hardness, the load was 10 kg, and the dwell time was 15 s; each specimen is measured 5 points to average. The indentation method was used to measure the fracture toughness of the composite, and the flexural strength of the composite was characterized by the three-point bending test (WDW-3100 Microcomputer Controlled Electronic Universal Material Testing Machine). The microstructure, fracture morphology, and interface bonding of the composites were observed by scanning electron microscopy (SEM; FEI Inspect F50) and transmission electron microscopy (TEM; FEI Titan G2 60-300). The compositions and elemental distributions of the composites were analyzed using an energy dispersive X-ray spectrometry (EDS; FEI Inspect F50).

3 Experimental results

3.1 Composition and microstructure of mixed powders

Figure 2(a) shows the XRD pattern of the WC powder before ball milling and the XRD pattern of the mixed powder after adding NiCl2·6H2O and ball milling with WC powder for 24 h. It is found that after 24 h of ball milling, the diffraction peak of WC weakened and cracked, indicating a decrease in the crystal size [30,31,32]. No diffraction peak of Ni was detected because of its low content. Figure 2(b) and (c) shows the SEM images of the WC powder and the WC–1 wt% Ni powder, respectively, and the particles show a polyhedral shape. The average particle size of the WC powder is 0.53 µm, which was obtained by the line intercept method, and the average particle size of WC–1 wt% Ni powder is 0.39 µm. It can be seen that the particle size of the powder after ball milling is reduced. The ball mill makes the powder more uniform, which laid a good foundation for postsintering densification.

Figure 2 (a) XRD patterns of the WC–1 wt% Ni composite powders ball milled at different times, (b) SEM images of the WC power, (c) SEM images of the WC–1 wt% Ni powder, and (d) EDS images of WC–1 wt% Ni powder.
Figure 2

(a) XRD patterns of the WC–1 wt% Ni composite powders ball milled at different times, (b) SEM images of the WC power, (c) SEM images of the WC–1 wt% Ni powder, and (d) EDS images of WC–1 wt% Ni powder.

3.2 Phase composition and microstructure of WC–Ni composite ceramics

Figure 3 shows the XRD patterns of six specimens at different compositions after different sintering processes. According to the XRD pattern analysis, it can be observed that the main diffraction peaks belong to the WC phase, whereas all specimens sintered by vacuum contain a small amount of W2C. In Figure 3, the Ni content in each of the two groups increased from top to bottom sequentially. Conversely, the intensity of the diffraction peak of the carbon-deficient phase W2C gradually decreased and almost disappeared. The occurrence of W2C was due to its great relationship with the sintering temperature and holding time; the carbon-deficient phase WC1−x is easier to generate at high temperatures and long holding time. In addition, as a small amount of oxide layer or oxygen was adsorbed on the surface of the raw powders, it would react with graphite at high temperatures and long holding time to make the stoichiometric ratio of W to C greater than 1; thus, W2C phase was formed during the sintering process. In Figure 3(e) and (f), when the Ni content is 1 wt%, only the WC phase diffraction peak is present in the specimen and the carbon-deficient phase W2C substantially disappears. In this work, since the WC–Ni multiphase ceramics were studied with less binder phase, the Ni content in the specimen material is at a quite low level; hence, no Ni phase diffraction peaks were found [28]. It can be seen from the spectrum that the W2C content of the carbon-deficient phase decreases with an increase in the Ni content, and the W2C diffraction peak intensity of the carbon-deficient phase in the specimen sintered at 1,560°C is lower than that of the specimen with the same Ni content. It indicates that the more carbon-deficient the phases, the lower the density of the material [33].

Figure 3 X-ray diffraction pattern of the WC–Ni composite ceramic specimen: (a) WCN-1#, (b) WCN-2#, (c) WCN-3#, (d) WCN-4#, (e) WCN-5#, and (f) WCN-6#.
Figure 3

X-ray diffraction pattern of the WC–Ni composite ceramic specimen: (a) WCN-1#, (b) WCN-2#, (c) WCN-3#, (d) WCN-4#, (e) WCN-5#, and (f) WCN-6#.

According to the analysis of the raw material powder, the carbon content in the raw material powder satisfies the stoichiometric ratio but contains 0.29 wt% of oxygen. The presence of W2C may be due to the reduction of oxides on the surface of the powder during sintering, thereby consuming C in the powder, making it impossible to maintain the stoichiometric ratio required to form WC [34]. Besides, the raw material powder was exposed to vacuum and reducing atmosphere in the graphite mold. At a high temperature, the surface oxide may also be reduced by C in the mold or the generated CO to form W powder, so that the atomic ratio of W to C is greater than 1, thereby forming W2C. In addition, when referring to the W–C binary phase diagram [35], the phase behavior of the (W2C + WC) state is exhibited at a sintering temperature higher than 1,250°C. During the sintering process, α–W has a relatively low activation energy (Q v = 247 kJ/mol) and tends to form a W2C phase by C diffusion reaction with 2 W + C → W2C. Second, the formation of W2C may be achieved by rapid cooling of the WC, which is formed by temperature-driven, nonequilibrium transitions for kinetically advantageous products such as W2C and WC1−x [36,37]. This also explains the presence of W2C in the specimen. The W2C (content less than 3.5 wt%) in WC ceramics has little effect on hardness and fracture toughness, but when the content of W2C increases to a certain extent, its hardness and fracture toughness decrease significantly with an increase in the W2C content [33].

