Improvement of Mechanical Properties of Composites with Surface Modified B4C for Precision Machining

In order to solve the problem of difficult sintering and high brittleness of B4C-based ceramics, B4C@ZrB2-TiB2 composite powder was synthesized by molten salt method, and B4C–(Zr, Ti)B2 composite ceramics were successfully prepared by spark plasma sintering. The effects of different raw material ratios on the composition, microstructure, and mechanical properties of the prepared composite ceramics were characterized by XRD, XPS, SEM, and TEM. The results show that ZrB2 and TiB2 were grown on the surface of B4C by template mechanism to form a dense nanocrystalline coating, and the original surface of B4C was exposed gradually with the decrease of the ratio of metal powder. When the composite powders were sintered at 1700 °C, ZrB2 and TiB2 formed a solid solution, which can refine grains and improve strength. When the raw material ratio is n(B4C): n(Zr): n(Ti) = 12:1:1, the composite ceramics have excellent comprehensive properties, the Vickers hardness reaches 41.2 GPa.


Introduction
Given the higher grinding requirements of rough wafers grinding in the integrated circuit field, it is urgent to develop ultra-hard ceramics with excellent mechanical properties to meet the new growth requirements. The hardness of boron carbide ceramics is third only to diamond and cubic boron nitride, and the lower cost makes it have great application potential. Boron carbide has the characteristics of high hardness, high melting point, high modulus of elasticity, low thermal expansion coefficient and good chemical stability [1,2], and is widely used in mechanical equipment, abrasive abrasives, microwave absorption, and refractory antioxidants and other fields [3][4][5]. Zamora et al. [6] prepared dense boron carbide composites at low temperatures and they showed good wear resistance with only slight wear after linear sliding over long distances. However, the lower sintering temperature and the introduction of more low-hardness phases make it difficult to meet the grinding needs of rough wafers. The mechanisms of porosity elimination, grain boundaries, and volume diffusion of ceramics only play a role when the sintering temperature reaches more than 2000 • C This ultimately makes it difficult to achieve the sintering densification of B 4 C ceramics with poor plasticity (K IC is about 2 MPa·m 1/2 ), and its ultra-hard properties are difficult to fully achieve [7,8].
In order to achieve the densification and sintering of B 4 C-based ceramics, various sintering aids and second phases were introduced to reduce the sintering temperature. It has been found that carbides and borides such as SiC, TiC, HfC, TaB 2 , TiB 2 , and ZrB 2 benefit the performance improvement of composite ceramics [9][10][11]. Chen et al. [12] found that ZrB 2 -ZrC-B 4 C composites have been successfully fabricated by SPRS from B 4 C and different content of Zr, while the flexure strength and hardness decrease first and then increase with increasing Zr content. It is found that boron carbide containing 66% Zr has the highest fracture toughness (5.83 MPa·m 1/2 ) and flexure strength (386.45 MPa) due Materials 2023, 16, 882 2 of 10 to the uniform distribution of ZrB 2 and ZrC. Sha et al. [13] explored the effect of carbon content on the mechanical properties of reaction-sintered B 4 C composite ceramics. When the carbon content is 10%, the composite ceramics have the highest bending strength (444 MPa) and elastic modulus (329 GPa). After adding an appropriate proportion of carbon and silicon in the preparation of B 4 C/TiB 2 composite ceramics, Zhu et al. [14] found that layered graphite appeared at the grain boundary of B 4 C and it could react with B 2 O 3 to cause volume shrinkage, which significantly improved the mechanical properties of the material. Yan et al. [15] introduced TiSiC 2 additive to prepare B 4 C-TiB 2 composite ceramics. Compared with the pure B 4 C ceramics prepared by the same method, the hardness of the composite ceramics decreased, but the flexural strength and fracture toughness increased significantly. Ren et al. [16,17] used NaCl-KCl as a molten salt medium to prepare the composite powder with Al 3 BC and TiB 2 coating uniformly on the surface of B 4 C particles. Due to the conductivity of boride coating, the material transportation is dynamically enhanced through SPS sintering.
