In Situ ZrB2 Formation in B4C Ceramics and Its Strengthening Mechanism on Mechanical Properties

In order to reduce the sintering temperature and improve the mechanical properties of B4C ceramics, ZrB2 was formed in situ using the SPS sintering method with ZrO2 and B4C as raw materials. Thermodynamic calculations revealed that CO pressure affected the formation of ZrB2 at temperatures from 814 °C to 1100 °C. The experimental results showed that the ZrB2 grain size was <5 µm and that the grains were uniformly distributed within the B4C ceramics. With an increase in ZrO2 content, the Vickers hardness and flexural strength of the B4C ceramics first increased and then decreased, while the fracture toughness continuously increased. When the content of ZrO2 was 15 wt%, the Vickers hardness, fracture toughness and flexural strength of B4C ceramics were 35.5 ± 0.63 GPa, 3.6 ± 0.24 MPa·m1/2 and 403 ± 10 MPa, respectively. These results suggest that ZrB2 inhibits B4C grain growth, eliminates crack tip stress, and provides fine grain to strengthen and toughen B4C ceramics.


Introduction
Due to their low density, high hardness and high elastic modulus, boron carbide (B 4 C) ceramics have been used in the manufacturing of armor, cutting tools and neutron absorbers [1][2][3][4][5]. However, further application of B 4 C ceramics is limited by the high sintering temperature (>2000 • C) and low fracture toughness (K IC < 2.2 MPa·m 1/2 ), which can be attributed to the strong B-C covalent bonds, low atomic diffusion coefficient and low dislocation plasticity [6][7][8][9][10]. Therefore, it is essential to lower the sintering temperature and improve the mechanical properties of B 4 C ceramics.
The addition of second phases, such as SiC, Al 2 O 3 , TiC and TiO 2 , is an effective method of reducing the sintering temperature and improving the mechanical properties of B 4 C ceramics [11,12], but ultimately results in reduced material hardness. Moreover, brittle compounds are formed during the densification process, leading to a low fracture toughness. Transitional metal diborides (TiB 2 , ZrB 2 ) are useful additives which have similar properties to B 4 C ceramics [13][14][15]. Among them, Zirconium Diboride (ZrB 2 ) has a high melting point (3245 • C), high hardness (14-23 GPa) and high flexural strength (higher than 500 MPa), which could improve the flexural strength and fracture toughness of B 4 C ceramics [16,17].
Many reports have shown that the addition of ZrB 2 could efficiently enhance the mechanical properties of B 4 C ceramics. For example, He et al. [18] found that ZrB 2 promoted the sintering process and increased the ceramic's high critical crack size, resulting in excellent thermal shock resistance. In order to inhibit the grain growth of B 4 C, Lin et al. [19] prepared B 4 C ceramics by pressureless sintering using ZrO 2 as a sintering aid. The in situ formation of ZrB 2 inhibited the abnormal growth of B 4 C grains, prolonged the 2 of 11 crack extension path and improved the sintering activity. Guo et al. [20] prepared B 4 C-ZrB 2 ceramics at low temperature, resulting in an average B 4 C grain size of 1-2 µm. The hardness and fracture toughness reached 33.0 ± 1.3 GPa and 4.1 ± 0.2 MPa·m 1/2 , respectively. Rehman et al. [21] prepared B 4 C ceramics by SPS at 1700 • C; the average grain size of ZrB 2 was 1.3 ± 0.2 µm. The Vickers hardness, fracture toughness and flexural strength were 31.28 GPa, 4.2 MPa·m 1/2 and 511 MPa, respectively. However, the role of ZrB 2 as a sintering aid and grain growth inhibitor requires further investigation, and the effect of ZrB 2 on the fracture mode transformation of B 4 C ceramics also needs to be clarified.
In this work, we aim to obtain densified B 4 C ceramics with the in situ formation of ZrB 2 at 1700 • C using the spark plasma sintering (SPS) method, with ZrB 2 grains homogeneously distributed in the B 4 C ceramics. The reaction between ZrO 2 and B 4 C could effectively reduce the sintering temperature, and the pinning effect of ZrB 2 is expected to inhibit the B 4 C grain growth, achieving desirable mechanical properties.
The compositions of the samples are shown in Table 1. The raw materials were mixed with stainless steel balls in anhydrous ethanol with a ball-to-material ratio of 5:1 and a mixing time of 6 h at 250 r/min. After mixing, the mixtures were dried in a drying oven at 80 • C for 24 h and then passed through a 200-mesh sieve. Then, the mixtures were sintered by SPS (Dr. Sinter-3.20 MKII, Sumitomo, Tokyo, JPN) under vacuum. The heating rate was maintained at 100 • C/min and a uniaxial load of 5 kN was applied during the heating stage to ensure good contact conductivity with the graphite indenter. After the temperature rose to 1600 • C, the uniaxial load was slowly increased to 15 kN. Then, the sintering temperature was increased to 1700 • C for 6 min. The sintered ceramics were cooled in the furnace to about 150 • C and the samples were taken out to obtain B 4 C ceramics.

