2.1. The influence of BNNTs content
Figure 2 is the XRD pattern of B4C samples fabricated by sintering at 1700 °C with different adding amounts of BNNTs. It can be seen from the Fig. 2 that there were obvious XRD characteristic diffraction peaks of B4C and h-BN in the samples of S-B4C-5wt.% BNNT and S-B4C-10wt.% BNNTs, and no other impurity peaks were found. This shown that the content of amorphous carbon in commercial boron carbide powder was relatively low. For samples Z-B4C-5wt.% BNNTs and Z-B4C-10wt.% BNNTs, there were obvious graphitic carbon peaks in the corresponding XRD spectra. This showed that part of the amorphous carbon contained in the home-made boron carbide powder was completely transformed into crystalline carbon (graphite) under high temperature conditions.
The fracture morphology of ceramic samples with different ceramic matrix and different contents of BNNTs is shown in Fig. 3. It could be seen from Fig. 3 that all the sample sections showed low porosity and high density. It could be seen from Fig. 3(e) that there were some pores of about 0.5 µm and a small amount of white impurity particles (shown in the white wire frame) on the cross-section of C-B4C ceramic. It was also found that the section of the C-B4C sample was flat, and it was speculated that the main fracture mode was transgranular fracture. When adding 5wt.% BNNTs, the C-B4C-5wt.% BNNTs sample appeared some intergranular fractures (Fig. 3(g)), indicating that the addition of a certain amount of BNNTs changed the fracture mode of B4C ceramics. When the content of BNNTs was further increased to 10 wt%, there were basically no pores in the cross section of C-B4C-10wt.% BNNTs, which was basically close to the theoretical density. It was worth noting that the cross-sections of H-B4C and H-B4C-5wt.% BNNTs ceramics showed some grooves formed after the B4C grains were pulled out during the fracture process of the sample (shown in the square wire frame in Figs. 3(a) and 3(b). When the content of BNNTs was too high, the agglomeration between nanotubes was significant (Figs. 3(d) and 3(h)). This kind of agglomeration was equivalent to micron-sized defects, and this loose agglomerate will also produce more void defects at the junction of the nanotube and the matrix, which will hinder the densification of the matrix [14, 19].
It can be seen from Fig. 4(a) that when the content of BNNTs was the same, the B4C-based ceramic composite material with H-B4C nano-powders as the matrix had a higher relative density than the B4C-based ceramic composite material with C-B4C powders as the matrix. The reason may be that the particles of H-B4C nano-powders are much smaller than micron-sized C-B4C powders particles, under the same conditions, the smaller the particle size of the raw material, the more conducive to obtaining high-density ceramic samples. When C-B4C was used as the matrix, the relative density of the B4C-BNNTs ceramic composite material increased accordingly with the increase of BNNTs content. The reason is that because BNNTs have a relatively small particle size, in the ceramic sintering process, BNNTs are easy to fill the gaps between B4C micron grains. However, when H-B4C was used as the matrix, the relative density of B4C-BNNTs ceramic composite material decreased slightly with the increase of BNNTs content [20], and the relative density of the four was close to the theoretical density.
It can be seen from Figs. 4(b) that under the same conditions, H-B4C as a matrix had a higher hardness than C-B4C. The reason may be that the particle size of H-B4C powders are nanometer, while the particle size of C-B4C powders are micrometer. Under the same conditions, the smaller the particle size of the raw material, the more conducive it is to obtain ceramic samples with high density and high hardness. When the matrix was the same, with the increase of BNNTs content, the hardness of the ceramic composite material gradually decreased. The reason may be as the content of BNNTs continued to increase, the possibility of nanotube agglomeration became higher, and the defects and matrix pores introduced by agglomerations will increase, which will reduced the continuity and density of the ceramic matrix, which will eventually lead to the hardness of the ceramic material decreases [21].