Spark plasma sintering (SPS) is a new rapid densification technology that uses highfrequency and high currents through the sample to heat the sample and apply mechanical pressure [18][19][20][21]. Yavas et al. [22] found that SPS sintering has higher surface energy than B 4 C powder with small particle size (HS) at heating rate of 75, 150 and 225 • C·min −1 , which improves the sintering driving force and enables the sintering of B 4 C ceramics to be completed at 1590 • C. Moshtaghioun et al. [23] combined high-energy ball milling, annealing treatment and SPS to provide sinterability for ultrafine B 4 C powder, and the optimal sintering conditions were 100 • C·min −1 temperature to 1700 • C for 3 min, which could prepare B 4 C ceramics with ultra-fine grains with a relative density of more than 98.5%. At the same time, due to the significant decrease of grain size and the increase of transcrystalline fracture mode, the ceramic has an ultra-high hardness of 38 GPa without corresponding ductile loss (~3 MPa·m 1/2 ). Experiments show that the two-step SPS treatment enables B 4 C powder to be densified while retaining nanoscale particles, with ultra-high hardness and good toughness [24].
Although many studies have improved the performance of B 4 C ceramics to varying degrees, there are still problems such as high sintering temperature, poor fracture toughness and lower hardness. In this paper, B 4 C-(Zr, Ti)B 2 composite ceramic with high hardness was prepared using the SPS and using (Zr, Ti)B 2 -coated B 4 C powder prepared by molten salt as raw materials. The effects of the ratio of raw materials on the phase composition and microstructure of composite powders and ceramics were discussed, and the hardness of the ceramics was compared.

Materials
The composite powder for preparing composite ceramics was synthesized by molten salt method using B 4 C (Mudanjiang Diamond Co., Ltd., W1.5, Mudanjiang, China), ZrH 2 (ST-NANO, 0.5 µm, Shanghai, China), TiH 2 (ST-NANO, 0.5 µm, Shanghai, China) as raw materials, and NaCl (Sinopharm Chemical ReagentCo., Ltd, AR, Shanghai, China), KCl (Macklin, AR, Shanghai, China) as a molten salt medium. First, mixed the raw materials in the mortar for about 25 min (according to the mass ratio of m (mixed salt):m (mixed powders) = 7:3 and m (NaCl):m (KCl) = 3:2. ZrH 2 and TiH 2 will decompose and precipitate H 2 at low temperatures to obtain metal Zr and Ti [25]. The actual reaction occurs between Zr/Ti and B 4 C. Due to the extremely low H content and the low addition of raw materials in each group of mixed powders, the ratio was calculated according to the molar mass of Zr and Ti. The molar ratio of B 4 C to Zr/Ti are 8:1:1, 12:1:1 and 16:1:1, named ZT8, ZT12, and ZT16, respectively. Put the mixture into the alumina crucible with a lid, then the crucible containing the mixture was placed in an atmosphere furnace and held to 1100 • C for 2 h under Ar atmosphere. The solid obtained after heat treatment was washed with deionized water after ultrasonic cleaning of 15 min in a water bath at 45 • C (these operations should be repeated three times). Lastly, put the products at 110 • C and dried for 12 h to obtain composite powder.
Accurately weighed composite powder was placed in a special cylindrical graphite crucible (Φ20 mm), and the crucible containing the powder was placed in an SPS sintering furnace for sintering (the inner wall of the crucible, and the composite powder were isolated with clean graphite paper, and the powder was compacted and sealed with a graphite plunger at the upper and lower ports). The sintering system of this experiment was: holding at 1400 • C for 3 min, holding at 1700 • C for 6 min, the heating rate was 100 • C/min, and the sintering pressure was 50 MPa. Pressure was increased gradually during heating and reached maximum pressure before maximum temperature. After the heat preservation, the pressure was gradually removed and the samples demoulded.
After removing the graphite paper from the ceramic samples, the ceramic samples were ground, polished, then characterized and tested for the performance of the samples.