Characterizations
The bulk density of B 4 C ceramics was measured using the Archimedes method. The Vickers hardness and fracture toughness were measured by the indentation method using a Vickers hardness tester (HV-50A, Huayin, Laizhou, China). The spacing between each indentation was ensured to be greater than 100 µm to avoid crack deformation or the overlapping of extended areas caused by indentation. Five effective indentations were also used to calculate the mean and error of Vickers hardness and fracture toughness. The B 4 C ceramics were processed into long strip samples of 2 mm × 3 mm × 20 mm using a diamond wire cutting machine, and the bending strength of the samples was measured using a three-point bending test with a span of 16 mm and a loading rate of 0.5 mm/min. The mean values of flexural strength and error analyses were calculated with 3 sets of valid data. The phase compositions were characterized by X-ray diffractometer (XRD, X'Pert MPD Pro, Philips, Amsterdam, The Netherlands). The voltage and current of the Cu Kα1 radiation were 40 kV and 40 mA, respectively. The scanning range (2θ) was 10 • -90 • . The microscopic morphologies were characterized using a field-emission scanning electron microscope (FESEM, Nova NanoSem 400, Thermo, Waltham, MA, USA).

Thermodynamic Calculations
The in situ generation of ZrB 2 from B 4 C and nano-ZrO 2 is represented in Reaction (1): According to thermodynamic calculations by FactSage, the variation in Gibbs' free energy of reaction (1) as a function of temperature at different P CO (pressure of CO) and as a function of P CO at different temperatures is shown in Figure 1. Figure 1a shows that in the standard state (P CO = 0.986 atm), a temperature above 1212 • C is favorable for the formation of ZrB 2 . It is expected that reaction (1) proceeds more rapidly when the reaction is carried out under vacuum. Therefore, pure ZrB 2 can be synthesized at a temperature as low as 814 • C when P CO = 10 −5 atm, according to Figure 1a. However, the CO generated by the sample surrounded by graphite die cannot be expelled in time, leading to P CO > 10 −4 atm; thus, the reaction will proceed at a temperature higher than 814 • C. Figure 1b further illustrates that reaction (1) needs to proceed at temperatures above 1100 • C when P CO is in the range of 0-0.1 atm. In the experiments, the vacuum pump worked continuously and the vacuum level in the furnace was as low as 8 × 10 −5 atm during SPS, maintaining the P CO at a low level and thus promoting the low-temperature densification of B 4 C ceramics.
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Thermodynamic Calculations
The in situ generation of ZrB2 from B4C and nano-ZrO2 is represented in Reaction (1): According to thermodynamic calculations by FactSage, the variation in Gibbs' free energy of reaction (1) as a function of temperature at different PCO (pressure of CO) and as a function of PCO at different temperatures is shown in Figure 1. Figure 1a shows that in the standard state (PCO = 0.986 atm), a temperature above 1212 °C is favorable for the formation of ZrB2. It is expected that reaction (1) proceeds more rapidly when the reaction is carried out under vacuum. Therefore, pure ZrB2 can be synthesized at a temperature as low as 814 °C when PCO = 10 −5 atm, according to Figure 1a. However, the CO generated by the sample surrounded by graphite die cannot be expelled in time, leading to PCO > 10 −4 atm; thus, the reaction will proceed at a temperature higher than 814 °C. Figure 1b further illustrates that reaction (1) needs to proceed at temperatures above 1100 °C when PCO is in the range of 0-0.1 atm. In the experiments, the vacuum pump worked continuously and the vacuum level in the furnace was as low as 8 × 10 −5 atm during SPS, maintaining the PCO at a low level and thus promoting the low-temperature densification of B4C ceramics.

