# Ti(C, N)-Based Cermets with Two Kinds of Core-Rim Structures Constructed byβ-Co Microspheres

Authors: Hui Yan; Ying Deng; Yong Yao Su; Shan Jiang; Qiao Wang Chen; Shi Xiu Cao; Bing Liu
Journal: Advances in Materials Science and Engineering
 (2020)

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## Abstract

Ti(C, N)-based cermet materials represent the best choice of materials for the manufacturing industry and military products. In this study, we use cubicβ-cobalt(β-Co) as the binder phase to strengthen constructed cermets. At the same time, to optimize the microstructure, (Ti, W, Mo, Ta) (C, N) powders are added to form two kinds of core-rim morphologies. Here, β-cobalt powders with face-centered cubic structures are obtained by the solid-phase reaction of high-energy ball milling. A solid-phase chemical reaction and a carbothermic reduction-nitridation method are used to prepare the (Ti, W, Mo, Ta) (C, N) powders. In our process, first we mix the cobalt and (Ti, W, Mo, Ta) (C, N) powders; then we press the powder mixtures into rectangular samples and sinter them in a pressure sintering furnace to obtain Ti(C, N)-based cermets with two kinds of core-rim structure, that is, black core/white rim and white core/gray rim. The results show that the new cermets demonstrate excellent toughness and performance.


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## Body

## 1. Introduction

Ti(C, N)-based cermets are ceramic-metal composite materials that have excellent strength and high-temperature hardness. These materials are prepared by powder metallurgy methods [1–3], which are composed of a ceramic phase and a metal binder phase [4, 5]. Not only do cermets retain the high strength, high hardness, high-temperature resistance, antioxidation, and chemical stability characteristics of the ceramic phase, but they also possess a certain degree of plasticity from the metal phase [6–11]. However, Ti(C, N)-based cermets are still characterized by many deficiencies, their toughness is a fatal weakness, and their microstructures need further improvement [12].Ti(C, N)-based cermet powders have a typical core-rim structure, where the composition and lattice constants between the core and the rim are quite different. In the process of polyphase sintering, the rim phase is easy to cause interface stress, induce the residual stress and component segregation, which lead to defects, and reduce their strength and toughness. Previous work has shown that micro-nano cermet materials can alleviate the contradiction between hardness and strength, whose bending strength and hardness can reach 2100 MPa and 92.7 HRA, respectively. However, the contradiction between hardness and strength still exists such that one of the critical concerns of ceramics in high stress condition is their low fracture toughness (3.2–3.9 MPa·m1/2) [13–17]. How to design the material structure and solve the toughness problem of cermet needs to be further studied.Based on the toughness characteristics ofβ-Co and the improved hardness and strength properties of micro-nano materials [18], this study aimed to improve the microstructure of the material using the face-centered cubic β-Co powder and (Ti, W, Mo, Ta) (C, N) solid solution powder. These materials were obtained in our laboratory, which allowed us to develop a more advanced Ti(C, N)-based cermet structure that had two kinds of core-rim structures. The newly formed material is composed of a black core/white rim and gray core/white rim, which reduce the number of brittle black cores in their structures. Moreover, the gray cores have almost the same composition as the rim phases, which reduce the difference between the cores to a large extent. At the same time, because of the different structure of the two kinds of core rims, stress transfer is alleviated, crack deflection is improved, and the toughness is improved.


## 2. Preparation Strategy and Experimental Procedure

In this study, CoCll2 · 6H2O and Na2CO3 are used as raw materials. The precursors were prepared by the solid-state reaction of high-energy ball milling, and the precursor powders were washed several times with deionized water and then spray-dried. The dried precursor powders were put into a furnace to decompose for 0.5–8 hr, under a nitrogen atmosphere at high temperature (400–450°C); then they were cooled to room temperature, whereby the target Co powders were obtained.30% MoO3, 10% WO3, 30% TiO2, and 30% Ta2O5 (mass fraction) were weighed, and an appropriate amount of alcohol was added. The mixed materials were then placed into a stainless steel vacuum metal can for ball milling (a planetary ball mill with a cermet grinding ball; the ball mass ratio was 10 : 1, the rotation speed was 500 rev·min−1, and the ball milling times were 3 hr, 6 hr, 9 hr, 12 hr). After ball milling, the slurry was sieved. After sieving, washing, precipitating, and spray drying, the sample and carbon powder were carbonized in a vacuum tube furnace according to a certain process curve, using a nitrogen atmosphere (flow rate 600 ml·min−1, furnace pressure 0.1–0.2 MPa) and temperatures of 1200°C and 1600°C, for 5 hr.The cermet sample powders were prepared according to the components shown in Table1. The mixed-powder samples A, B, and C were poured into the ball-grinding tank, where 3.5% paraffin and an appropriate amount of alcohol were added. Then, high-energy ball milling was carried out (ball mass ratio: 5 : 1, rotation speed of 200 rev·min−1, ball milling time of 48 hr). After milling, the slurry was discharged for spray drying, and then the mixtures were compacted under a pressure of 100 MPa. Finally, the pressure samples were loaded into the pressure sintering furnace for overpressure sintering. The sintering process is shown in Figure 1 [19].Table 1
Composition of powder mixtures for the Ti(C, N)-based cermet (wt.%).
Samples
Ti(C, N)
WC
MoC
TaC
α/β-Co
β-Co
(Ti, W, Mo, Ta) (C, N)
A
50
20
10
7
13
0
0
B
50
20
10
7
0
13
0
C
50
0
0
0
0
13
37Figure 1
Curve of cermet sintering process.The phases and specimens were analyzed using a Japanese D/MAX2500 X-ray diffractometer. The morphologies and microstructures of the samples were observed with a Jeol-6490lv scanning electron microscope. The transverse fracture strength (TRS) of each sample was tested with a We-100b universal material testing machine, and the three-point bending method was adopted. The size of the sample was 20 × 6.5 × 5.5 mm (length width and height), the span was 30 mm, and the loading speed was 0.5 mm·min−1. The Rockwell hardness of each sample was measured with an AR-600 Rockwell hardness machine. The Vickers hardness of each sample material was tested with a HV-10 tester, and the fracture toughness values were calculated using the Shetty fracture toughness formula.