Figure 4 presents the backscattered topography and EDS elemental surface distribution of the specimenWCN-6#. The gray matrix in Figure 4(a) corresponds to WC, and the Ni element distribution is very uniform (Figure 4(c)–(e)), and no agglomerate of Ni is present [28]. During high-temperature sintering, the Ni phase adhesive becomes a liquid phase to make the densification easier, and the void defects that generate after high-temperature sintering can sufficiently grind the powder particles during mixing, reducing aggregation and making the specimen more compact after sintering. Since Ni belongs to the FCC crystal system and has good plasticity, it is easy to deform plastically during the wet grinding process to form a sheet-like Ni powder mass. Therefore, it is necessary to avoid the aggregation effect of the powder during wet grinding, so that the phenomenon of “nickel pool” can be eliminated. Increasing the sintering temperature or prolonging the holding time and HIP can reduce the porosity and improve the overall performance of the cemented carbide [3,26].

Figure 4 WCN-6# backscatter topography and surface scanning map, (a) backscatter topography, (b) surface scanning map, (c) Element W, (d) Element C, (e) Element Ni, and (f) EDS results.
Figure 4

WCN-6# backscatter topography and surface scanning map, (a) backscatter topography, (b) surface scanning map, (c) Element W, (d) Element C, (e) Element Ni, and (f) EDS results.

Figure 5 shows a transmission electron microscopic analysis of the specimen WCN-4#. It can be seen from Figure 5(a) that the WC grain morphology is mainly lath and triangular, which is because the WC monomorphism tends to be in triangular prism shape, so the crystal grains appear as triangles and slats [38]. The intrinsic factor that determines the grain morphology is determined by the point group. It can be seen that the space group of WC is P 6 ¯ m 2 , and the space group knows that the point group of WC is 6 ¯ m 2 , and the lattice constant a in a WC unit cell is greater than c [39]. Based on the relation among the point group, the initial surface, and the simple shape, when the initial surface is ( 10 1 ¯ 0 ) , the simple shape of the WC is a triangular prism, the shape of the predicted WC grain is a strip shape, and the initial surface is (0001). When the WC has a simple shape on both sides, the predicted WC grain shape is a triangle. The crystallographic plane reflects the anisotropy of the crystal, which is caused by the different surface energies of the crystal faces. The WC grains are lath-like because the surface energies of ( 01 1 ¯ 0 ) and (1000) are different [28,40,41]. Figure 5(b) shows the electron diffraction pattern of the WC grain in specimen WCN-4#, inferring that the grain only contains WC and its crystal axis direction is [010]. Figure 5(c) and (d) shows the crystal defects in specimen WCN-4#. A large number of crystal defects such as dislocation arrays and twins can be observed in the WC crystal grains. In Figure 5(c), the parallel distribution of dislocations can be seen. The observed twin morphology is shown in Figure 5(d). It can be seen from the figure that the twins are almost parallel and the interface is straight grid interface. The information of dislocation column may be due to high temperature and high pressure, the crystal grains are plastically deformed, the dislocations are proliferated, and these dislocations slide down in the force. When the obstacles are encountered, they are plugged in front of the obstacles. Delivery may occur when the wrong encounter occurs. When the dislocation moves to the grain boundary, the motion of the dislocation is blocked, and it will be concatenated near the grain boundary, thereby forming a dislocation column [28]. Due to the existence of dislocation columns, the dislocations of WC grains are difficult to slip at room temperature, stress concentration is easily formed in the crystal grains, and the stress concentration caused by each crystal grain is different, so the crystal defects are distributed uniformly. It may be one of the reasons why the dispersion of the flexural strength is large.

Figure 5 WC grain morphology and corresponding electron diffraction pattern of WCN-4# specimen and crystal defects of specimen: (a) WC grain morphology, (b) electron diffraction pattern of WC grain, (c) dislocation column, and (d) twins.
Figure 5

WC grain morphology and corresponding electron diffraction pattern of WCN-4# specimen and crystal defects of specimen: (a) WC grain morphology, (b) electron diffraction pattern of WC grain, (c) dislocation column, and (d) twins.