XRD Phase Analysis
The XRD patterns of B4C ceramics prepared by SPS with different ZrB2 contents are shown in Figure 2. The ceramics are composed of B4C and ZrB2, and no residual nano-ZrO2 phase can be detected, indicating that nano-ZrO2 and B4C completely reacted to form the ZrB2 phase. The B4C in the samples all had the rhombohedral crystal structure (PDF #35-0798). The diffraction peaks of samples at 2θ of 21.21°, 32.64°, and 41.69° were all taken from the ZrB2 crystal (PDF #89-3930), and correspond to the (110), (200) and (211) crystal planes of its hexagonal crystal structure, respectively. The intensity of the ZrB2 diffraction peak increases with the addition of nano-ZrO2, while the intensity of the B4C diffraction peak is the opposite, which is consistent with the expected results. The XRD results indicate that nano-ZrO2 can react with B4C in situ to form ZrB2, which can promote the diffusion effect between B4C particles during sintering and achieve the densification of B4C ceramics.

XRD Phase Analysis
The XRD patterns of B 4 C ceramics prepared by SPS with different ZrB 2 contents are shown in Figure 2. The ceramics are composed of B 4 C and ZrB 2 , and no residual nano-ZrO 2 phase can be detected, indicating that nano-ZrO 2 and B 4 C completely reacted to form the ZrB 2 phase. The B 4 C in the samples all had the rhombohedral crystal structure (PDF #35-0798). The diffraction peaks of samples at 2θ of 21.21 • , 32.64 • , and 41.69 • were all taken from the ZrB 2 crystal (PDF #89-3930), and correspond to the (110), (200) and (211) crystal planes of its hexagonal crystal structure, respectively. The intensity of the ZrB 2 diffraction peak increases with the addition of nano-ZrO 2 , while the intensity of the B 4 C diffraction peak is the opposite, which is consistent with the expected results. The XRD results indicate that nano-ZrO 2 can react with B 4 C in situ to form ZrB 2 , which can promote the diffusion effect between B 4 C particles during sintering and achieve the densification of B 4 C ceramics.

Microstructural Analysis
The BSE images of B4C ceramics and the grain size of ZrB2 prepared by SPS are shown in Figure 3. The distribution of B, C and Ti elements in sample BZ15 are shown in Figure  3f-i. Combined with XRD, it can be deduced that the dark gray phase is B4C, and the bright white phase is ZrB2. Figure 3a shows that the addition of 3 wt% nano-Y2O3 can achieve a polished surface of B4C ceramics with no obvious pores without the addition of nano-ZrO2. Figure 3b-e show that the amount of ZrB2 produced gradually increases as the addition of nano-ZrO2 increases from 5 wt% to 20 wt%, and the average particle size of ZrB2 increases from 1.03 µm to 1.37 µm. The in situ reaction of nano-ZrO2 to generate ZrB2 facilitates the flowability between B4C particles and promotes the densification and pore discharge of B4C ceramics, such that there are no obvious pores on the polished surface. In addition, ZrB2 not only has small grain size, but is also uniformly distributed within the B4C ceramics, which will effectively improve the mechanical properties.

