The boron carbide matrix composites containing boron carbide (B 4 C), titanium diboride (TiB 2 ; 20 wt% and 30 wt%) and titanium carbide (TiC) were fabricated at 2130 °C using the pressureless sintering method. Different amounts of TiB 2 and TiC particles were added to B 4 C, and the TiB 2 content was chosen as the main variable to study the effect on the composites. The density, hardness, bending strength and fracture toughness measurements were performed to obtain the physical and mechanical properties of the samples. The obtained results indicate that by adding 30 wt% TiB 2 submicron particles at 2130 °C, the bending strength and fracture toughness of 277.6 MPa and 5.38 MPa·m 1/2 are obtained for the composite. The particle pullout and crack microbridging play an effective role in toughening the composite ceramics. The main toughening mechanisms of the B 4 C-TiB 2 composites are noted to be microcrack and crack deflection toughening owing to the residual stresses resulting from the mismatch of the thermal expansion coefficient between TiB 2 and B 4 C.
Research Article
Effects of directly added or reaction formed TiB2 on the microstructure and mechanical properties of pressureless-sintered B4C–TiB2 composites
https://doi.org/10.21203/rs.3.rs-17972/v1
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Due to its low density as well as high hardness and melting point, boron carbide (B4C) is widely used for lightweight armor and cutting tools. B4C belongs to the group of non-metallic hard materials, with hardness inferior to only diamond and cubic boron nitride[1]. However, the limitations of poor toughness and sintering performance of B4C hinder its application to some extent[2–5].
Efforts have been made to enhance the sintering ability and mechanical properties of B4C. Compared to the sintering methods such as hot pressing and spark plasma sintering, pressureless sintering is a relatively simple and economical method that can be used for the preparation of large and complex parts[6]. Furthermore, many additives are added to B4C to improve the ceramic properties, and the results have shown that TiB2 is an ideal choice as an additive for the B4C matrix[7–10]. According to the literature studies, there are two main methods for adding TiB2 in the B4C ceramics. One involves the direct addition of TiB2 particles[11], whereas the other one is based on the chemical reaction of B4C with additives such as Ti[12, 13], TiO2[9] and TiC[14, 15], indicating the in-situ method to prepare TiB2. The addition of TiC can effectively improve the properties of B4C, and the studies show that TiC reacts with B4C to form TiB2 in the system (Reaction 1). However, in B4C–TiB2 composites field, most researches focus on reactive hot-pressing, while little researches focus on directly added or reaction formed TiB2 by easy and cheap pressureless sintered. Large-scale industrial production of ceramics is hampered by the high economic and poor efficiency costs of hot-pressing.
Extensive studies on improving the toughness of the B4C ceramics, evolved by combining the addition of TiB2 particles and pressureless sintering, are necessary. Few studies have explored the different effects of directly added or B4C-TiC reaction formed TiB2 on B4C using the pressureless sintering method based on the corresponding impact on the microstructure and mechanical properties. And what remains unclear about toughening mechanism is whether two methods of forming TiB2 promote each other or only one method played a major role during the synthesis. In this study, the influence of TiB2 addition on the microstructure and mechanical properties of pressureless sintered B4C–TiB2 composites at 2130 ℃ has been investigated. The different addition methods were evaluated by comparing the corresponding impact on the microstructure and mechanical properties of the B4C–TiB2 composites, such as density, hardness, bending strength and fracture toughness.
Commercially available B4C powder (~ 3.19 µm, Zhengzhou Songshan Material Co., China), TiC powder (~ 5.02 µm, Shanghai Shuitian Material Co., China) and TiB2 powder (~ 0.20 µm, Shanghai Shuitian Material Co., China) were used as raw materials. Si powder (~ 2.40 µm, Shanghai Shuitian Material Co., China) and C powder (~ 0.50 µm, Shanghai Shuitian Material Co., China) were used as sintering aids, as results show that C and Si are effective sintering aids for improving the sintering ability and mechanical properties [16]. The powder mixtures of six different compositions (Table 1) were prepared by wet ball-milling at 200 rpm for 12 h in a polyethylene bottle, using ZrO2 balls and alcohol medium. The weight ratio of the balls to powders was 3:1. The mixtures were dried in a rotary evaporator, followed by crushing and screening through a 100 mesh sieve. Afterwards, green compacts were formed by uniaxial pressing at about 50 MPa, followed by cold isostatic pressing at 200 MPa and sintering using a graphite element vacuum furnace at 2130 °C for 40 min. The initial heating rate was 6–8° C/min. After reaching 1800° C, the heating rate was enhanced to 9–10 °C/min. The temperature was maintained for 40 min. It was followed by cooling at a rate of 9–11 ℃/ min. After cooling down to 500 ℃, the materials were slowly cooled along with the furnace. The sintering process curve has been shown in Fig. 1.