Characterization and Testing
The phase composition and microstructure of the composite powders and the composite ceramics were characterized using the X-ray diffractometer (XRD, X'Pert Pro, Philips, The Netherlands), field emission scanning electron microscope (SEM, Nova nano 400, FEI, Hillsboro, OR, USA) equipped with an energy dispersive spectrometer (EDS, IE350 Penta FET X-3, Oxford, UK) and high resolution transmission electron microscope (TEM, JEM2100, JEOL, Tokyo, Japan). The element composition and chemical bond bonding state on the surface of the composite powder was analyzed by an X-ray photoelectron spectrometer (XPS, AXIS SUPRA+, Shimadzu-Kratos, Hadano, Kanagawa, Japan). The Vickers hardness of the ceramic samples was measured by an automatic micro/macro hardness tester, the applied load was 0.5 kg and the loading time was 15 s. The indentation diagram is shown in Figure 1. 2 h under Ar atmosphere. The solid obtained after heat treatment was washed with deionized water after ultrasonic cleaning of 15 min in a water bath at 45 °C (these operations should be repeated three times). Lastly, put the products at 110 °C and dried for 12 h to obtain composite powder. Accurately weighed composite powder was placed in a special cylindrical graphite crucible (Φ20 mm), and the crucible containing the powder was placed in an SPS sintering furnace for sintering (the inner wall of the crucible, and the composite powder were isolated with clean graphite paper, and the powder was compacted and sealed with a graphite plunger at the upper and lower ports). The sintering system of this experiment was: holding at 1400 ℃ for 3 min, holding at 1700 °C for 6 min, the heating rate was 100 °C/min, and the sintering pressure was 50 MPa. Pressure was increased gradually during heating and reached maximum pressure before maximum temperature. After the heat preservation, the pressure was gradually removed and the samples demoulded.
After removing the graphite paper from the ceramic samples, the ceramic samples were ground, polished, then characterized and tested for the performance of the samples.

Characterization and Testing
The phase composition and microstructure of the composite powders and the composite ceramics were characterized using the X-ray diffractometer (XRD, X'Pert Pro, Philips, The Netherlands), field emission scanning electron microscope (SEM, Nova nano 400, FEI, Hillsboro, OR, USA) equipped with an energy dispersive spectrometer (EDS, IE350 Penta FET X-3, Oxford, UK) and high resolution transmission electron microscope (TEM, JEM2100, JEOL, Tokyo, Japan). The element composition and chemical bond bonding state on the surface of the composite powder was analyzed by an X-ray photoelectron spectrometer (XPS, AXIS SUPRA+, Shimadzu-Kratos, Hadano, Kanagawa, Japan). The Vickers hardness of the ceramic samples was measured by an automatic micro/macro hardness tester, the applied load was 0.5 kg and the loading time was 15 s. The indentation diagram is shown in Figure 1.  Figure 2 shows the XRD patterns of the composite powder prepared at 1100 °C. It can be seen from Figure 2 that the main phase of the composite powder is composed of B4C, ZrB2, and a small amount of ZrO2. The oxidation of metal zirconium powder and the reaction of B2O3 on the surface of B4C with Zr powder may cause the diffraction peak of ZrO2. Meanwhile, only a weak diffraction peak of TiB2 was detected in the composite powder, which might be the dissolution of TiB2 in ZrB2 [26], and the significant atomic number of Zr leads to the reduction of X-ray scattering factor of TiB2 [27].  