Microstructural Analysis
The BSE images of B 4 C ceramics and the grain size of ZrB 2 prepared by SPS are shown in Figure 3. The distribution of B, C and Ti elements in sample BZ15 are shown in Figure 3f-i. Combined with XRD, it can be deduced that the dark gray phase is B 4 C, and the bright white phase is ZrB 2 . Figure 3a shows that the addition of 3 wt% nano-Y 2 O 3 can achieve a polished surface of B 4 C ceramics with no obvious pores without the addition of nano-ZrO 2 . Figure 3b-e show that the amount of ZrB 2 produced gradually increases as the addition of nano-ZrO 2 increases from 5 wt% to 20 wt%, and the average particle size of ZrB 2 increases from 1.03 µm to 1.37 µm. The in situ reaction of nano-ZrO 2 to generate ZrB 2 facilitates the flowability between B 4 C particles and promotes the densification and pore discharge of B 4 C ceramics, such that there are no obvious pores on the polished surface. In addition, ZrB 2 not only has small grain size, but is also uniformly distributed within the B 4 C ceramics, which will effectively improve the mechanical properties. BSE diagrams of the fracture morphology of B4C ceramics with different nano-ZrO2 contents are shown in Figure 4. The fracture mode is mainly through crystal fracture, and the fracture surface is relatively neat and smooth. This indicates a favorable degree of connection between the grains after SPS. In addition, there are more pores in the BZ0 sample, but the number of pores decreases significantly with the addition of nano-ZrO2. This indicates that nano-ZrO2 promotes the densification of B4C ceramics.  BSE diagrams of the fracture morphology of B 4 C ceramics with different nano-ZrO 2 contents are shown in Figure 4. The fracture mode is mainly through crystal fracture, and the fracture surface is relatively neat and smooth. This indicates a favorable degree of connection between the grains after SPS. In addition, there are more pores in the BZ0 sample, but the number of pores decreases significantly with the addition of nano-ZrO 2 . This indicates that nano-ZrO 2 promotes the densification of B 4 C ceramics. BSE diagrams of the fracture morphology of B4C ceramics with different nano-ZrO2 contents are shown in Figure 4. The fracture mode is mainly through crystal fracture, and the fracture surface is relatively neat and smooth. This indicates a favorable degree of connection between the grains after SPS. In addition, there are more pores in the BZ0 sample, but the number of pores decreases significantly with the addition of nano-ZrO2. This indicates that nano-ZrO2 promotes the densification of B4C ceramics. TEM images of BZ15 are shown in Figure 5. Figure 5a shows the interface between B4C and ZrB2 in the B4C ceramics. The HRTEM image shown in Figure 5b indicates that the grain boundary between B4C (d = 0.3781 nm, (012) zone axis) and ZrB2 (d = 0.2166 nm, (101) zone axis) is clean. This result indicates that the grain boundary strength between B4C and ZrB2 is strong. It is well-known that a high binding force between the matrix phase and second phase has a significant effect on mechanical properties. Therefore, this suggests that excellent mechanical properties can be obtained from these B4C ceramic samples. TEM images of BZ15 are shown in Figure 5. Figure 5a shows the interface between B 4 C and ZrB 2 in the B 4 C ceramics. The HRTEM image shown in Figure 5b indicates that the grain boundary between B 4 C (d = 0.3781 nm, (012) zone axis) and ZrB 2 (d = 0.2166 nm, (101) zone axis) is clean. This result indicates that the grain boundary strength between B 4 C and ZrB 2 is strong. It is well-known that a high binding force between the matrix phase and second phase has a significant effect on mechanical properties. Therefore, this suggests that excellent mechanical properties can be obtained from these B 4 C ceramic samples.

Mechanical Properties of the B4C Ceramics
The relative density and bulk density of B4C ceramics prepared by SPS with different ZrO2 contents are shown in Figure 6. The relative densities measured are all above 98.3%, and increase and then decrease with increasing ZrO2 content. The bulk density continues to increase, between 2.52 g/cm 3 and 2.83 g/cm 3 . The relative density reaches a maximum of 99.6% when the content of ZrO2 is 5 wt%; this can be attributed to the in situ reaction of nano-ZrO2 to generate ZrB2, which that promotes the flowability between B4C particles and obtains dense ceramics. When the content of ZrO2 is higher than 5 wt%, the CO gas generated by the sintering reaction also increases, causing the relative density to decrease. The observed increase in bulk density can be explained by the fact that the theoretical density of ZrB2 (6.1 g/cm 3 ) is greater than that of B4C, meaning that an increase in ZrB2 content leads to a gradual increase in the bulk density of the ceramics.