Figure 6 shows a transmission electron microscopic analysis of WC, Ni, WC, and Cr3C2 performed on the specimen WCN-4#. It can be seen from Figure 6(a) that the nickel phase is mainly distributed at three points in the boundary, the volume is relatively small, and the distribution is uniform, and no agglomeration occurs. Figure 6(a) shows that the nickel phase diffuses in this region, which is consistent with the characteristics of liquid phase sintering. At high temperatures, the nickel phase becomes liquid and rapidly permeates into the pores between the WC particles. The nickel phase has good wettability and will cover the WC particles well; hence, the nickel phase can be well distributed at the grain boundaries [28,42]. Figure 6(b) shows an interface diagram of WC and Ni. It can be seen that WC and Ni are well combined, and the interplanar spacing of Ni phase and WC is 0.1956 nm and 0.2535 nm, respectively. TEM (Figure 6(c)) and EDS (Figure 6(d)) were used to observe and analyze the distribution of grain growth inhibitors VC and Cr3C2. Figure 6(c) shows the typical microscopic morphology and corresponding energy spectrum at the triple grain boundary. It could be observed that the grain boundary at this point is a flat grain boundary where V and Cr are detected. Analysis shows that there are grain growth inhibitors VC and Cr3C2. Figure 6(d) shows a high-resolution diagram of WC and Cr3C2. The interplanar spacing of WC and Cr3C2 is 0.2313 and 0.4845 nm, respectively, and it is found that Cr3C2 is diffused, which better limits the growth of WC grains. It is found that the flat grain boundary is a rough interface on the atomic scale, and the facet grain boundary is a smooth interface on the atomic scale. When the grain boundary is straight, the grain grows normally. However, when the grain boundary is a facet, it will cause grain growth abnormally [12,43]. Therefore, when the grain inhibitors VC and Cr3C2 are added, VC and Cr3C2 will hinder the growth of grains under high temperature and long-term sintering conditions; hence, fine microstructures can still be obtained. Therefore, the density and fracture toughness of WC–Ni composite ceramics can be improved [42,44]. In addition, VC and Cr3C2 at the grain boundaries can reduce the surface energy of WC particles and reduce the driving force for the grain growth [45].

Figure 6 (a) TEM results of the topography, (b) TEM results between WC and Ni, and (c and d) TEM results between WC and Cr3C2.
Figure 6

(a) TEM results of the topography, (b) TEM results between WC and Ni, and (c and d) TEM results between WC and Cr3C2.

3.3 Performance analysis of WC–Ni composite ceramics

3.3.1 Density

When compared with the conventional method, the preparation method in this paper uses NiCl2·6H2O as the nickel source. This is because NiCl2·6H2O has the advantage of being easily soluble in water; hence, it is not easy to agglomerate in the ball mill to avoid the nickel pool phenomenon. The appearance of the nickel pool will greatly reduce the density and mechanical properties of WC–Ni composite ceramics. Therefore, avoiding the appearance of the nickel pool is of great significance to the density and mechanical properties of WC–Ni composite ceramics. As shown in Figure 7, the densities of the WC–Ni composite ceramic materials of different compositions and different sintering processes generally show a tendency to increase with an increase in the Ni content. It can be seen that the adhesive Ni exhibits a good promotion effect on the sintering densification of the WC–Ni composite ceramic material. That is because Ni can be easily melted into a liquid phase at a high temperature and penetrates the pores between the WC particles so that the WC particles can be well infiltrated and the WC particles are more closely arranged under the action of surface tension and capillary effect at this time to increase the density of the specimen [26]. At the same time, as the sintering temperature increases, both the fluidity of liquid Ni and the driving force of sintering increase, which accelerates the densification of ceramic materials and increases the density of ceramic materials. From the results of the specimen densification of WC–Ni composite ceramics with 0.25, 0.5, and 1 wt% Ni metal binder phases after VPS at 1,530 and 1,560°C and then HIP, the analysis shows that WCN-4#, the ceramic material of WC–0.5 wt% Ni sintered at 1,560°C, has the best densification degree, reaching 100.13%, indicating that the material is completely dense. The relative density is greater than 100% because of the reaction occurring during the sintering process, which provides a phase different from that of the original material; hence, the density theoretically obtained by the mixing rule is not the best reference for the relative density or is obtained from the density value porosity data [46]. WC–0.25 wt% Ni and WC–1 wt% Ni specimens have a density lower than WC–0.5 wt% Ni specimen. Comparing the densities of ceramic specimens at different sintering temperatures of the same composition, it can be found that the WC–Ni composite ceramic specimens sintered at 1,560°C are denser than that of the vacuum sintered WC–Ni ceramic specimens at 1,530°C. On the whole, the densities of 0.25, 0.5, and 1 wt% of the composite materials of Ni content are more than 97%, and the degree of densification is greatly improved when compared with pure WC ceramics. The main reason is that the VPS process makes the material dense, then the HIP treatment can make the material completely dense, and the density is increased to achieve secondary densification. The density of the specimen WCN-5# in the data was significantly reduced, probably due to an increase in the Ni content, which led to a longer period of liquid phase sintering, resulting in a decrease in the density. At the same time, Lisovsky [47] found that the flow driving force of the binder liquid phase between the WC particles is inversely proportional to the WC particle size. Therefore, the rearrangement of the WC particles in the liquid phase could be accelerated as the sintering temperature and the binder increase, and at the same time, the growth of the WC particles could also be promoted. Therefore, the flow driving force of the binder liquid phase between the WC particles is reduced, resulting in a decrease in the driving force of the binder liquid phase filling hole and a decrease in the density.

