Microstructure, Mechanical, and Tribological Properties of SiC-AlN-TiB₂ Multiphase Ceramics

Abstract

SiC multiphase ceramics were prepared via spark plasma sintering using AlN and TiB₂ as the second phase and Y₂O₃ as a sintering additive. The effects of TiB₂ content (10 vol.% and 20 vol.%) and sintering temperature (1900 °C to 2100 °C) on the phase composition, microstructure, and mechanical and tribological properties of SiC multiphase ceramics were investigated. The results showed that Y₂O₃ reacts with Al₂O₃ on the surface of AlN to form the intercrystalline phase Y₄Al₂O₉ (YAM), which promotes the densification of the multiphase ceramics. The highest density of SiC multiphase ceramics was achieved at 10 vol.% TiB₂ content. Moreover, TiB₂ and SiC exhibited good interfacial compatibility. In turn, a thin solid-solution layer (~50 nm) was formed by SiC and AlN at the interface. The periodic structure of SiC prevented the dislocation movement and inhibited the base plane slip. The most optimal mechanic characteristics (a density of 98.3%, hardness of 28 GPa, fracture toughness of 5.7 MPa·m^1/2, and bending strength of 553 MPa) were attained at the TiB₂ content of 10 vol.%. The specific wear rates of SiC multiphase ceramics were (4–8) × 10⁻⁵ mm³/N·m at 25 °C and 2.5 × 10⁻⁵ mm³/N·m at 600 °C. The wear mechanism changed from abrasion at 25 °C to a tribo-chemical reaction at 600 °C. Therefore, adding lubricious oxides of TiB₂ is beneficial for the improvement in wear resistance of SiC ceramics at 600 °C.

1. Introduction

SiC ceramics have been widely used because of their high oxidation resistance, excellent thermal conductivity, low thermal expansion coefficient, high hardness and strength, and good tribological properties [1,2,3]. However, due to the binding characteristics of strong covalent bonds, it is difficult to achieve the full density of pure SiC below 2200 °C, which also negatively affects the fracture toughness of the material, thereby limiting its application range to some extent.

Liquid-phase sintered SiC ceramics (LPSiC) enable the reduction of the sintering temperature during the densification process and the obtainment of SiC ceramics with excellent properties [4]. In turn, the rare-earth metal aluminates, formed by the reaction between rare-earth oxides and Al₂O₃, promote the densification of SiC under pressureless conditions [5]. Meanwhile, liquid-phase additives, such as AlN/Al₂O₃, and Y₂O₃, are employed to improve the sintering performance and mechanical properties of SiC ceramics [6,7]. For instance, a sample modified via AlN/Y₂O₃ doping exhibited the lowest weight loss rate and the most optimal bending strength, which were 6% and 433 MPa, respectively [8]. Liang et al. [9] produced LPS-SiC ceramics with Al₂O₃/Y₂O₃ sintering additives, whose coefficient of friction (COF) was in the range of 0.36 to 0.63. Yadav et al. [10,11] studied the tribological and mechanical behaviors of aluminum alloy 6061 composites with the incorporated ceramic particulates by performing a Taguchi analysis. Among them, Al₂O₃- and SiC-doped composites possessed the most prominent properties with a tensile strength and flexural strength of 297 MPa and 397 MPa, respectively.

To improve the fracture toughness and tribological properties of SiC-based ceramics, the second phase is introduced into the SiC matrix. In particular, the reinforcement agents, such as WC, BN, and TiB₂, can noticeably enhance the wear resistance of SiC ceramics [12]. Besides that, the higher thermal expansion coefficient of TiB₂ (8.6 × 10⁻⁶ °C⁻¹) relative to that of SiC (4.2 × 10⁻⁶ °C⁻¹) induces the thermal residual stress field in the SiC-TiB₂ composite, likely weakening the grain and interphase boundaries and hence contributing to crack bridging and deflection [13]. Therefore, hardness and fracture toughness are believed to play an important role in the tribological characteristics of the above ceramics [14].