Mechanical Properties of the B 4 C Ceramics
The relative density and bulk density of B 4 C ceramics prepared by SPS with different ZrO 2 contents are shown in Figure 6. The relative densities measured are all above 98.3%, and increase and then decrease with increasing ZrO 2 content. The bulk density continues to increase, between 2.52 g/cm 3 and 2.83 g/cm 3 . The relative density reaches a maximum of 99.6% when the content of ZrO 2 is 5 wt%; this can be attributed to the in situ reaction of nano-ZrO 2 to generate ZrB 2 , which that promotes the flowability between B 4 C particles and obtains dense ceramics. When the content of ZrO 2 is higher than 5 wt%, the CO gas generated by the sintering reaction also increases, causing the relative density to decrease. The observed increase in bulk density can be explained by the fact that the theoretical density of ZrB 2 (6.1 g/cm 3 ) is greater than that of B 4 C, meaning that an increase in ZrB 2 content leads to a gradual increase in the bulk density of the ceramics. The observed increase in bulk density can be explained by t density of ZrB2 (6.1 g/cm 3 ) is greater than that of B4C, meanin content leads to a gradual increase in the bulk density of the ce The Vickers hardness, fracture toughness and flexural str pared by SPS at different ZrO2 contents are shown in Figure 7 seen that with the increase of ZrO2 content, the Vickers hard increasing and then decreasing, and the fracture toughness gra ers hardness is maintained above 35.5 GPa with a of ZrO2 co creases to 32.3 GPa with a ZrO2 content of 20 wt%. Fracture tou from 2.3 MPa·m 1/2 (BZ0) to 3.96 MPa·m 1/2 (BZ20). This is becaus density of B4C ceramics leads to an increase in Vickers hardnes can also have a negative effect on Vickers hardness. The ZrB2 g ZrO2 has a pinning effect, which inhibits the growth of B4C p The Vickers hardness, fracture toughness and flexural strength of B 4 C ceramics prepared by SPS at different ZrO 2 contents are shown in Figure 7. From Figure 7a, it can be seen that with the increase of ZrO 2 content, the Vickers hardness shows a trend of first increasing and then decreasing, and the fracture toughness gradually increases. The Vickers hardness is maintained above 35.5 GPa with a of ZrO 2 content of 5-15 wt%, but decreases to 32.3 GPa with a ZrO 2 content of 20 wt%. Fracture toughness gradually increases from 2.3 MPa·m 1/2 (BZ0) to 3.96 MPa·m 1/2 (BZ20). This is because an increase in the relative density of B 4 C ceramics leads to an increase in Vickers hardness, but the presence of pores can also have a negative effect on Vickers hardness. The ZrB 2 generated in situ from nano-ZrO 2 has a pinning effect, which inhibits the growth of B 4 C particles. At the same time, the residual stresses present during cooling can block the crack expansion of B 4 C ceramics due to the difference in thermal expansion coefficients of B 4 C and ZrB 2 . Both processes are beneficial for the improvement of the fracture toughness of B 4 C ceramics. As can be seen in Figure 7b, the bending strength of the B 4 C ceramics shows a trend of increasing and then decreasing with an increase in ZrO 2 content. The bending strength reaches a maximum value of 404 MPa at 15 wt% ZrO 2 , which is 60.3% higher than the BZ0 sample (252 MPa). The addition of nano-ZrO 2 promotes the sintering between B 4 C particles and inhibits their further growth. Furthermore, the bending strength of ZrB 2 is greater than that of B 4 C, ultimately improving the bending strength of the B 4 C ceramics. However, when the content of ZrO 2 reaches 20 wt%, the generation of CO gas leads to an increase in the number of pores and a decrease in the density, which results in a lower bending strength. When the content of ZrO 2 is 15 wt%, the Vickers hardness, fracture toughness and bending strength of sample BZ15 are 35.5 GPa, 3.6 MPa·m 1/2 and 404 MPa, respectively, with optimal overall performance. B4C, ultimately improving the bending strength of the B4C ceramics. However, when the content of ZrO2 reaches 20 wt%, the generation of CO gas leads to an increase in the number of pores and a decrease in the density, which results in a lower bending strength. When the content of ZrO2 is 15 wt%, the Vickers hardness, fracture toughness and bending strength of sample BZ15 are 35.5 GPa, 3.6 MPa·m 1/2 and 404 MPa, respectively, with optimal overall performance.

Densification and Toughening Mechanism Analysis
Crack propagation within the BZ15 sample following the Vickers hardness indentation tests is shown in Figure 8. There are obvious branching and deflection phenomena in crack extension, which can increase the energy required for crack extension and thus improve the fracture toughness of B4C ceramics. In addition, the fracture mode of the B4C ceramics changes from a traditional transgranular fracture to a hybrid mode of intergranular and transgranular fracture, which is beneficial for the improvement of fracture toughness.