## 3. Results and Discussion

### 3.1. Structure and Morphology of the Cobalt Powder

The reaction equation that may occur during the preparation of the cobalt powders is as shown in the following formula:(1)CoCl2⋅H2O+NH4HCO3=CoCO3+NH3+2HCl+H2O(2)CoCO3+H2=Co+CO2+H2OThe structure and morphology of cobalt is shown in Figure2. Sample A is the Co powder created in this study, while sample B is a commercial Co powder for comparison. From Figure 2, we can see that the structure of sample A is face-centered cubic (fcc, β-Co) and it has a spherical morphology. Meanwhile, the structure of the commercial cobalt powder sample b is hexagonal close packed (hcp, α-Co) with a dendritic morphology. Because fcc structure has more sliding ability than hcp, it has a stronger resistance to plastic deformation. In addition, the spherical powder particles are in point contact with each other, which is not easy to adhere. Therefore, it will have better dispersion when applied to cermets to reduce the cobalt pool, which in turn will affect the mechanical properties of the cermet. This verifies the excellent structure and morphology of the β-Co powders obtained from this study. As the binder phase of cermet, it has better strength and toughness than traditional cobalt powders.Figure 2
(a, b) The structure and morphology of the cobalt powder as detected by XRD and SEM. (c) Graph of (i) theβ-Co powder and (ii) the commercial cobalt powder (α/β-Co).In this study, we used a mechanical solid-phase chemical method and ultrasonic spraying. These resulted in the spherical morphology of the CoCO3 powders, which have high internal energies. The mechanical force of the high-energy ball mill assured that the power retained sufficient internal energy, which ensured that the CoCO3 precursor thermally decomposed to β-Co, which has higher formation energy than α-Co. The internal lattice of the material was deformed by the high-speed grinding ball, whereby the number of internal lattice defects increased and the thermal vibration energy within the lattice was generated, so that the diffusion coefficient increased and the reaction activation energy decreased, which resulted in the solid-phase chemical reaction occurring at low temperature [19].