Figure 7 WC–Ni composite ceramics density.
Figure 7

WC–Ni composite ceramics density.

3.3.2 Hardness and fracture toughness

Table 3 presents the mechanical property results of WC–Ni composite ceramic materials with different binder contents (0.25, 0.5, and 1 wt%). It can be seen from Figure 7 that Ni has a significant influence on WC sintering to achieve high density. When the Ni content is 0.25 wt%, the hardness, fracture toughness, and flexural strength of specimen WCN-2# are 23.90 GPa, 4.10 MPa m1/2, and 935.87 MPa, respectively. As the Ni content increased to 1 wt.%, the hardness, fracture toughness, and flexural strength of specimen WCN-6# are 23.00 GPa, 4.30 MPa m1/2, and 1514.77 MPa, respectively. This shows that an increase in the Ni content greatly improves the flexural strength of WC–Ni composite ceramics [28]. At the same time, WC–Ni composite ceramic materials have a high hardness of more than 21.0 GPa, which is significantly higher than the WC–15 wt% Co composite (less than 17 GPa) prepared by Shih-Hsien Chang [48] and is close to fully dense (approximately 24 GPa) pure WC prepared by other researchers [18,49]. When the sintering temperature is 1,560°C, the hardness of the WC–Ni composite ceramics specimen decreases with an increase in the Ni content, which is due to the increase in the Ni content, which leads to a longer period of liquid phase sintering, resulting in a decrease in the density. At the same time, the flow driving force of the binder liquid phase between the WC particles is inversely proportional to the WC particle size. Therefore, with an increase in the sintering temperature and the binder, the rearrangement of WC particles in the liquid phase is accelerated, and the growth of WC particles is promoted. Due to the growth of WC particles, this hinders the flow of the binder liquid phase and reduces the driving force for the binder liquid phase to fill the pores, resulting in the presence of pores between the WC particles, which can reduce the density and hardness [47].

Table 3

Performance of WC–Ni composite ceramics

Numbering Ingredient (wt%) Hardness (GPa) Fracture toughness K IC (MPa m1/2) Bend-resistant strength (MPa)
WCN-1# WC–0.25 wt% Ni 22.50(±0.22) 5.76(±0.37) 1168.29(±246.28)
WCN-2# 23.90(±0.22) 4.10(±0.21) 935.87(±272.03)
WCN-3# WC–0.5 wt% Ni 21.83(±0.05) 4.14(±0.22) 1225.06(±139.24)
WCN-4# 23.00(±0.22) 5.28(±0.25) 1396.58(±50.17)
WCN-5# WC–1 wt% Ni 23.73(±0.12) 4.38(±0.28) 1181.17(±105.51)
WCN-6# 23.00(±0.12) 4.30(±0.09) 1514.77(±259.6)

In this work, the fracture toughness test of the indentation method [50] was used. The K IC calculation formula is calculated by the indentation method:

(1) K IC = 0.075 P ⋅ C − 3 / 2 ,

where K IC is the fracture toughness (kgf m−3/2), P is the load size (kgf), and C is the diagonal length of indentation crack (mm), as 1 kgf m−3/2 = 0.31 MPa m1/2, hence it can be converted to

(2) K IC = 0.02325 P ⋅ C − 3 / 2 ,

where K IC is the fracture toughness (MPa ml/2), P is the load size (kgf), and C is the diagonal length (mm) of indentation crack.

Figure 8 compares the hardness and fracture toughness of each group of the WC–Ni composites. Observing and analyzing the fracture toughness data and the graph, it is known that the fracture toughness of WC–Ni composite ceramics specimens decreases first and then increases with an increase in Ni content at a sintering temperature of 1,530°C, but in the WC–Ni complex phase at a sintering temperature of 1,560°C, the fracture toughness of ceramic specimens increases first and then decreases with an increase in the Ni content; hence, the Ni content and sintering temperature will affect the fracture toughness at the same time. Moreover, the W2C existing in WC ceramics has an effect on the hardness and fracture toughness. It is found that the hardness and fracture toughness of WC ceramics decrease with an increase in the W2C content [33]. According to the XRD phase analysis, it can be known that an increase in the Ni content makes the content of carbon-deficient phase W2C gradually decrease or even disappear. The decrease of carbon-deficient phase W2C is beneficial to improve the density and performance of the material; hence, the fracture toughness of the material gradually increases. When the Ni content increases from 0.25 to 0.5 wt%, the fracture toughness increases [33]. When the content of Ni in the metal binder phase is 0.25 wt% sintering at 1,560°C, the WC–Ni composite has the highest hardness of 23.9 GPa and fracture toughness of 4.10 MPa ml/2. The fracture toughness of the WC–Ni composite prepared by sintering at 1,530°C is as high as 5.76 MPa ml/2, and the hardness of the material is also high (22.5 GPa).