Densification and Toughening Mechanism Analysis
Crack propagation within the BZ15 sample following the Vickers hardness indentation tests is shown in Figure 8. There are obvious branching and deflection phenomena in crack extension, which can increase the energy required for crack extension and thus improve the fracture toughness of B 4 C ceramics. In addition, the fracture mode of the B 4 C ceramics changes from a traditional transgranular fracture to a hybrid mode of intergranular and transgranular fracture, which is beneficial for the improvement of fracture toughness. content of ZrO2 reaches 20 wt%, the generation of CO gas leads to an increase in the number of pores and a decrease in the density, which results in a lower bending strength. When the content of ZrO2 is 15 wt%, the Vickers hardness, fracture toughness and bending strength of sample BZ15 are 35.5 GPa, 3.6 MPa·m 1/2 and 404 MPa, respectively, with optimal overall performance.

Densification and Toughening Mechanism Analysis
Crack propagation within the BZ15 sample following the Vickers hardness indentation tests is shown in Figure 8. There are obvious branching and deflection phenomena in crack extension, which can increase the energy required for crack extension and thus improve the fracture toughness of B4C ceramics. In addition, the fracture mode of the B4C ceramics changes from a traditional transgranular fracture to a hybrid mode of intergranular and transgranular fracture, which is beneficial for the improvement of fracture toughness.  The different coefficients of thermal expansion between the matrix and the second phase in ceramics could induce residual stress during the cooling process. The effect of this is more obvious when the strength and elastic modulus of the second phase are higher than those of the matrix. Since the thermal expansion coefficients of B 4 C and ZrB 2 are quite different, the interface between ZrB 2 and B 4 C will therefore be subject to residual stress (p i ) during the cooling process [22]. where ∆α is the difference between the coefficients of thermal expansion of B 4 C and ZrB 2 , ∆T is the temperature difference between the cooling stages of SPS, ν is Poisson's ratio, E is elastic modulus, and the subscripts p and m stand for ZrB 2 and B 4 C, respectively. In this condition, σ r (radial stress) and σ t (tangential stress) are generated in the matrix at a distance r from the particles [22]. and where R is the radius of the particle and r is the distance from a point in the stress field to the center of the particle. It can be seen that both σ r and σ t are proportional to (R/r) 3 . When a crack passes through this stress field, a deflection will occur, as shown in Figure 9. When second phase particles are encountered ahead of the crack, the radial tensile stress σ r increases in the matrix (σ r is maximal at the interface), causing the crack to deflect in the direction of the particles and reach the particle/matrix interface before propagating along the original extension direction. The combined effect of σ r and σ t lengthens the crack expansion path in the matrix and consumes the energy for crack expansion. Therefore, the crack deflection caused by the residual stress field produces an absolute toughening effect. Figure 8 shows that the crack extension undergoes deflection and bypass, and that there is a secondary crack generation on the primary crack, which increases the path of crack extension and consumes the energy of crack extension. If the crack propagation path is more tortuous, and the volume fraction of the second phase component which leads to crack deflection is higher, then the effect of crack deflection and toughening will be significant.

Comparison of Properties of B4C Ceramics in Recent Years
The Vickers hardness and fracture toughness of B4C ceramics prepared under different sintering methods in recent years are shown in Figure 10 [20,21,[23][24][25][26][27][28][29][30][31][32][33]. The different sintering methods described are pressureless sintering, hot-pressure sintering and spark plasma sintering. Although the Vickers hardness and fracture toughness of B4C ceramics are high, the sintering temperature of dense B4C ceramics is almost above 1800 °C. In this study, B4C ceramics are successfully prepared at 1700 °C by in situ reaction using nano-ZrO2 as a sintering aid. The maximum values of Vickers hardness and fracture toughness are 35.5 GPa and 3.6 MPa·m 1/2 , respectively. The in situ-generated ZrB2 inhibits the growth of B4C grains, deflecting and bifurcating the cracks, thus enhancing the mechanical properties of B4C ceramics.
The phase-field model of Roy shows that interfacial stress is an important factor affect-