Chodisetti et al. [15] reported that the addition of TiB₂ increases the hardness and wear resistance of SiC multiphase ceramics. For example, a low wear rate and high COF value (0.47) were successfully achieved in the composite material with 20 vol.% TiB₂ content under a load of 5 N. Zhou et al. [16] concluded that TiB₂ particles impart crack deflection to the SiC matrix and are conducive to superior mechanical and tribological properties. Prakasarao et al. [17] demonstrated that at a TiB₂ content less than 10 vol.%, the friction coefficient at high temperatures dropped to 0.4, while increasing to 0.8 at a TiB₂ amount of 30 vol.%. In spark plasma sintering (SPS), a short sintering time and pressure are expected to limit the grain size and subsequently improve the wear resistance [18]. High-frequency induction sintering (HFIS) is another reliable method to achieve the rapid sintering of metal materials [19]. Meanwhile, the influence of interface reactions between SiC and second phases on the tribological and wear properties of SiC-TiB₂ ceramic composites has still been rarely reported.

In this paper, SiC-AlN-TiB₂ multiphase ceramics were prepared with Y₂O₃ (1 wt.%) as a sintering aid. Utilizing SPS, the effects of TiB₂ content and sintering temperature on the phase, microstructure, and mechanical and tribological properties of SiC multiphase ceramics were carefully studied.

2. Experimental

2.1. Sample Preparation

The raw materials used in this paper were SiC (d₅₀, 0.6 μm; purity, 99.9%; Shanghai Yunfu Nano Technology Co., Ltd., Shanghai, China), AlN (d₅₀, 0.6 μm; purity, 99.9%; S H.C. STARCK. Goslar, Germany), TiB₂ (d₅₀, 0.6 μm; purity, 99.5%; Shanghai Yunfu Nano Technology Co., Ltd., Shanghai, China), and Y₂O₃ (d₅₀, 0.6 μm; purity, 99.9%; Shanghai Yunfu Nano Technology Co., Ltd., Shanghai, China) [20]. The raw SiC contained 6H-SiC, 4H-SiC, and a small amount of 3C-SiC.

The amount of SiC was adjusted so that the volume fractions of the TiB₂ particles were 10 vol.% and 20 vol.%, while that of AlN was 10 vol.%. The sintering aid was Y₂O₃ (1 wt.%). The powders were ball-milled in absolute ethyl alcohol for 10 h. The mass ratio of the ball to powder was 5:1. The slurry was separated from the mill balls through a screen and then dried in an oven at 100 °C for 24 h. The obtained powder was afterward loaded into a graphite mold. The samples were sintered in an SPS furnace (SPS-4, Shanghai Chenhua Technology Co., Ltd. Shanghai, China), conforming to the following heating modes. The temperature was first raised from 25 °C to 650 °C at a rate of 135 °C/min, and pressure was subsequently applied. The temperature afterward rose to 1600 °C at a rate of 125 °C/min, whereas the pressure was increased to 40 MPa. The heating proceeded until the highest temperature (1900 °C, 1950 °C, 2000 °C, 2050 °C, and 2100 °C) was attained at a rate of 50 °C/min. The ultimately achieved temperature along with the pressure was held for 10 min, after which it was decreased to 1000 °C at a rate of 140 °C/min. The furnace with the sample inside was afterward left cooling naturally to room temperature. According to the composition and sintering temperature, they were designated as two series, S1 (10 vol.% TiB₂) and S2 (20 vol.% TiB₂). For example, specimens were labeled as S1-2000 (10 vol.% TiB₂; 2000 °C).