### 3.2. Structure and Morphology of the (Ti, W, Mo, Ta) (C, N) Solid Solution

The possible reaction process of carbothermic and nitrogen with (Ti, W, Mo, Ta) (C, N) is(3)12Ta2O5+WO3+MoO3+TiO2+232+C+12+N2=Ti, W, Mo, TaC, N+212COFigure3 shows the TG-DSC analysis of reaction 3. It can be seen from the figure that at 200°C several consecutive endothermic peaks begin to appear. At about 800°C, the reaction is violent and the powder begins to carbonize, and after 1500°C the DSC curve tends to be gentle and the reaction is more stable. This indicates that the raw powders are still in a reactive stage before 1500°C, whereafter they began to stabilize. At around 1600°C the solid solution reaction has been completed [19].Figure 3
TG-DSC curves of (Ti, W Mo, Ta) (C, N) solid solution powders.Figure4 shows the results of XRD patterns of the carbothermic reduction-nitridation reaction at different temperatures. From the figure, we can see that when the sintering temperature is 1100°C, the product is composed of metal carbide, oxide, and a small amount of (Ti, W, Mo, Ta) (C, N) solid solution phase. Compared with Figure 4(ii), the relative strength of the diffraction peak of the metal oxide is weaker, and a large number of metal carbide phases have appeared. This indicates that at this stage, the reaction between oxide and carbon is very intense, and a large amount of metal oxide has reduced. At 1450°C, the product is composed of a (Ti, W, Mo, Ta) (C, N) solid solution phase and very little WC phase. The diffraction peak of the metal carbide and oxide has disappeared, which indicates that the reaction has likely completed at this stage. When the temperature increased to 1500°C, the diffraction peak of the metal carbide disappeared, the relative strength of the C, N diffraction peak was further enhanced, and no other phase was detected. This indicates that the solid solution reaction had been carried out at this time, and the product is completely composed of a pure C, N solid solution phase. The elements Ti, W, Mo, and Ta completely dissolved into the C, N crystal, as shown in Figure 4. This also verifies the TG-DSC results of Figure 3.Figure 4
X-ray diffraction patterns of solid solution powders at different sintering temperatures: (i) 20°C; (ii) 900°C; (iii) 1100°C; (iv) 1450°C; (v) 1500°C.Figures5(a)–5(d) provide SEM images of the powders obtained from the carbothermic reduction-nitridation reaction for different milling times of the raw materials. It can be seen from the figure that the particle sizes of the (Ti, W, Mo, Ta) (C, N) separable particles decrease as the milling time increases; however there is little difference between those milled for 9 hr and 12 hr. This indicates that after 9 hr of milling, the milling effect is good. This provides a theoretical basis for the final setting of the process parameters for 9 hr of ball milling and a carbonization temperature of 1600°C [19].Figure 5
Morphology of (Ti, W Mo, Ta) (C, N) powders, as obtained by SEM. (a) Ball mill 3 h; (b) ball mill 6 h; (c) ball mill 9 h; (d) ball mill 12 h.
(a)
(b)
(c)
(d)Figure6 shows the morphology and energy spectrum of the (Ti, W, Mo, Ta) (C, N) solid solution obtained in this study. It can be seen from its topography that the morphology of the solid solution is obviously refined, which has spherical-like morphology and particle size of about 200 nm [19].Figure 6
SEM and EDS images of (Ti, Ta, Mo, W) (C, N) solid solution powder prepared at 1500°C.
(a)
(b)


### 3.3. Solid Solution Level of Heavy Metals in Cermets

Figure7 shows the XRD parameters of the structure of the cermet samples A, B, and C according to the compositions given in Table 1 and the sinter techniques presented in Figure 1. It can be seen from the figure that the characteristic peaks of all samples are consistent with the standard PDF card of Ti(C, N), proving that a C, N solid solution can be prepared under the conditions of this experiment. From Figure 7, we can also observe that the crystalline structure of the Co powder of sample A is hcp, while samples B and C have fcc structures.Figure 7
XRD parameters of the structure of the cermet.It is well known that when a solid solution reaction of Ti(C, N) occurs, heavy metal elements will dissolve into the (111) crystal plane of Ti(C, N), and the diffraction peak strength of (111) will be enhanced, while the diffraction peak strength of (200) will not change much. Therefore, the ratio between the diffraction peak strength I(111)/I(200) is large, indicating that the solid solubility of the heavy metal elements is greater. Table2 shows the relative strengths of the solid solution in the cermet samples. From the table, it can be seen that the ratio of I(111)/I(200) of the samples is A < B < C. This suggests that solid solution level of sample C is the best, which indicates that the addition of Co and (Ti, W, Mo, Ta) (C, N) will improve the solid solution level of the heavy metal elements and hence improve the structure of the cermet.Table 2
Relative strength of the solid solution phase of cermet samples.
Samples
(Ti, W, Mo, Ta) (C, N)
Co
(111)
(200)
I(111)/I(200)
(111) fcc
(200)fcc
(101) hcp
A
100
99.6
1.004
52.5
19
14.1
B
100
76.6
1.30
39.5
23.3
—
C
100
57.1
1.75
40.8
17.2
—