Comparison of Properties of B 4 C Ceramics in Recent Years
The Vickers hardness and fracture toughness of B 4 C ceramics prepared under different sintering methods in recent years are shown in Figure 10 [20,21,[23][24][25][26][27][28][29][30][31][32][33]. The different sintering methods described are pressureless sintering, hot-pressure sintering and spark plasma sintering. Although the Vickers hardness and fracture toughness of B 4 C ceramics are high, the sintering temperature of dense B 4 C ceramics is almost above 1800 • C. In this study, B 4 C ceramics are successfully prepared at 1700 • C by in situ reaction using nano-ZrO 2

Conclusions
In this study, B4C ceramics were prepared by the in situ generation of ZrB2 using SPS at low temperature. The effects of ZrB2 on the microstructure and mechanical properties of the B4C ceramics were characterized and tested. The following conclusions were drawn.
(1) The ceramics were composed of B4C and ZrB2. The nano-ZrO2 reacted to form fine-crystalline ZrB2, which was uniformly distributed in the B4C ceramics. The in situ generation of ZrB2 contributed to the flowability between B4C particles, improving the sintering activity and achieving the densification of the B4C ceramics at low temperature.
(2) The relative density, Vickers hardness and bending strength of the B4C ceramics first increased and then decreased with the increase of ZrB2 content, and the bulk density and fracture toughness increased continuously. The relative density was above 98.3% and the bulk density was 2.52-2.83 g/cm 3 . When the content of ZrO2 was 15 wt%, the Vickers hardness, fracture toughness and bending strength were 35.5 ± 0.63 GPa, 3.6 ± 0.24 MPa·m 1/2 and 403 ± 10 MPa, respectively, with the best overall performance.
(3) The pinning effect of ZrB2 particles changed the fracture mode of the B4C ceramics from traditional transgranular fracture to a hybrid mode of intergranular and transgranular fracture, which improved the bending strength of the ceramics. The high bonding strength between powders and the residual stress between B4C and ZrB2 led to crack deflection and crack branching during the crack propagation, which resulted in the high mechanical properties of the B4C ceramics.   The phase-field model of Roy shows that interfacial stress is an important factor affecting the interfacial stress distribution at interfaces and the phase-field solution [34]. Therefore, the effect of interfacial stress between B 4 C and ZrB 2 , as well as the ZrO 2 phase transition process on the evolution of the nano-microstructure, must be studied in further detail. In addition, Mahdi investigated the effect of high pressure on the phase transformation [35], which may offer some insight into the effect of residual stress on ZrB 2 grain growth due to the difference in the coefficients of thermal expansion between B 4 C and ZrB 2 .

Conclusions
In this study, B 4 C ceramics were prepared by the in situ generation of ZrB 2 using SPS at low temperature. The effects of ZrB 2 on the microstructure and mechanical properties of the B 4 C ceramics were characterized and tested. The following conclusions were drawn.
(1) The ceramics were composed of B 4 C and ZrB 2 . The nano-ZrO 2 reacted to form fine-crystalline ZrB 2 , which was uniformly distributed in the B 4 C ceramics. The in situ generation of ZrB 2 contributed to the flowability between B 4 C particles, improving the sintering activity and achieving the densification of the B 4 C ceramics at low temperature.
(2) The relative density, Vickers hardness and bending strength of the B 4 C ceramics first increased and then decreased with the increase of ZrB 2 content, and the bulk density and fracture toughness increased continuously. The relative density was above 98.3% and the bulk density was 2.52-2.83 g/cm 3 . When the content of ZrO 2 was 15 wt%, the Vickers hardness, fracture toughness and bending strength were 35.5 ± 0.63 GPa, 3.6 ± 0.24 MPa·m 1/2 and 403 ± 10 MPa, respectively, with the best overall performance.
(3) The pinning effect of ZrB 2 particles changed the fracture mode of the B 4 C ceramics from traditional transgranular fracture to a hybrid mode of intergranular and transgranular fracture, which improved the bending strength of the ceramics. The high bonding strength between powders and the residual stress between B 4 C and ZrB 2 led to crack deflection and crack branching during the crack propagation, which resulted in the high mechanical properties of the B 4 C ceramics.