2.2. Testing and Characterization

The relative densities of the samples were measured using Archimedes’ method. The phase compositions were determined via X-ray diffraction (XRD, XRX-6000, Shimazu, Japan, 40 kV and 30 mA) using Cu Kα radiation in the 2θ range of 10–80° (scan rate = 2°/min). The morphology of the samples was examined by transmission electron microscopy (TEM) and field-emission scanning electron microscopy (FESEM, Sigmas, Zeiss, Germany), coupled with energy-dispersive spectroscopy (EDS, Oxford, UK). The morphology and atomic arrangement at the grain boundaries and inside the grains were characterized by high-resolution transmission electron microscopy (HRTEM) and selected area electron diffraction (SAED). A universal material testing machine (WDW-600KN, Shenzhen New SANS, Shenzhen, China) was used to measure the bending strength of the specimens according to a three-point bending method. Prior to the tests, the samples were machined to achieve the dimensions of 3 mm × 4 mm × 20 mm and ground using diamond polishing liquid with a grain size of 0.5 μm. The loading speed was kept at 0.2 mm/min. The Vickers hardness of ceramics was assessed by means of a Vickers hardness machine (432VD, Changzhou Taylor, Changzhou, China) under a load of 10 N applied for 10 s. Seven measurements were performed on each sample, and the average value was taken as the final result. The fracture toughness (K_1C) was established conforming to the indentation fracture method proposed by Anstis [21]. The friction experiments were carried out using a high-temperature friction and wear testing machine (HT-1000, Lanzhou Zhongke Kaihua Instrument Co., Ltd., Lanzhou, China) at a load and sliding speed of 15 N and 0.1 m/s, respectively. A cemented carbide (WC, Co bong) ball with a diameter of 5 mm served as the frictional counterpart. The frictional temperatures were 25 °C and 600 °C, the rotation radius was 3 mm, and the sliding velocity and duration were set to 0.1 m/s and 180 min, respectively. Three measurements were performed on each sample, and the average value was taken as the final result. The CoF values were recorded with the testing machine in real-time mode. After the test was finished, the sample was cleaned ultrasonically in anhydrous ethanol. The mass of the sample was weighed before and after the friction test, and the specific wear rate was calculated according to the equation below [22]:

$$W={\displaystyle \frac{V}{F\cdot L}}$$

(1)

where V is the wear volumes obtained by the weight loss and measured density of the SiC ceramics, F (N) is the applied load, and L (m) is the sliding distance.

3. Results and Discussion

3.1. Phase Composition

Figure 1 depicts the XRD patterns of SiC multiphase ceramics. As can be seen from Figure 1a, all the sintered samples were mainly composed of α-SiC. There were also AlN, TiB₂ (PDF#35-0741), and C (PDF#99-0057) phases. For S1-2000, the full width at half-maximum of the relevant XRD peaks was narrower than that of S2-2000. This indicated that the as-obtained SiC multiphase ceramics possessed a relatively high crystallinity degree [23]. In Figure 1b, the XRD peaks from the S1-1900 to S1-2050 specimens were similar to those of S1-2000. However, the diffraction features of S1-2100 showed an obvious broadening in their half-height width. This may be due to the increase in the degree of solid solution between AlN and SiC with the increase in temperature [24].

The diffusion coefficients of AlN in SiC are 25 times as many as those of SiC in AlN, and there is more AlN than SiC at the grain interface [25]. According to An et al. [23], the lattice constant of AlN increases after AlN and SiC form a solid solution. Thus, solid solutions of AlN and SiC formed in the S1-2100 sample. According to the XRD data, the SiC powder in this paper predominately had an α-polytype structure (6H-SiC) whose theoretical lattice constants are a = 3.073 Å and c = 15.079 Å (PDF#75-1664). In turn, the AlN material had a cubic structure with a = 4.342 Å (PDF#88-2363). In addition, for the cubic SiC phase, a = 4.349 Å (PDF#73-1665). The lattice constants of SiC (PDF#73-1663), calculated from the XRD profiles in Figure 1a, were found to be a = 3.079 Å and c = 15.061 Å, while those for the AlN (PDF#80-0010) and cubic SiC (PDF#29-1129) phases were a = 4.365 Å and a= 4.359 Å, respectively. Therefore, the lattice constants of SiC changed to a small extent compared to the theoretical values, while those of AlN became larger.