Figure 8 Hardness and fracture toughness of each specimen of WC–Ni composite ceramics.
Figure 8

Hardness and fracture toughness of each specimen of WC–Ni composite ceramics.

3.3.3 Flexural strength

Figure 9 shows the flexural strength of each group of the WC–Ni composites. The flexural strength of the specimen generally increases with an increase in the Ni content. Since the uniformity and continuity of the microstructure of the material will affect the flexural strength of the material, the addition of an appropriate amount of Ni makes the WC grain size more uniform, so that the flexural strength of the WC–Ni composite ceramics is improved while it is easy to cause aggregation due to the excess Ni. Specimen WCN-2# has the lowest flexural strength value of 935.87 ± 272.03 MPa even though the sintered WC–Ni composite ceramic at 1,560°C has a lower density. In contrast, specimen WCN-6# with the WC–Ni composite ceramics sintered at 1,560°C has a flexural strength of 1514.77 ± 259.6 MPa, which is due to better liquid phase sintering results [14]. As can be seen from Figure 7, when the sintering temperature is from 1,530 to 1,560°C, the relative density of the WC–Ni multiphase ceramics exceeds 99%. In general, the internal pores of the material are prone to stress concentration. However, the WC–Ni composite ceramics sintered at 1,530 and 1,560°C achieve almost complete densification, which effectively reduces the strain point along with the fracture mechanism and increases the resistance bend strength value. At the same time, the improvement of the Ni phase contributes to the densification of the samples, and sintering at high temperatures generates a large amount of liquid phase filling the pores [51,52], which also greatly improves the flexural strength of the WC–Ni composite ceramics.

Figure 9 Flexural strength of each specimen of WC–Ni composite ceramics.
Figure 9

Flexural strength of each specimen of WC–Ni composite ceramics.

4 Discussion

4.1 Densification mechanism

For WC-based cemented carbides, WC grain growth is divided into two stages during the sintering densification process. The first is a solid phase sintering stage in which small WC particles are grown by surface diffusion to contact with each other. The second is a liquid phase sintering stage in which small WC particles are premelted due to the presence of a binder-rich liquid phase and then precipitated on the surface of the coarse WC particles by migration, resulting in WC particle size growth. When the binder is added to the alloy, the contact area between the WC particles is reduced, thereby suppressing the interdiffusion of WC, resulting in limited growth of the WC particles. Furthermore, W dissolves in the binder phase during sintering, thereby suppressing W precipitation and WC grain growth. Consequently, the WC grains in the alloy become more uniform and finer.

Figure 10 shows a broken point sweep diagram of specimen WCN-6#. The binderless WC is difficult to reach a fully dense state due to the lack of a metal binder phase. In the present study, WC–Ni composite ceramics were obtained by vacuum pressureless sintering (VPS) followed by HIP sintering. Ni is added as a binder phase because its liquid exhibits good fluidity and wettability at high temperatures. As can be seen from Figure 10(a) and (b), the binder phase Ni distributed at the triple point of the grain boundary shows the characteristics of liquid phase sintering. The Ni phase melts into a liquid phase at a high temperature and penetrates the pores between the WC particles. It also penetrates the WC particles and is more closely aligned under the influence of the surface tension and capillary action of the WC particles, thereby promoting the densification of the sintering process [26]. The fluidity of liquid Ni increases with an increase in the temperature, and the driving force also increases, which accelerates the movement process of WC particles and thus increases the density of materials [3]. Moreover, the closed cells remaining after the VPS sintering can be eliminated by HIP, and a completely dense WC–Ni composite material is obtained. Simultaneous addition of grain growth inhibitors maintains the uniformity and fineness of the grains and also increases the density of the material.

Figure 10 WCN-6# specimen fracture scanning and EDS results, (a) specimen fracture scanning, and (b) EDS results.
Figure 10

WCN-6# specimen fracture scanning and EDS results, (a) specimen fracture scanning, and (b) EDS results.

4.2 Strengthening and toughening mechanisms

Figure 11 shows the line-cut method to calculate that the grain sizes of specimen WCN-5# and specimen WCN-6# are 0.64 and 0.74 µm, respectively, so that the specimens can be seen after VPS and HIP, and the grain size of WC did not grow significantly. During liquid phase sintering, small WC particles are premelted due to the presence of a binder-rich liquid phase between the WC particles and then precipitated on the surface of the coarse WC particles by migration, resulting in WC particle size growth. However, when the binder particles are added to the alloy, the contact area between the WC particles is reduced, thereby suppressing the interdiffusion of WC, resulting in limited growth of the WC particles. In addition, a small amount of grain inhibitors (VC and Cr3C2) are contained in the WC powder. It is found that the addition of VC and Cr3C2 grain inhibitors can effectively inhibit the growth of WC particles, reduce the size of WC particles, and control the abnormal growth of grains [53,54]. Results have shown that when too little VC and Cr3C2 are added or do not contain inhibitors, there are cases where the grains grow rapidly and grow abnormally. As the content of grain inhibitors increases, the grain size gradually decreases. The abnormal growth of the grain almost disappeared [12,45].