It can be seen from Fig. 4(c) that whether the ceramic matrix is C-B4C or H-B4C, as the content of BNNTs increased, the fracture toughness of the composite ceramics first increased and then decreased. When the content of BNNTs was 5wt.%, the fracture toughness of C-B4C-5wt.%BNNTs and H-B4C-5wt.%BNNTs ceramics were both the best, respectively 4.31 Mpa·m1/2 and 5.92 Mpa·m1/2. The results showed that adding an appropriate amount of BNNTs could effectively improve the fracture toughness of B4C ceramics. The reason is that BNNTs are uniformly distributed on the grain boundaries and grains of the B4C matrix. During the crack propagation process, the excellent mechanical properties of the nanotubes can effectively prevent the further propagation of the crack, thereby improving the fracture toughness of the ceramic [14, 22–24]. However, with the further increase of the BNNTs content, the agglomeration of the nanotubes continued to increase, which caused the pores around the nanotubes to increase, which easily induced crack propagation and reduced the toughness of the composite material. Thus, the optimum BNNTs content for B4C composite ceramics in our study is determined to be 5wt.%. In addition, it was found that H-B4C has better mechanical properties than C-B4C under the same conditions.
2.2. The influence of sintering temperature
The mechanical properties of superhard structure ceramics are directly determined by its microstructure, which is affected by the sintering temperature. Figure 5 illustrates the SEM patterns of H-B4C-5wt.% BNNTs samples section at different temperatures. It can be seen from Fig. 5(a) that at low temperatures, there existed more pores and pits in sintered samples, and the powder particles failed to combine with each other to form obvious grain boundaries. With the increase of the sintering temperature, the sample grains tended to fuse, and the grain size increased gradually, the pores reduced and closed, and the density increased. When the sintering temperature reached 1750 °C, the sample was almost completely sintered. The mainly reason is as the sintering temperature increases, the process of surface diffusion and interface diffusion mass transfer speeds up, the density increases, and the pores are continuously eliminated. When the sintering temperature was 1800℃, the sample section was uneven, which may be due to the high temperature, the grain boundary migration rate was greater than the pore migration rate, the grain size increased significantly, and small closed pores were formed inside the grains [15].
It can be seen from Fig. 6 that with the increase of the sintering temperature, the change trend of the microhardness of the H-B4C-5 wt% BNNTs ceramic sample was the same as the change trend of the relative density of the B4C ceramic, indicating that the particle rearrangement of the B4C ceramic mixed with BNNTs was enhanced in the sintering process. With the increase of the sintering temperature, the driving force of B4C sintering continued to increase, and the continuous growth of crystal grains increased the sintering densification of ceramics, which made the microhardness increase, gradually. When the sintering temperature was 1750℃, the microhardness and relative density of the sintered sample were the largest, which are 99.14% and 32.68 GPa, respectively. As the temperature continued to rise, the particle size grew rapidly, and more pore defects were produced, which caused a decrease in the density and hardness of the composite material.
It can be seen from Fig. 6(b) that the trend of the fracture toughness of the composite material with the sintering temperature was similar to that of the microhardness. When the temperature was lower than 1750℃, the fracture toughness of the ceramic continued to increase as the temperature rose. When temperature reached 1750 °C, the fracture toughness was the largest at 6.87 Mpa·m1/2. Thus, the optimum sintering temperature for H-B4C-5wt.% ceramics in our study is determined to be 1750℃.
The reason is that when the sintering temperature increases, the bonding strength of the heterogeneous interface between BNNTs and the B4C ceramic matrix increases. When the crack extends to the surface of the nanotube, the crack propagation path or crack growth energy is increased through crack deflection, bridging and pull-out effects. As the temperature continued to rise, the fracture toughness of ceramic materials decreased significantly. The reason is the B4C grains grow rapidly, and the toughening effect of BNNTs is difficult to offset the abnormal growth of B4C grains and the abnormal interface strength reduction, which leads to a significant reduction in the fracture toughness of the composite material [20].