3.2. Microstructure

3.2.1. FESEM Data

Table 1. Density and mechanical properties of SiC multiphase ceramics.

Sample	Measured Density /g/cm3	Relative Density /%	Vickers’ Hardness /GPa	Fracture Toughness /MPa·m1/2	Bending Strength /MPa
S2-1950	3.16	91.06	21.9 ± 0.9	-	-
S2-2000	3.31	95.38	25.3 ± 1.0	5.13 ± 0.4	-
S2-2050	3.31	95.50	26.8 ± 1.6	5.12 ± 0.6	-
S2-2100	3.32	95.56	27.2 ± 1.1	4.95 ± 0.6	-
S1-1900	3.31	95.28	26.7 ± 1.4	5.10 ± 0.4	438 ± 22
S1-1950	3.34	96.11	26.8 ± 0.7	5.39 ± 0.6	450 ± 54
S1-2000	3.41	98.27	28.3 ± 1.6	5.74 ± 0.5	553 ± 22
S1-2050	3.32	95.53	27.0 ± 1.2	4.95 ± 0.3	444 ± 43
S1-2100	3.34	96.56	27.1 ± 1.6	4.98 ± 1.2	467 ± 29

summarizes the relative densities of SiC multiphase ceramics. As seen from the data, the relative density of S1 varied in a non-monotonic manner with temperature, reaching a maximum value of 98.3% at 2000 °C. With 20 vol.% TiB₂, the relative density of S2 gradually increased. The highest relative density (95.6%) was achieved in the S2-2100 specimen. Thus, the relative density of S1 exceeded that of the S2 ceramics at any given temperature.

Figure 2 depicts the SEM images and elemental mappings of the S1-2000 and S2-2000 samples. Three different regions were distinguished within the polished surfaces, among which the continuous “gray” particles represented the matrix, while the “gray–black” and “white” particles were distributed dispersedly in the “gray” one. Based on the EDS results in Figure 2a, the “gray”, “gray–black”, and “white” particles were ascribed to SiC, AlN, and TiB₂, respectively. Thus, TiB₂ and AlN were secondary phases. It is noteworthy that the black part was the pit left by the release of TiB₂ particles. Moreover, there was a close connection between the SiC matrix and AlN. However, since no solid solution was detected in the SEM images, TEM analyses were further required. In Figure 2b, the gray–black phase was also observed in sample S2. By analyzing the Al/Si/C/N/Ti mass fraction at point A (Figure 2c), it was concluded that Al (5%) and Si (36%) were the main elements in this region, indicating that SiC and AlN were dissolved mutually [26]. Besides that, the intercrystalline phase left by the sintering of liquid-phase additives was distributed throughout the small white linear area between the crystals.

Furthermore, almost no pores could be observed on the polished surface of S1-2000. The second phases were tightly bound, and the growth of the SiC grains was restricted (Figure 2a). Meanwhile, abundant micropores were observed in S2-2000 (Figure 2b). This was because of weak links between SiC and the second phase, resulting in the formation of numerous defects and the incomplete densification of S2-2000, which is consistent with the relative density data from Table 1. Therefore, the increase in TiB₂ amounts in SiC multiphase ceramics was not conducive to densification.

Figure 3 displays the fractured surface of S1 samples. According to Figure 3a–e, the S1 specimens mainly underwent transgranular fracture with cleavage steps, though there were pull-outs of grains.

3.2.2. TEM Data

To investigate the microstructure of SiC, TiB₂, and AlN grains, as well as the interfacial reaction between each phase, the formation of a solid solution of AlN and SiC, and the change of SiC at high temperatures, a TEM analysis was performed on the sample S1-2000.