Figure 11 SEM images of (a) WCN-5# fracture and (b) WCN-6# fracture.
Figure 11

SEM images of (a) WCN-5# fracture and (b) WCN-6# fracture.

Figure 12 shows the fracture SEM results of specimens WCN-1#, WCN-2#, WCN-3#, WCN-4#, WCN-5#, and WCN-6#. A small number of small pores were observed on the fracture surface. The fracture modes of specimens WCN-1# and WCN-2# were mainly the crystal fracture, while the specimens WCN-3–6# were mainly transgranular fracture. It was observed that the inner crystal grains of the material were planar, and there were few and very small wrapped inner holes. The reason is mainly because in the vacuum high-temperature sintering process, the degree of densification is mainly improved by the continuous migration of grain boundaries and pores, and the contact between the powders is not tight due to uneven grinding and mixing between the raw material powders and diffusion driving during sintering. When the force is insufficient, the grain boundary migration speed is greater than the velocity of the pore migration, and the pores are generated. If the pores are to be eliminated, the pores are mainly diffused to the grain boundary by vacancies. Figure 12 shows that as the Ni content increases, the grain size of the WC–Ni composite ceramics gradually decreases. The reason may be the effect of VPS–HIP process technology after adding a small amount of binder. And at high temperatures, the powder can be fully contacted during a longtime sintering process and will generate sufficient diffusion driving force, thereby reducing the porosity. The reduction, and at the same time, the grain growth inhibitor component, hinders the abnormal growth of the WC particles, making the structure more uniform and dense, thereby improving the flexural strength of the composites. At the same time, the composites contain mainly the elongated triangular prism WC particles and Ni phase. The existence of nickel makes the fracture of the WC–Ni composite mostly of transgranular type with a small fraction of intergranular fracture mode, which can improve the flexural strength.

Figure 12 Fracture scanning: (a) WCN-1#, (b) WCN-2#, (c) WCN-3#, (d) WCN-4#, (e) WCN-5#, and (f) WCN-6#.
Figure 12

Fracture scanning: (a) WCN-1#, (b) WCN-2#, (c) WCN-3#, (d) WCN-4#, (e) WCN-5#, and (f) WCN-6#.

Figure 13 shows the fracture toughness indentation of the specimen WCN-6#. The brittle fracture occurs at the edge of the indentation and a slight deflection occurs at the crack. When the crack is deflected, the path increases and the crack propagation decreases. In this way, more energy increases the toughness of the composites. Comparing Figure 8 to analyze the effect of Ni content on the microstructure and mechanical properties of WC–Ni composite ceramics materials. Determine that the fracture mechanism is controlled by the Ni phase. When the WC–Ni composite ceramics is under pressure, the presence of the Ni phase will cause some micro-cracks in the WC–Ni composite ceramics, which in turn causes crack deflection, bridging and branch deflection, thereby improving the WC–Ni composite ceramics fracture toughness. At the same time, fracture toughness is also related to internal defects and grain size. It can be seen from the fracture observation of Figure 12 that the specimen WCN-1# has very low internal porosity and has high fracture toughness. According to the data analysis, the fracture toughness of the specimen WCN-1# was 5.76 MPa m1/2.

Figure 13 WCN-6# specimen: (a) fracture toughness indentation map and (b) partial enlarged view.
Figure 13

WCN-6# specimen: (a) fracture toughness indentation map and (b) partial enlarged view.

5 Conclusions

In this work, the microstructure and properties of WC–Ni composite ceramics with WC–0.25 wt% Ni, WC–0.5 wt% Ni, and WC–1 wt% Ni were studied. The following results were obtained:

  1. The addition of NiCl2·6H2O greatly reduces the difficulty of sintering WC–Ni composite ceramics. 0.5 wt% Ni was added to the WC matrix, and VPS at 1,560°C, followed by HIP at 1,600°C × 80 MPa × 1 h to obtain a fully dense WC–Ni composite ceramic.

  2. TEM results show that the WC grains are lath or triangular, and there are dislocations, twins, and other dislocations in the interior. The distribution of the Ni phase in the composite is uniform. The grain growth inhibitors VC and Cr3C2 are mainly located at the crystal angle (edge) and facet interface of WC, which hinders the migration of grain boundaries, limits the growth of WC grains, and inhibits the abnormal growth of WC grains.