Figure 2 shows the XRD patterns of the composite powder prepared at 1100 • C. It can be seen from Figure 2 that the main phase of the composite powder is composed of B 4 C, ZrB 2 , and a small amount of ZrO 2 . The oxidation of metal zirconium powder and the reaction of B 2 O 3 on the surface of B 4 C with Zr powder may cause the diffraction peak of ZrO 2 . Meanwhile, only a weak diffraction peak of TiB 2 was detected in the composite powder, which might be the dissolution of TiB 2 in ZrB 2 [26], and the significant atomic number of Zr leads to the reduction of X-ray scattering factor of TiB 2 [27]. Figure 3 shows the SEM images of the composite powder prepared at 1100 • C. Figure 3a,b show the morphology of the ZT8 composite powder and its elemental mapping analysis. On the one hand, it confirms the generation of TiB 2 and ZrB 2 on the surface of B 4 C particles after heat treatment at 1100 • C in molten salt. On the other hand, it illustrates the homogeneous encapsulation of the two phases on the B 4 C particles. Figure 3c,d show the ZT12 and ZT16 composite powder morphology. The composite powder particles maintain the original shape of the irregular polyhedra of the B 4 C particles, and the original surface is gradually exposed with a decrease in the proportion of metal powders.  Figure 3 shows the SEM images of the composite powder prepared at 1100 °C. Figure  3a,b show the morphology of the ZT8 composite powder and its elemental mapping analysis. On the one hand, it confirms the generation of TiB2 and ZrB2 on the surface of B4C particles after heat treatment at 1100 °C in molten salt. On the other hand, it illustrates the homogeneous encapsulation of the two phases on the B4C particles. Figure 3c,d show the ZT12 and ZT16 composite powder morphology. The composite powder particles maintain the original shape of the irregular polyhedra of the B4C particles, and the original surface is gradually exposed with a decrease in the proportion of metal powders. To further verify the formation of ZrB2 and TiB2 on the surface of B4C particles, the chemical composition of ZT8 composite powder was analyzed by XPS. Figure 4a shows the pattern of B1s, which is divided into two peaks at 187.45 eV and 192.55 eV. Since the energies of the B-Zr bond [28,29] of ZrB2 and the B-Ti bond [30] of TiB2 are similar, the   Figure 3 shows the SEM images of the composite powder prepared at 1100 °C. Figure  3a,b show the morphology of the ZT8 composite powder and its elemental mapping analysis. On the one hand, it confirms the generation of TiB2 and ZrB2 on the surface of B4C particles after heat treatment at 1100 °C in molten salt. On the other hand, it illustrates the homogeneous encapsulation of the two phases on the B4C particles. Figure 3c,d show the ZT12 and ZT16 composite powder morphology. The composite powder particles maintain the original shape of the irregular polyhedra of the B4C particles, and the original surface is gradually exposed with a decrease in the proportion of metal powders. To further verify the formation of ZrB2 and TiB2 on the surface of B4C particles, the chemical composition of ZT8 composite powder was analyzed by XPS. Figure 4a shows the pattern of B1s, which is divided into two peaks at 187.45 eV and 192.55 eV. Since the energies of the B-Zr bond [28,29] of ZrB2 and the B-Ti bond [30] of TiB2 are similar, the To further verify the formation of ZrB 2 and TiB 2 on the surface of B 4 C particles, the chemical composition of ZT8 composite powder was analyzed by XPS. Figure 4a shows the pattern of B1s, which is divided into two peaks at 187.45 eV and 192.55 eV. Since the energies of the B-Zr bond [28,29] of ZrB 2 and the B-Ti bond [30] of TiB 2 are similar, the superposition of peaks may occur at 187.45 eV. The 192.55 eV corresponds to the B-O bond of B 2 O 3 [31], which is the oxidation on the surface of the composite powder. The pattern of Ti2p is divided into four peaks, which are two peaks generated by the energy level splitting of the Ti-B bond and Ti-O bond in Figure 4b. The Ti2p 3/2 sub-peak at 454.38 eV and the Ti2p 1/2 sub-peak at 458.95 eV correspond to the Ti-B bond of TiB 2 [32], which proves the existence of TiB 2 and the coating of the surface of B 4 C particle. The Ti2p 3/2 sub-peak at 459.67 eV and Ti2p 1/2 sub-peak at 464.86 eV correspond to the Ti-O bond in TiO 2 [33], due to the oxidation on the surface of the composite powder.