Figure 4a,d depict the bright-field TEM images of SiC grains, whereas Figure 4b,c,e,f display the SAED patterns and HRTEM images of areas A and B, respectively. It can be seen that there were many stacking faults within the SiC grains from areas A and B (Figure 4a,b), but they were not observed in AlN and TiB₂ grains. Figure 4b depicts the relatively weak diffraction spots corresponding to SiC (0 0 4) (4H PDF: 73-1664) and the transmission spot, which were distributed at the n/4 positions, indicating that the SiC grains within area A had a four-period stacking structure. Moreover, there were five weak reflections between the strong diffraction spots (1 0 5) corresponding to SiC (6H PDF: 72-0018), and the spacing was divided into six equal fractions (Figure 4e), indicating that the SiC grains within area B had a 6H long-period ordered stacking (LPSO) structure [27]. The HRTEM data revealed the structural characteristics of the stack, where the partial dislocation and the absence of the atomic arrangement in the incident electron beam direction led to a weak image contrast (Figure 4c,f). Moreover, the SiC possessed a periodic structure, which inhibited dislocation movement and a base plane slip, improving the properties of the SiC ceramics [28].

Figure 5 depicts the TEM images and SAED patterns of TiB₂. According to the data, TiB₂ particles presented an elliptical shape mainly with a crystal band axis of [0, 1, $\overline{1}$]. The fringe spacings (Figure 5c) on the two sides of the interface were 0.2633 nm and 0.2582 nm, which corresponded to the TiB₂ (0 1 0) and SiC (1 1 1) crystal planes, respectively. The degree of convergence of the lattice (F), or the mismatch of lattices at the interface between TiB₂ and SiC, was calculated using the formula below [29]:

$$F=2({\mathrm{a}}_2-{\mathrm{a}}_1)/({\mathrm{a}}_2+{\mathrm{a}}_1)$$

(2)

where a₂ (TiB₂) = 0.2633 nm and a₁ (SiC) = 0.2582 nm. The corresponding F value at the TiB₂/SiC interface was found to be 1.95%. The results showed that the TiB₂ and SiC phases formed a coherent and highly stable interface with low strain energy [30], free of impurity phases. Therefore, the degree of lattice matching between TiB₂ and SiC endowed them with favorable interface compatibility, which was conducive to an improvement in the mechanical properties of the multiphase ceramics.

Figure 6 shows the TEM images, elemental mappings, and SAED patterns of the intercrystalline phase in the S1-2000 specimen. In Figure 6a–d, the abundant Al, N, Y, and O elements exist at the grain interfaces of the matrix and in the triangular region. According to the HRTEM image at P2 (Figure 6a,e), the intercrystalline phase possessed a high degree of structural ordering along with high crystallinity. Combined with Figure 6a–e, the formation of the intercrystalline Y₄Al₂O₉ (YAM) phase was thus confirmed. Moreover, no amorphous structure was observed at the interface between the adjacent SiC grains.

Figure 7 depicts the distribution of Si, Al, C, and N elements in the sample. In particular, the SiC matrix was embraced by a thin layer with Al, N, and Y elements. The thickness of the solid-solution layer was approximately 30–70 nm. The SiC-AlN core-shell structure indicated that Al and N were partially dissolved in the marginal region of SiC grains, whereby SiC and AlN formed a disordered solid solution. Therefore, there were only slight changes in the lattice constants of AlN and SiC, revealed by the XRD data (Figure 1). The boundary energies of SiC and AlN were reduced by the formation of solid solutions, resulting in the promotion of the displacement and migration of Si, C, Al, and N atoms in the lattice, and improvement in the sintering properties of the materials [14,24].

3.3. Mechanical Properties

Table 1 displays the mechanical properties of the SiC multiphase ceramics. The hardness of S2 increased with temperature, achieving a maximum of 27.2 GPa at 2100 °C (S2-2100). In turn, the hardness of the S1 multiphase ceramics increased from 26.7 to 28.3 GPa, reaching the maximum at 2000 °C. Therefore, at any given temperature, the hardness decreased with the increasing content of TiB₂. At the same time, the variation of bending strength corroborated with the hardness of S1 multiphase composites. Based on Table 1, the bending strength increased from 438 to 553 MPa in S1. In turn, the fracture toughness of the S2 samples was approximately 5 MPa·m^1/2 and remained almost unchanged with temperature. On the contrary, the fracture toughness of the S1 specimens varied in a non-monotonic manner, enriching the maximum value of 5.7 MPa·m^1/2 for S1-2000. However, the fracture toughness of the S1-2050 and S1-2100 samples tended to be stable.