### 3.4. Cermet Microstructure

Figure8 shows the BSE microstructure of cermet samples A, B, and C under a certain process curve. According to the backscattering principle, the higher the atomic number of an element, the brighter its image. So, in Figure 9, the brighter the region, the higher the heavy metal content of elements (W, Mo, Ta).Figure 8
BSE morphology of Ti(C, N)-based cermets: (a) ordinary cermet, (b) addedβ-Co powder, and (c) added W Mo, Ta (C, N) and β-Co power.
(a)
(b)
(c)Figure 9
BSE morphology of crack propagation paths.
(a)
(b)
(c)From Figure8, it can be seen that the three samples show different core-rim characteristics. Sample A is a traditional cermet, and as the sinter process progressed via the influence of dissolution-precipitation reaction, it showed the standard structure core-rim structure: inner rim/outer rim. First the Ti(C, N), WC, MoC, and TaC in sample A dissolved into binder phase Co, which was followed by the metal elements Ti, W, Mo, and Ta precipitating out and covering the surface of Ti(C, N) (black core), which did not dissolve, to form a white rim phase while the concentration of heavy metal elements saturated. Then, the metal elements underwent an additional dissolution-precipitation reaction, and it was found that the concentration of heavy metal elements was lower, and a gray rim phase had formed. Finally, the standard structure (black core/white rim/gray rim) of sample A formed. Sample B had β-Co as the binding phase. Based on the analysis given in Section 3.3, β-Co contains a higher internal energy, which will increase the driving force of the solid-phase sintering. This will accelerate the spreading speed of the metal elements during the solid-phase sintering stage, and thus, improve the solid solubility of the metallic elements in TiCN and Co, eventually forming smaller cores and thicker rims than sample A. Sample C used (Ti, W, Mo, Ta) (C, N) instead of the WC, MoC, and TaC powders and had β-Co as its binder phase. During the sintering process, the elements Ti, W, Mo, and Ta in the lattice of (Ti, W, Mo, Ta) (C, N) and the Ti elements in Ti(C, N) began to diffuse with the increase of temperature, and gradually dissolve into the Co lattice. When the concentration of metallic elements in the Co phase reached saturation, the precipitated phases (Ti, W, Mo, Ta) (C, N) appeared and covered on the surface of the undissolved Ti(C, N) and (Ti, W, Mo, Ta) (C, N) particles and formed two types of core/rim structures. Similarly, the higher internal energy of β-Co resulted in higher solid solubility of the metallic elements, which presented a perfect thin core and thick rim structure.


### 3.5. Mechanical Properties of Cermet

Figure10 shows the mechanical properties of the Ti(C, N) cermet samples A, B, and C. It can be seen from the figure that the comprehensive performance of sample C is the best, followed by sample B. The transverse rupture strength (TRS) of sample C is 7% higher, and its fracture toughness (KIC) is 10% larger, than sample B. The comprehensive mechanical properties of sample C are better than the cermets prepared by Jie Wang et al.. However, the trends of the fracture toughness and bending strength of the samples are consistent. It can be clearly seen from Figure 9 that there are many black cores in the microstructure of sample A, where some of the core-rim structures are not fully developed, which directly affects the performance of the cermet. The size of the black cores in samples B and C became smaller, and at the same time, part of the black core of sample C was replaced by a white core with higher toughness, thereby improving its strength and toughness.Figure 10
Histogram showing comparison of mechanical properties. Five samples were recorded for each sample, and the error bars represent differences between the maximum and minimum values of five recordings from the three independent samples.Figure9 shows a BSE figure of fractures and cracks in the samples. As shown in the figure, the fracture modes of samples B and C are dominated by intergranular fractures, while that of sample A is dominated by transgranular fractures. This indicates that β-Co has a strength-toughening effect, and the solid solution strengthening effect of the metal bonding phase is more obvious. At the same time, the chemical composition of the white cores of the two types of core-rim structure is almost the same as the gray rim phase, which minimizes the difference between the cores and the rims. Moreover, because of the different core-rims types, stress transfer is alleviated, crack deflection can be improved, and toughness can be improved.


### 3.6. Effects of the Microstructure on the Mechanical Properties

#### 3.6.1. Core/Rim Interface

Figure11 illustrates the process of forming the core-rim structure of sample C. The core-rim structure of the Ti(C, N) cermet is an important factor that determines its comprehensive performance. The rim phase can greatly improve the wettability between the hard ceramic phase and the metal binder phase and prevent the hard phase from aggregating and growing. The energy spectrum analysis (Figure 12) shows that the chemical composition of the white core phase is close to that of the gray rim phase, and their crystal structures are essentially the same. Therefore, the lattice mismatch of the core-rim structure, which is caused by the different core-rim lattice constants, is small, and the internal stress of the core-rim interface is correspondingly reduced, which results in lower core-rim interfacial stress and stronger bonding at the core-rim interface.Figure 11
Formation process of the core-rim structure.Figure 12
Energy spectra of a white core and gray rim. (a) BSE morphology of core-rim, (b) energy spectrum of (i) white core, and (c) energy spectrum of (ii) gray rim.
(a)
(b)
(c)