B 4 C@ZrB 2 -TiB 2 Composite Powder
level splitting of the Ti-B bond and Ti-O bond in Figure 4b. The Ti2p3/2 sub-peak at 454.38 eV and the Ti2p1/2 sub-peak at 458.95 eV correspond to the Ti-B bond of TiB2 [32], which proves the existence of TiB2 and the coating of the surface of B4C particle. The Ti2p3/2 subpeak at 459.67 eV and Ti2p1/2 sub-peak at 464.86 eV correspond to the Ti-O bond in TiO2 [33], due to the oxidation on the surface of the composite powder.  Figure 5 shows the XRD patterns of the composite ceramics sintering at 1700 °C. It can be seen from the picture that the primary phases are B4C (PDF #86-1024), (Zr, Ti)B2(PDF #89-3924), and C(PDF #26-1079). The diffraction peaks of ZrB2 (PDF #34-0423) shifts to high-angle azimuth (Figure 5b). As a result of the small size of Ti atoms, the solid solution replaces the Zr atoms into the lattice of ZrB2, causing distortion and compressive stress, resulting in the reduction of cell parameters (0.3100 nm) [27]. The diffraction peaks of ZrO2 disappear after high-temperature sintering. The lowest temperatures required for the complete conversion of Zr and Ti oxides into diborides were 1620 °C and 1690 °C under 0.8 CO ambient pressure. Furthermore, excessive B4C and lower CO partial pressure will further reduce the reaction temperature [34]. Asl et al. [35] found that ZrO2 will react with B4C to form a delicate ZrB2 connecting phase until it completely reacts in the subsequent heating process at 1281 °C.  Figure 6 shows the ternary phase diagram of the B4C-Zr-Ti system at 1700 °C obtained by FactSage calculations. When the B4C content is high, the region above the phase line L is a stable four-phase composition of B4C, ZrB2, TiB2, and C at 1700 °C. Combined  Figure 5 shows the XRD patterns of the composite ceramics sintering at 1700 • C. It can be seen from the picture that the primary phases are B 4 C (PDF #86-1024), (Zr, Ti)B 2 (PDF #89-3924), and C(PDF #26-1079). The diffraction peaks of ZrB 2 (PDF #34-0423) shifts to high-angle azimuth (Figure 5b). As a result of the small size of Ti atoms, the solid solution replaces the Zr atoms into the lattice of ZrB 2 , causing distortion and compressive stress, resulting in the reduction of cell parameters (0.3100 nm) [27]. The diffraction peaks of ZrO 2 disappear after high-temperature sintering. The lowest temperatures required for the complete conversion of Zr and Ti oxides into diborides were 1620 • C and 1690 • C under 0.8 CO ambient pressure. Furthermore, excessive B 4 C and lower CO partial pressure will further reduce the reaction temperature [34]. Asl et al. [35] found that ZrO 2 will react with B 4 C to form a delicate ZrB 2 connecting phase until it completely reacts in the subsequent heating process at 1281 • C.  Figure 5 shows the XRD patterns of the composite ceramics sintering at 1700 °C. It can be seen from the picture that the primary phases are B4C (PDF #86-1024), (Zr, Ti)B2(PDF #89-3924), and C(PDF #26-1079). The diffraction peaks of ZrB2 (PDF #34-0423) shifts to high-angle azimuth (Figure 5b). As a result of the small size of Ti atoms, the solid solution replaces the Zr atoms into the lattice of ZrB2, causing distortion and compressive stress, resulting in the reduction of cell parameters (0.3100 nm) [27]. The diffraction peaks of ZrO2 disappear after high-temperature sintering. The lowest temperatures required for the complete conversion of Zr and Ti oxides into diborides were 1620 °C and 1690 °C under 0.8 CO ambient pressure. Furthermore, excessive B4C and lower CO partial pressure will further reduce the reaction temperature [34]. Asl et al. [35] found that ZrO2 will react with B4C to form a delicate ZrB2 connecting phase until it completely reacts in the subsequent heating process at 1281 °C.  