Relative density plays a decisive role in mechanical properties of multiphase ceramics. In LPS ceramics, an increase in temperature is effective to promote densification. However, excessive temperatures lead to the volatilization of the liquid phase, increasing the internal defects of ceramics and inhibiting densification [31]. Meanwhile, pores promote stress concentration and abnormal crystal growth. Therefore, the highest relative density of the S1-2000 specimen was the main reason for the increase in mechanical properties in this paper.

Figure 8a–e depict the three propagation modes of the crack, including deflection, bridging, and transgranular cracking, through TiB₂ particles. The contribution of SiC particles to crack deflection and bridging was small. In turn, TiB₂ particles were dispersed in the matrix, which increased the probability of their contact with TiB₂ grains during crack propagation, as well as the probability of crack deflection. The extension of the crack propagation path considerably improves the fracture toughness of SiC multiphase ceramics [32]. Based on Ref. [14], the most fundamental way to enhance the fracture toughness of ceramics is to hinder the crack growth. In this paper, abundant tortuous crack paths and crack bridges were observed in the S1-2000 material. Therefore, it can be concluded that the above specimen possessed the highest fracture toughness.

3.4. Tribological Properties

3.4.1. CoF and WRs

Table 2. Average CoF and WRs of S1 multiphase ceramics at 25 °C.

Sample	Average CoF	Average WRs/(mm3/N·m)
S1-1900	0.45	3.62 × 10−5
S1-1950	0.46	6.63 × 10−5
S1-2000	0.51	6.81 × 10−5
S1-2050	0.42	7.18 × 10−5
S1-2100	0.47	8.02 × 10−5

depicts the average coefficients of friction (CoF) and specific wear rates (WRs) of S1 multiphase ceramics at 25 °C. In turn,

Table 3. Average CoF and WRs of S1-2000 multiphase ceramics at 600 °C.

Sample	Average CoF	Average WRs/(mm3/N·m)
S1-2000	0.47	2.46 × 10−5

displays the CoF and WRs of S1-2000 at 600 °C. As can be seen from Table 2, the average CoF was basically maintained between 0.4 and 0.5. With an increase in sintering temperature, the average WRs of S1 multiphase ceramics gradually increased from 3.62 × 10⁻⁵ to 8.02 × 10⁻⁵ mm³/N·m.

According to [33], the CoF of LPS-SiC ceramics (PL) was between 0.37 and 0.63, and the wear rate was (10⁻⁵–10⁻⁴) mm³/N·m. Therefore, when cemented carbide was used as a counterpart, SiC multiphase ceramics exhibited an intermediate CoF and a low WR. This indicated that the SiC multiphase ceramics containing 10 vol.% TiB₂ had better friction and wear characteristics.

Figure 9 shows the CoF of S1-2000, measured in real-time mode. At 25 °C, the run-in period was short, and the CoF in the stable stage slightly fluctuated. The average CoF value was 0.51. At 600 °C, the run-in period of S1-2000 was longer, which might have been caused by the formation of debris during the initial sliding process and the change of surface roughness within the run-in period. The CoF of the stability stage was 0.47 (Table 3), which was comparable to that at 25 °C. Meanwhile, the WR was 2.46 × 10⁻⁵ mm³/Nm, being only 36% of that at 25 °C. The reduction in the WR might have been related to tribo-oxidation films on the worn surface.

3.4.2. Morphology of Worn Surface

Figure 10 depicts the typical SEM images of the wear tracks of S1-2000. At 25 °C, there were numerous grooves inside the wear tracks (Figure 10a), as well as traces of fractures and pull-outs of the grains at the edges. The liquid YAM phase between the SiC grains had a relatively low strength and was easy to break under the action of friction [34]. Once the liquid phase had broken, the SiC grains could be peeled off. Since the flaking SiC fragments were abrasive, grooves successfully formed on the surface.