#### 3.6.2. Rim/Binder Face Interface

According to classical wetting thermodynamics, the wettability of the metal to ceramic phases is the basic criterion that determines the interface bonding between metals and ceramics. The physical quantity that represents the degree of wettability of the metal to the ceramic phase is the contact angle (θ) and adhesion force, as shown in Figure 13(a), which is the famous Young equation:(4)Wad=γB1+cosθ,where γB is the interfacial free energy of the metal phase and Wad is the work of adhesion between the metal phase and the ceramic phase. As can be seen from equation (4), in the case of a certain cosθ value, the larger γB is, the larger Wad is, the greater the adhesion between metal and ceramic is, the better the wettability of metal/ceramic interface is, the more stable the formed interface is, and the better the mechanical properties of the cermets are.Figure 13
(a) Touch angle, (b) rim/binder face of sample C, and (c) the rim/binder face of sample A.Figures13(b) and 13(c) show illustrations of the rim/binder face of sample C and sample A, respectively. As previously discussed in this paper, while the white and gray rims contain the (Ti, W, Mo, Ta) (C, N) solid solution, they have different Ti, W, Mo, and Ta concentrations, and their microstructure is the same as Ti(C, N) and β-Co, which have fcc crystal structures and approximately the same lattice parameters and plane orientations. Therefore, their interface atomic matching degree is high and their lattice mismatch is low, which results in a coherent relationship and the touch angle between the rim/binder phases being very small, as shown in Figure 13(b). According to equation (4), the smaller θ is, the larger is Wad. Therefore, the adhesion force at the rim/binder interface is stronger, and it can improve the wettability between ceramic grains and the binder. In contrast, the touch angle between the rim/binder phases of sample A is large (as shown in Figure 13(c)), which will result in smaller Wad.In addition,β-Co used in this study has a higher interfacial free energy than ordinary Co, so its γB is higher than that of ordinary α-Co. Hence, it has a more significant wetting effect on the ceramic phase, thus improving the mechanical properties of the cermet.


## 3.1. Structure and Morphology of the Cobalt Powder

The reaction equation that may occur during the preparation of the cobalt powders is as shown in the following formula:(1)CoCl2⋅H2O+NH4HCO3=CoCO3+NH3+2HCl+H2O(2)CoCO3+H2=Co+CO2+H2OThe structure and morphology of cobalt is shown in Figure2. Sample A is the Co powder created in this study, while sample B is a commercial Co powder for comparison. From Figure 2, we can see that the structure of sample A is face-centered cubic (fcc, β-Co) and it has a spherical morphology. Meanwhile, the structure of the commercial cobalt powder sample b is hexagonal close packed (hcp, α-Co) with a dendritic morphology. Because fcc structure has more sliding ability than hcp, it has a stronger resistance to plastic deformation. In addition, the spherical powder particles are in point contact with each other, which is not easy to adhere. Therefore, it will have better dispersion when applied to cermets to reduce the cobalt pool, which in turn will affect the mechanical properties of the cermet. This verifies the excellent structure and morphology of the β-Co powders obtained from this study. As the binder phase of cermet, it has better strength and toughness than traditional cobalt powders.Figure 2
(a, b) The structure and morphology of the cobalt powder as detected by XRD and SEM. (c) Graph of (i) theβ-Co powder and (ii) the commercial cobalt powder (α/β-Co).In this study, we used a mechanical solid-phase chemical method and ultrasonic spraying. These resulted in the spherical morphology of the CoCO3 powders, which have high internal energies. The mechanical force of the high-energy ball mill assured that the power retained sufficient internal energy, which ensured that the CoCO3 precursor thermally decomposed to β-Co, which has higher formation energy than α-Co. The internal lattice of the material was deformed by the high-speed grinding ball, whereby the number of internal lattice defects increased and the thermal vibration energy within the lattice was generated, so that the diffusion coefficient increased and the reaction activation energy decreased, which resulted in the solid-phase chemical reaction occurring at low temperature [19].


## 3.2. Structure and Morphology of the (Ti, W, Mo, Ta) (C, N) Solid Solution

The possible reaction process of carbothermic and nitrogen with (Ti, W, Mo, Ta) (C, N) is(3)12Ta2O5+WO3+MoO3+TiO2+232+C+12+N2=Ti, W, Mo, TaC, N+212COFigure3 shows the TG-DSC analysis of reaction 3. It can be seen from the figure that at 200°C several consecutive endothermic peaks begin to appear. At about 800°C, the reaction is violent and the powder begins to carbonize, and after 1500°C the DSC curve tends to be gentle and the reaction is more stable. This indicates that the raw powders are still in a reactive stage before 1500°C, whereafter they began to stabilize. At around 1600°C the solid solution reaction has been completed [19].Figure 3
TG-DSC curves of (Ti, W Mo, Ta) (C, N) solid solution powders.Figure4 shows the results of XRD patterns of the carbothermic reduction-nitridation reaction at different temperatures. From the figure, we can see that when the sintering temperature is 1100°C, the product is composed of metal carbide, oxide, and a small amount of (Ti, W, Mo, Ta) (C, N) solid solution phase. Compared with Figure 4(ii), the relative strength of the diffraction peak of the metal oxide is weaker, and a large number of metal carbide phases have appeared. This indicates that at this stage, the reaction between oxide and carbon is very intense, and a large amount of metal oxide has reduced. At 1450°C, the product is composed of a (Ti, W, Mo, Ta) (C, N) solid solution phase and very little WC phase. The diffraction peak of the metal carbide and oxide has disappeared, which indicates that the reaction has likely completed at this stage. When the temperature increased to 1500°C, the diffraction peak of the metal carbide disappeared, the relative strength of the C, N diffraction peak was further enhanced, and no other phase was detected. This indicates that the solid solution reaction had been carried out at this time, and the product is completely composed of a pure C, N solid solution phase. The elements Ti, W, Mo, and Ta completely dissolved into the C, N crystal, as shown in Figure 4. This also verifies the TG-DSC results of Figure 3.Figure 4
X-ray diffraction patterns of solid solution powders at different sintering temperatures: (i) 20°C; (ii) 900°C; (iii) 1100°C; (iv) 1450°C; (v) 1500°C.Figures5(a)–5(d) provide SEM images of the powders obtained from the carbothermic reduction-nitridation reaction for different milling times of the raw materials. It can be seen from the figure that the particle sizes of the (Ti, W, Mo, Ta) (C, N) separable particles decrease as the milling time increases; however there is little difference between those milled for 9 hr and 12 hr. This indicates that after 9 hr of milling, the milling effect is good. This provides a theoretical basis for the final setting of the process parameters for 9 hr of ball milling and a carbonization temperature of 1600°C [19].Figure 5
Morphology of (Ti, W Mo, Ta) (C, N) powders, as obtained by SEM. (a) Ball mill 3 h; (b) ball mill 6 h; (c) ball mill 9 h; (d) ball mill 12 h.
(a)
(b)
(c)
(d)Figure6 shows the morphology and energy spectrum of the (Ti, W, Mo, Ta) (C, N) solid solution obtained in this study. It can be seen from its topography that the morphology of the solid solution is obviously refined, which has spherical-like morphology and particle size of about 200 nm [19].Figure 6
SEM and EDS images of (Ti, Ta, Mo, W) (C, N) solid solution powder prepared at 1500°C.
(a)
(b)