  3. As the nickel content increases, the density gradually increases. When the nickel content increases to 1% of the mass fraction, the density decreases, because the liquid phase sintering is too long, resulting in a decrease in density. Specimen WCN-4# multiphase ceramics have a hardness of up to 23.0 GPa, a fracture toughness of 5.28 MPa m1/2, and a flexural strength of 1396.58 MPa, and the WC grains remain fine and do not grow significantly.

  4. The toughening mechanism of Ni to WC–Ni composite ceramics is due to microcracks, crack deflection, crack bridging, and crack bifurcation, which absorbs the energy of crack propagation, thereby increasing the surface area of crack propagation and increasing the fracture toughness of the composite.

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Acknowledgments

This work was supported by Key Laboratory of Infrared Imaging Materials and Detectors, Shanghai Institute of Technical Physics, Chinese Academy of Sciences (No. IIMDKFJJ-19-08), and China Postdoctoral Science Foundation (No. 2015M570794 and 2018T110993).

  1. Conflict of interest: The authors declare no conflict of interest regarding the publication of this paper.

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Received: 2020-04-25
Revised: 2020-04-26
Accepted: 2020-05-17
Published Online: 2020-06-13

© 2020 Xingxing Lyu et al., published by De Gruyter

This work is licensed under the Creative Commons Attribution 4.0 International License.