Figure 6 shows the ternary phase diagram of the B4C-Zr-Ti system at 1700 °C obtained by FactSage calculations. When the B4C content is high, the region above the phase line L is a stable four-phase composition of B4C, ZrB2, TiB2, and C at 1700 °C. Combined  Figure 6 shows the ternary phase diagram of the B 4 C-Zr-Ti system at 1700 • C obtained by FactSage calculations. When the B 4 C content is high, the region above the phase line L is a stable four-phase composition of B 4 C, ZrB 2 , TiB 2 , and C at 1700 • C. Combined with XRD analysis, ZrB 2 and TiB 2 will generate (Zr, Ti)B 2 solid solution after sintering at 1700 • C. with XRD analysis, ZrB2 and TiB2 will generate (Zr, Ti)B2 solid solution after sintering at 1700 °C.  Figure 7 shows the BSE images of the B4C-(Zr, Ti)B2 composite ceramics with different raw material ratios. Due to the larger atomic number of Zr and Ti atoms show a brighter white lining in the backscattering mode compared to B and C atoms. Combining the XRD analysis in Figure 5 and the EDS analysis (in Figure 7b,c) , it can be found that the bright white phases in Figure 7 are (Zr, Ti)B2. Figure 7a,b show the BSE image of the ZT8 sample and mapping analysis of the rectangular area. Zr and Ti elements are concentrated and evenly distributed in the bright white area. Furthermore, the same distribution interval of Zr and Ti elements and the shift of the XRD diffraction peak ( Figure 5) confirms the generation of the solid solution of ZrB2 and TiB2. Figure 7c,d are the BSE plots of ZT12 and ZT16 samples, respectively. With the decrease of the additional amount of Zr and Ti, the distribution of (Zr, Ti)B2 in the B4C matrix was gradually sparse and dispersed.  Figure 7 shows the BSE images of the B 4 C-(Zr, Ti)B 2 composite ceramics with different raw material ratios. Due to the larger atomic number of Zr and Ti atoms show a brighter white lining in the backscattering mode compared to B and C atoms. Combining the XRD analysis in Figure 5 and the EDS analysis (in Figure 7b,c), it can be found that the bright white phases in Figure 7 are (Zr, Ti)B 2 . Figure 7a,b show the BSE image of the ZT8 sample and mapping analysis of the rectangular area. Zr and Ti elements are concentrated and evenly distributed in the bright white area. Furthermore, the same distribution interval of Zr and Ti elements and the shift of the XRD diffraction peak ( Figure 5) confirms the generation of the solid solution of ZrB 2 and TiB 2 . Figure 7c,d are the BSE plots of ZT12 and ZT16 samples, respectively. With the decrease of the additional amount of Zr and Ti, the distribution of (Zr, Ti)B 2 in the B 4 C matrix was gradually sparse and dispersed.  Figure 8 shows the TEM image of ZT8 composite ceramic at 1700 °C. Figure 8b shows the elemental mapping analysis of Figure 8a. The surface of the ceramic sample shows the aggregation of B, Zr and Ti. Further illustration of (Zr, Ti)B2 solid solution formation and good bonding with B4C grains. The segregation of the C elements in the grain is mainly due to the influence of carbon supporting film.  Figure 8 shows the TEM image of ZT8 composite ceramic at 1700 • C. Figure 8b shows the elemental mapping analysis of Figure 8a. The surface of the ceramic sample shows the aggregation of B, Zr and Ti. Further illustration of (Zr, Ti)B 2 solid solution formation and good bonding with B 4 C grains. The segregation of the C elements in the grain is mainly due to the influence of carbon supporting film.  Figure 8 shows the TEM image of ZT8 composite ceramic at 1700 °C. Figure 8b shows the elemental mapping analysis of Figure 8a. The surface of the ceramic sample shows the aggregation of B, Zr and Ti. Further illustration of (Zr, Ti)B2 solid solution formation and good bonding with B4C grains. The segregation of the C elements in the grain is mainly due to the influence of carbon supporting film.     (Figure 7), when the amount of Zr/Ti is reduced, the generation of (Zr, Ti)B 2 is diminished, weakening the pinning effect on the B 4 C matrix-resulting in the coarsening of B 4 C grains, the reduction of the grain boundaries, and the reduction of the ability to resist local deformation. Ti 2+ solid solution into the ZrB 2 lattice will generate compression stress to enhance the internal energy. The active (Zr, Ti)B 2 grains formation will block the movement of the crystal world in the sintered process of the SPS, reducing the grain size, which plays the role of good crystal reinforcement, because internal stress will be generated around the grains when the solid solution is formed. Compared with the mechanical properties data of B 4 C composite ceramics reported in references (Listed in Table 1), such as higher temperature and longer holding time, this work realizes the preparation of high hardness B 4 C ceramics at low temperatures. The Vickers hardness is better than the data reported in the references due to the solution strengthening effect of ZrB 2 and TiB 2 . of Zr/Ti is reduced, the generation of (Zr, Ti)B2 is diminished, weakening the pinning effect on the B4C matrix-resulting in the coarsening of B4C grains, the reduction of the grain boundaries, and the reduction of the ability to resist local deformation. Ti 2+ solid solution into the ZrB2 lattice will generate compression stress to enhance the internal energy. The active (Zr, Ti)B2 grains formation will block the movement of the crystal world in the sintered process of the SPS, reducing the grain size, which plays the role of good crystal reinforcement, because internal stress will be generated around the grains when the solid solution is formed. Compared with the mechanical properties data of B4C composite ceramics reported in references (Listed in Table 1), such as higher temperature and longer holding time, this work realizes the preparation of high hardness B4C ceramics at low temperatures. The Vickers hardness is better than the data reported in the references due to the solution strengthening effect of ZrB2 and TiB2.    Through the XRD phase analysis and microstructure of the composite powder and composite ceramics, and the ternary phase diagram of the B 4 C-Zr-Ti system at 1700 • C, the reaction mechanism of composite powder and ceramic was been studied: The mass transfer of Zr and Ti could be promoted because of the existence of a molten salt medium, Zr and Ti would react with B 4 C to form ZrB 2 and TiB 2 at high temperature and the product would coat the B 4 C surface. Thus, the composite powder exhibited a core-shell structure. When the composite powder was sintered by high temperature and pulse activation into composite ceramics, ZrB 2 and TiB 2 of the powder surface would produce a solid solution (Zr, Ti)B 2 ). These (Zr, Ti)B 2 solid solutions would prevent the movement of crystal boundaries during SPS sintering, thereby reducing the grain size of the composite ceramics. The reaction mechanism is shown in Figure 10.

Conclusions
B4C-(Zr, Ti)B2 composite ceramics were obtained by SPS sintering of powder synthesized by the molten salt method, and its microstructure a properties were studied. The following conclusions were drawn.
(1) Thermodynamic calculation and phase analysis revealed that Zr and B4C to form ZrB2 and TiB2 in a B4C-rich environment. Thanks to the transfer of molten salt medium, the product is uniformly coated on the particles. (2) When the powder is sintered by SPS at 1700 ℃, ZrB2 and TiB2 form sol Ti)B2 at high temperature, the primary phases of composite ceramic

Conclusions
B 4 C-(Zr, Ti)B 2 composite ceramics were obtained by SPS sintering of the composite powder synthesized by the molten salt method, and its microstructure and mechanical properties were studied. The following conclusions were drawn.
(1) Thermodynamic calculation and phase analysis revealed that Zr and Ti reacted with B 4 C to form ZrB 2 and TiB 2 in a B 4 C-rich environment. Thanks to the heat and mass transfer of molten salt medium, the product is uniformly coated on the surface of B 4 C particles. (2) When the powder is sintered by SPS at 1700°C, ZrB 2 and TiB 2 form solid solution (Zr, Ti)B 2 at high temperature, the primary phases of composite ceramics are B 4 C, (Zr, Ti)B 2 , and C. Institutional Review Board Statement: Not applicable.
Informed Consent Statement: Not applicable.
Data Availability Statement: Data available on request due to restrictions privacy. The data presented in this study are available on request from the corresponding author.

Conflicts of Interest:
The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.