## 3.3. Solid Solution Level of Heavy Metals in Cermets

Figure7 shows the XRD parameters of the structure of the cermet samples A, B, and C according to the compositions given in Table 1 and the sinter techniques presented in Figure 1. It can be seen from the figure that the characteristic peaks of all samples are consistent with the standard PDF card of Ti(C, N), proving that a C, N solid solution can be prepared under the conditions of this experiment. From Figure 7, we can also observe that the crystalline structure of the Co powder of sample A is hcp, while samples B and C have fcc structures.Figure 7
XRD parameters of the structure of the cermet.It is well known that when a solid solution reaction of Ti(C, N) occurs, heavy metal elements will dissolve into the (111) crystal plane of Ti(C, N), and the diffraction peak strength of (111) will be enhanced, while the diffraction peak strength of (200) will not change much. Therefore, the ratio between the diffraction peak strength I(111)/I(200) is large, indicating that the solid solubility of the heavy metal elements is greater. Table2 shows the relative strengths of the solid solution in the cermet samples. From the table, it can be seen that the ratio of I(111)/I(200) of the samples is A < B < C. This suggests that solid solution level of sample C is the best, which indicates that the addition of Co and (Ti, W, Mo, Ta) (C, N) will improve the solid solution level of the heavy metal elements and hence improve the structure of the cermet.Table 2
Relative strength of the solid solution phase of cermet samples.
Samples
(Ti, W, Mo, Ta) (C, N)
Co
(111)
(200)
I(111)/I(200)
(111) fcc
(200)fcc
(101) hcp
A
100
99.6
1.004
52.5
19
14.1
B
100
76.6
1.30
39.5
23.3
—
C
100
57.1
1.75
40.8
17.2
—


## 3.4. Cermet Microstructure

Figure8 shows the BSE microstructure of cermet samples A, B, and C under a certain process curve. According to the backscattering principle, the higher the atomic number of an element, the brighter its image. So, in Figure 9, the brighter the region, the higher the heavy metal content of elements (W, Mo, Ta).Figure 8
BSE morphology of Ti(C, N)-based cermets: (a) ordinary cermet, (b) addedβ-Co powder, and (c) added W Mo, Ta (C, N) and β-Co power.
(a)
(b)
(c)Figure 9
BSE morphology of crack propagation paths.
(a)
(b)
(c)From Figure8, it can be seen that the three samples show different core-rim characteristics. Sample A is a traditional cermet, and as the sinter process progressed via the influence of dissolution-precipitation reaction, it showed the standard structure core-rim structure: inner rim/outer rim. First the Ti(C, N), WC, MoC, and TaC in sample A dissolved into binder phase Co, which was followed by the metal elements Ti, W, Mo, and Ta precipitating out and covering the surface of Ti(C, N) (black core), which did not dissolve, to form a white rim phase while the concentration of heavy metal elements saturated. Then, the metal elements underwent an additional dissolution-precipitation reaction, and it was found that the concentration of heavy metal elements was lower, and a gray rim phase had formed. Finally, the standard structure (black core/white rim/gray rim) of sample A formed. Sample B had β-Co as the binding phase. Based on the analysis given in Section 3.3, β-Co contains a higher internal energy, which will increase the driving force of the solid-phase sintering. This will accelerate the spreading speed of the metal elements during the solid-phase sintering stage, and thus, improve the solid solubility of the metallic elements in TiCN and Co, eventually forming smaller cores and thicker rims than sample A. Sample C used (Ti, W, Mo, Ta) (C, N) instead of the WC, MoC, and TaC powders and had β-Co as its binder phase. During the sintering process, the elements Ti, W, Mo, and Ta in the lattice of (Ti, W, Mo, Ta) (C, N) and the Ti elements in Ti(C, N) began to diffuse with the increase of temperature, and gradually dissolve into the Co lattice. When the concentration of metallic elements in the Co phase reached saturation, the precipitated phases (Ti, W, Mo, Ta) (C, N) appeared and covered on the surface of the undissolved Ti(C, N) and (Ti, W, Mo, Ta) (C, N) particles and formed two types of core/rim structures. Similarly, the higher internal energy of β-Co resulted in higher solid solubility of the metallic elements, which presented a perfect thin core and thick rim structure.