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  52. Effects of substrate properties and sputtering methods on self-formation of Ag particles on the Ag–Mo(Zr) alloy films
  53. Enhancing carbonation and chloride resistance of autoclaved concrete by incorporating nano-CaCO3
  54. Effect of SiO2 aerogels loading on photocatalytic degradation of nitrobenzene using composites with tetrapod-like ZnO
  55. Radiation-modified wool for adsorption of redox metals and potentially for nanoparticles
  56. Hydration activity, crystal structural, and electronic properties studies of Ba-doped dicalcium silicate
  57. Microstructure and mechanical properties of brazing joint of silver-based composite filler metal
  58. Polymer nanocomposite sunlight spectrum down-converters made by open-air PLD
  59. Cryogenic milling and formation of nanostructured machined surface of AISI 4340
  60. Braided composite stent for peripheral vascular applications
  61. Effect of cinnamon essential oil on morphological, flammability and thermal properties of nanocellulose fibre–reinforced starch biopolymer composites
  62. Study on influencing factors of photocatalytic performance of CdS/TiO2 nanocomposite concrete
  63. Improving flexural and dielectric properties of carbon fiber epoxy composite laminates reinforced with carbon nanotubes interlayer using electrospray deposition
  64. Scalable fabrication of carbon materials based silicon rubber for highly stretchable e-textile sensor
  65. Degradation modeling of poly-l-lactide acid (PLLA) bioresorbable vascular scaffold within a coronary artery
  66. Combining Zn0.76Co0.24S with S-doped graphene as high-performance anode materials for lithium- and sodium-ion batteries
  67. Synthesis of functionalized carbon nanotubes for fluorescent biosensors
  68. Effect of nano-silica slurry on engineering, X-ray, and γ-ray attenuation characteristics of steel slag high-strength heavyweight concrete
  69. Incorporation of redox-active polyimide binder into LiFePO4 cathode for high-rate electrochemical energy storage
  70. Microstructural evolution and properties of Cu–20 wt% Ag alloy wire by multi-pass continuous drawing
  71. Transparent ultraviolet-shielding composite films made from dispersing pristine zinc oxide nanoparticles in low-density polyethylene
  72. Microfluidic-assisted synthesis and modelling of monodispersed magnetic nanocomposites for biomedical applications
  73. Preparation and piezoresistivity of carbon nanotube-coated sand reinforced cement mortar
  74. Vibrational analysis of an irregular single-walled carbon nanotube incorporating initial stress effects
  75. Study of the material engineering properties of high-density poly(ethylene)/perlite nanocomposite materials
  76. Single pulse laser removal of indium tin oxide film on glass and polyethylene terephthalate by nanosecond and femtosecond laser
  77. Mechanical reinforcement with enhanced electrical and heat conduction of epoxy resin by polyaniline and graphene nanoplatelets
  78. High-efficiency method for recycling lithium from spent LiFePO4 cathode
  79. Degradable tough chitosan dressing for skin wound recovery
  80. Static and dynamic analyses of auxetic hybrid FRC/CNTRC laminated plates
  81. Review articles
  82. Carbon nanomaterials enhanced cement-based composites: advances and challenges
  83. Review on the research progress of cement-based and geopolymer materials modified by graphene and graphene oxide
  84. Review on modeling and application of chemical mechanical polishing
  85. Research on the interface properties and strengthening–toughening mechanism of nanocarbon-toughened ceramic matrix composites
  86. Advances in modelling and analysis of nano structures: a review
  87. Mechanical properties of nanomaterials: A review
  88. New generation of oxide-based nanoparticles for the applications in early cancer detection and diagnostics
  89. A review on the properties, reinforcing effects, and commercialization of nanomaterials for cement-based materials
  90. Recent development and applications of nanomaterials for cancer immunotherapy
  91. Advances in biomaterials for adipose tissue reconstruction in plastic surgery
  92. Advances of graphene- and graphene oxide-modified cementitious materials
  93. Theories for triboelectric nanogenerators: A comprehensive review
  94. Nanotechnology of diamondoids for the fabrication of nanostructured systems
  95. Material advancement in technological development for the 5G wireless communications
  96. Nanoengineering in biomedicine: Current development and future perspectives
  97. Recent advances in ocean wave energy harvesting by triboelectric nanogenerator: An overview
  98. Application of nanoscale zero-valent iron in hexavalent chromium-contaminated soil: A review
  99. Carbon nanotube–reinforced polymer composite for electromagnetic interference application: A review
  100. Functionalized layered double hydroxide applied to heavy metal ions absorption: A review
  101. A new classification method of nanotechnology for design integration in biomaterials
  102. Finite element analysis of natural fibers composites: A review
  103. Phase change materials for building construction: An overview of nano-/micro-encapsulation
  104. Recent advance in surface modification for regulating cell adhesion and behaviors
  105. Hyaluronic acid as a bioactive component for bone tissue regeneration: Fabrication, modification, properties, and biological functions
  106. Theoretical calculation of a TiO2-based photocatalyst in the field of water splitting: A review
  107. Two-photon polymerization nanolithography technology for fabrication of stimulus-responsive micro/nano-structures for biomedical applications
  108. A review of passive methods in microchannel heat sink application through advanced geometric structure and nanofluids: Current advancements and challenges
  109. Stress effect on 3D culturing of MC3T3-E1 cells on microporous bovine bone slices
  110. Progress in magnetic Fe3O4 nanomaterials in magnetic resonance imaging
  111. Synthesis of graphene: Potential carbon precursors and approaches
  112. A comprehensive review of the influences of nanoparticles as a fuel additive in an internal combustion engine (ICE)
  113. Advances in layered double hydroxide-based ternary nanocomposites for photocatalysis of contaminants in water
  114. Analysis of functionally graded carbon nanotube-reinforced composite structures: A review
  115. Application of nanomaterials in ultra-high performance concrete: A review
  116. Therapeutic strategies and potential implications of silver nanoparticles in the management of skin cancer
  117. Advanced nickel nanoparticles technology: From synthesis to applications
  118. Cobalt magnetic nanoparticles as theranostics: Conceivable or forgettable?
  119. Progress in construction of bio-inspired physico-antimicrobial surfaces
  120. From materials to devices using fused deposition modeling: A state-of-art review
  121. A review for modified Li composite anode: Principle, preparation and challenge
  122. Naturally or artificially constructed nanocellulose architectures for epoxy composites: A review
https://doi.org/10.1515/ntrev-2020-0044
Keywords for this article
WC–Ni composite ceramics; hot isostatic pressing; densification
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  58. Polymer nanocomposite sunlight spectrum down-converters made by open-air PLD
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  62. Study on influencing factors of photocatalytic performance of CdS/TiO2 nanocomposite concrete
  63. Improving flexural and dielectric properties of carbon fiber epoxy composite laminates reinforced with carbon nanotubes interlayer using electrospray deposition
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  65. Degradation modeling of poly-l-lactide acid (PLLA) bioresorbable vascular scaffold within a coronary artery
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  67. Synthesis of functionalized carbon nanotubes for fluorescent biosensors
  68. Effect of nano-silica slurry on engineering, X-ray, and γ-ray attenuation characteristics of steel slag high-strength heavyweight concrete
  69. Incorporation of redox-active polyimide binder into LiFePO4 cathode for high-rate electrochemical energy storage
  70. Microstructural evolution and properties of Cu–20 wt% Ag alloy wire by multi-pass continuous drawing
  71. Transparent ultraviolet-shielding composite films made from dispersing pristine zinc oxide nanoparticles in low-density polyethylene
  72. Microfluidic-assisted synthesis and modelling of monodispersed magnetic nanocomposites for biomedical applications
  73. Preparation and piezoresistivity of carbon nanotube-coated sand reinforced cement mortar
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  76. Single pulse laser removal of indium tin oxide film on glass and polyethylene terephthalate by nanosecond and femtosecond laser
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  78. High-efficiency method for recycling lithium from spent LiFePO4 cathode
  79. Degradable tough chitosan dressing for skin wound recovery
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  83. Review on the research progress of cement-based and geopolymer materials modified by graphene and graphene oxide
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  85. Research on the interface properties and strengthening–toughening mechanism of nanocarbon-toughened ceramic matrix composites
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  94. Nanotechnology of diamondoids for the fabrication of nanostructured systems
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  96. Nanoengineering in biomedicine: Current development and future perspectives
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  106. Theoretical calculation of a TiO2-based photocatalyst in the field of water splitting: A review
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  117. Advanced nickel nanoparticles technology: From synthesis to applications
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