At 600 °C, there were compacted layers of debris along with grooves. The wear track showed a micro-fracture (Figure 10b) accompanied by grain shedding. This was similar to the grain pull-out observed during the unidirectional sliding of SiC-TiB₂ multiphase ceramics in Ref. [35]. According to Figure 10b-(O), the worn surface suffered from oxidation.

3.4.3. Tribo-Oxides on the Worn Surface

Figure 11 depicts the binding energies of elements on the worn tracks at 25 °C and 600 °C. Compared to the results acquired at 25 °C, there were more oxidized compounds on the worn surface. Figure 11a shows the bimodal Si2p peaks ascribed to Si-C (101 eV) and Si-O (102.6 eV) bonding [36]. In Figure 11b, the bimodal Ti2p peaks were attributed to TiB₂ (453.9 eV) and TiO₂ (458.5–458.7 eV, 464.0–464.4 eV) [37]. In Figure 11c, the C1s peak at 282 eV was associated with the C-Si bond, while that at 284 eV was due to contaminants present in the specimen [38]. In Figure 11h, the W4f peak of 35.8 eV at 600 °C was attributed to tungsten oxides [39], which were transferred from the counterpart ball of the cemented carbide (WC, Co-bonded). Thus, it could be concluded that at 600 °C, there were TiO₂, silicon oxides, and tungsten oxides on the worn surface, which confirmed the tribo-oxidized state of the latter. Moreover, there was also the transfer of substance from duality. According to Skopp et al. [40], TiB₂ grains are more easily oxidized than SiC grains. Moreover, lubricious TiO₂ and B₂O₃ phases were generated via the tribo-oxidation of TiB₂ during rubbing. Compared to the data obtained at 25 °C, TiB₂ grains were oxidized more severely at 600 °C (Figure 11b). In this paper, it was also observed that the degree of tribo-oxidation of TiB₂ at 600 °C was more severe than that at 25 °C.

3.4.4. Wear Mechanism

As shown in Figure 10, the morphology of worn surfaces at 25 °C and 600 °C exhibited different features. Grooves, pits of fracture, and pull-outs of grains were predominant at 25 °C, which indicated the abrasion wear mechanism of the sample surface. By contrast, there were compacted tribo-layers (Figure 10b) composed of tribo-oxides (Figure 11) as well as grooves on the worn surface at 600 °C, indicating the mixed wear mechanisms of abrasion and tribo-chemical reaction. In a word, lubricious oxides of TiB₂ endowed the ceramic with wear resistance at 600 °C.

4. Conclusions

Dense SiC-AlN-TiB₂ multiphase ceramics were prepared via SPS using Y₂O₃ as the sintering additive. The main conclusions can be drawn as follows.

• AlN particles were diffused throughout the SiC surface to form a solid solution layer (~50 nm), which effectively reduced the grain boundary energy and improved the sintering performance. Y₂O₃ reacted with Al₂O₃ on the surface of AlN to form the intercrystalline phase YAM, which promoted the densification of multiphase ceramics. TiB₂ and SiC exhibited good interfacial compatibility.
• The maximum mechanical characteristics (a density of 98.3%, hardness of 28 GPa, fracture toughness of 5.7 MPa·m^1/2, and bending strength of 553 MPa) were achieved in the S1-2000 specimen at a TiB₂ content of 10 vol.%.
• The average CoF values of SiC multiphase ceramics at 25 °C and 600 °C were mainly in the range of 0.4–0.5. The WRs were in the order of 10⁻⁵ mm³/N·m at 25 °C, decreasing by 36% at 600 °C. The wear mechanism changed from abrasion at 25 °C to a tribo-chemical reaction at 600 °C. The lubricious oxides of TiB₂ enabled the improvement of the wear resistance of SiC ceramics at 600 °C.