## 3.5. Mechanical Properties of Cermet

Figure10 shows the mechanical properties of the Ti(C, N) cermet samples A, B, and C. It can be seen from the figure that the comprehensive performance of sample C is the best, followed by sample B. The transverse rupture strength (TRS) of sample C is 7% higher, and its fracture toughness (KIC) is 10% larger, than sample B. The comprehensive mechanical properties of sample C are better than the cermets prepared by Jie Wang et al.. However, the trends of the fracture toughness and bending strength of the samples are consistent. It can be clearly seen from Figure 9 that there are many black cores in the microstructure of sample A, where some of the core-rim structures are not fully developed, which directly affects the performance of the cermet. The size of the black cores in samples B and C became smaller, and at the same time, part of the black core of sample C was replaced by a white core with higher toughness, thereby improving its strength and toughness.Figure 10
Histogram showing comparison of mechanical properties. Five samples were recorded for each sample, and the error bars represent differences between the maximum and minimum values of five recordings from the three independent samples.Figure9 shows a BSE figure of fractures and cracks in the samples. As shown in the figure, the fracture modes of samples B and C are dominated by intergranular fractures, while that of sample A is dominated by transgranular fractures. This indicates that β-Co has a strength-toughening effect, and the solid solution strengthening effect of the metal bonding phase is more obvious. At the same time, the chemical composition of the white cores of the two types of core-rim structure is almost the same as the gray rim phase, which minimizes the difference between the cores and the rims. Moreover, because of the different core-rims types, stress transfer is alleviated, crack deflection can be improved, and toughness can be improved.


## 3.6. Effects of the Microstructure on the Mechanical Properties

### 3.6.1. Core/Rim Interface

Figure11 illustrates the process of forming the core-rim structure of sample C. The core-rim structure of the Ti(C, N) cermet is an important factor that determines its comprehensive performance. The rim phase can greatly improve the wettability between the hard ceramic phase and the metal binder phase and prevent the hard phase from aggregating and growing. The energy spectrum analysis (Figure 12) shows that the chemical composition of the white core phase is close to that of the gray rim phase, and their crystal structures are essentially the same. Therefore, the lattice mismatch of the core-rim structure, which is caused by the different core-rim lattice constants, is small, and the internal stress of the core-rim interface is correspondingly reduced, which results in lower core-rim interfacial stress and stronger bonding at the core-rim interface.Figure 11
Formation process of the core-rim structure.Figure 12
Energy spectra of a white core and gray rim. (a) BSE morphology of core-rim, (b) energy spectrum of (i) white core, and (c) energy spectrum of (ii) gray rim.
(a)
(b)
(c)


### 3.6.2. Rim/Binder Face Interface

According to classical wetting thermodynamics, the wettability of the metal to ceramic phases is the basic criterion that determines the interface bonding between metals and ceramics. The physical quantity that represents the degree of wettability of the metal to the ceramic phase is the contact angle (θ) and adhesion force, as shown in Figure 13(a), which is the famous Young equation:(4)Wad=γB1+cosθ,where γB is the interfacial free energy of the metal phase and Wad is the work of adhesion between the metal phase and the ceramic phase. As can be seen from equation (4), in the case of a certain cosθ value, the larger γB is, the larger Wad is, the greater the adhesion between metal and ceramic is, the better the wettability of metal/ceramic interface is, the more stable the formed interface is, and the better the mechanical properties of the cermets are.Figure 13
(a) Touch angle, (b) rim/binder face of sample C, and (c) the rim/binder face of sample A.Figures13(b) and 13(c) show illustrations of the rim/binder face of sample C and sample A, respectively. As previously discussed in this paper, while the white and gray rims contain the (Ti, W, Mo, Ta) (C, N) solid solution, they have different Ti, W, Mo, and Ta concentrations, and their microstructure is the same as Ti(C, N) and β-Co, which have fcc crystal structures and approximately the same lattice parameters and plane orientations. Therefore, their interface atomic matching degree is high and their lattice mismatch is low, which results in a coherent relationship and the touch angle between the rim/binder phases being very small, as shown in Figure 13(b). According to equation (4), the smaller θ is, the larger is Wad. Therefore, the adhesion force at the rim/binder interface is stronger, and it can improve the wettability between ceramic grains and the binder. In contrast, the touch angle between the rim/binder phases of sample A is large (as shown in Figure 13(c)), which will result in smaller Wad.In addition,β-Co used in this study has a higher interfacial free energy than ordinary Co, so its γB is higher than that of ordinary α-Co. Hence, it has a more significant wetting effect on the ceramic phase, thus improving the mechanical properties of the cermet.


## 3.6.1. Core/Rim Interface

Figure11 illustrates the process of forming the core-rim structure of sample C. The core-rim structure of the Ti(C, N) cermet is an important factor that determines its comprehensive performance. The rim phase can greatly improve the wettability between the hard ceramic phase and the metal binder phase and prevent the hard phase from aggregating and growing. The energy spectrum analysis (Figure 12) shows that the chemical composition of the white core phase is close to that of the gray rim phase, and their crystal structures are essentially the same. Therefore, the lattice mismatch of the core-rim structure, which is caused by the different core-rim lattice constants, is small, and the internal stress of the core-rim interface is correspondingly reduced, which results in lower core-rim interfacial stress and stronger bonding at the core-rim interface.Figure 11
Formation process of the core-rim structure.Figure 12
Energy spectra of a white core and gray rim. (a) BSE morphology of core-rim, (b) energy spectrum of (i) white core, and (c) energy spectrum of (ii) gray rim.
(a)
(b)
(c)


## 3.6.2. Rim/Binder Face Interface

According to classical wetting thermodynamics, the wettability of the metal to ceramic phases is the basic criterion that determines the interface bonding between metals and ceramics. The physical quantity that represents the degree of wettability of the metal to the ceramic phase is the contact angle (θ) and adhesion force, as shown in Figure 13(a), which is the famous Young equation:(4)Wad=γB1+cosθ,where γB is the interfacial free energy of the metal phase and Wad is the work of adhesion between the metal phase and the ceramic phase. As can be seen from equation (4), in the case of a certain cosθ value, the larger γB is, the larger Wad is, the greater the adhesion between metal and ceramic is, the better the wettability of metal/ceramic interface is, the more stable the formed interface is, and the better the mechanical properties of the cermets are.Figure 13
(a) Touch angle, (b) rim/binder face of sample C, and (c) the rim/binder face of sample A.Figures13(b) and 13(c) show illustrations of the rim/binder face of sample C and sample A, respectively. As previously discussed in this paper, while the white and gray rims contain the (Ti, W, Mo, Ta) (C, N) solid solution, they have different Ti, W, Mo, and Ta concentrations, and their microstructure is the same as Ti(C, N) and β-Co, which have fcc crystal structures and approximately the same lattice parameters and plane orientations. Therefore, their interface atomic matching degree is high and their lattice mismatch is low, which results in a coherent relationship and the touch angle between the rim/binder phases being very small, as shown in Figure 13(b). According to equation (4), the smaller θ is, the larger is Wad. Therefore, the adhesion force at the rim/binder interface is stronger, and it can improve the wettability between ceramic grains and the binder. In contrast, the touch angle between the rim/binder phases of sample A is large (as shown in Figure 13(c)), which will result in smaller Wad.In addition,β-Co used in this study has a higher interfacial free energy than ordinary Co, so its γB is higher than that of ordinary α-Co. Hence, it has a more significant wetting effect on the ceramic phase, thus improving the mechanical properties of the cermet.


## 4. Conclusions

(1)
β-Co microspheres obtained by the solid-phase reaction of high-energy ball milling have high activation energies and good wettability, which accelerate solid-phase diffusion, promote atomic diffusion, and accelerate the solid-solution speed during sintering. These greatly improve the wettability of the ceramics phase and increase the solid-solution strengthening effect.
(2)
The (Ti, W, Mo, Ta) (C, N) solid solution can be prepared by the carbothermic reduction-nitridation method (N2 atmosphere protection, heating temperature of 1600°C).
(3)
Cermets were prepared usingβ-Co as the binder phase and using (Ti, W, Mo, Ta) (C, N) solid instead of single metal carbide, which resulted in two kinds of core-rim structures with fine cores and thick rims. With this structure, when the cermets are acted on by an external force, stress transfer is alleviated and crack deflection is improved, which means the toughness of the cermets is improved.
(4)
The properties of the new cermet were compared with those of traditional cermets. The former had a flexural strength (TRS) of 2155 MPa and a fracture toughness (KIC) of 11.79 MPa·m1/2, which are 10% better than those of traditional cermets.



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*Source: 4684529-2020-03-14.xml*