The samples were polished and density values were measured by the standard Archimedes’ method. The bending strength and the fracture toughness (KIC) were measured on a universal testing machine (INSTRON-5985, Norwood, America). The three-point bending test with a span of 30 mm and a loading rate of 0.5 mm/min was conducted on the specimens (36.0 mm × 3.0 mm × 4.0 mm) to evaluate the bending strength; the single-edge notched beam method with a span of 30 mm and a loading rate of 0.05 mm/min was used on the specimens (36.0 mm × 3.0 mm × 6.0 mm) to measure the fracture toughness. The bending strength and the fracture toughness values were determined from the results averaged from seven bars, respectively. Crystalline phases in the sintered specimens were performed with X-Ray diffraction (XRD, Ultima III, Rigaku, Japan). Analysis of the microstructure and compositional analysis were examined by Scanning Electron Microscopy (SEM, Hitachi-S3400N, Japan).
3.1 Phase transition process
Figure 2 shows the XRD patterns of the as-mixed B4C/TiB2/TiC/Si/C composite powders. As expected, B4C and TiB2 are the main crystalline phases in all powder samples. Si reacted with the C powder or C element in the B4C matrix, thus, forming a SiC phase, as detected in the XRD pattern. Furthermore, it could be observed that the TiC phase disappeared, which indicated that the Reaction 1 occurred during the sintering of the ceramic samples and the TiC phase reacted completely. This suggests that the main reason affecting the microstructure and properties of the multiphase ceramics is the presence of TiB2 as the second phase, for the pre-designed three-phase or two-phase ceramic systems.
3.2 Microstructure
The backscattered electron (BSE) images of the polished B4C-TiB2 composites are shown in Fig. 3. It can be seen that the B4C-TiB2 composites consist of three phases, the dark areas represent the B4C phase, the bright areas are the TiB2 phase and the gray areas are the SiC phase. The SiC phase with small size is noted to be dispersed in the B4C matrix and presents the plate-like tissue morphology with a large aspect ratio, thus, helping to improve the mechanical properties of the whisker toughened ceramic. A small amount of laminated (white dotted bordered rectangle) and flocculated (black dotted bordered rectangle) graphite phase can also be seen in Fig. 4. Generally, there are two sources of the graphite phase. In the first case, the fine and dispersed TiB2 and graphite phases were formed by the reaction of B4C with TiC during pressureless sintering, observed to be closely bound to the matrix. In the second case, the C particles used as sintering aid are graphitized at high temperature, thus, forming the graphite phase. The graphite grains promote the sintering rate as graphite located on the grain boundary inhibits the grain growth and increases the grain boundary quantity, thereby, increasing the material transport between the adjacent B4C grains, thus, promoting sintering. The graphite phase in the matrix can absorb the energy of the crack propagation by torsion, deflection and fracture under the action of the external forces. Furthermore, the crack propagation direction can be made perpendicular to the laminated graphite direction to further hinder the crack propagation.
Figure 5 and Fig. 6 show the secondary electron (SE) and BSE fracture images of the composites of Group 1 and Group 2, respectively. The porosity was observed in all samples, with Group 2 exhibiting lower porosity, thus, indicating a higher degree of densification for Group 2. It also suggested that 30 wt% TiB2 could effectively improve the density of B4C. The agglomerated graphite filled in the pores of the ceramic matrix is commonly observed to enhance the densification. A few sharp sintering concave necks are noted in Fig. 5a, which also confirms the reason for the low relative density of 80B4C-10TiC, as shown in Fig. 9a. The more rounded is the sintering neck, the fuller is the sintering, thus, the higher is the sintering density. A few regional neck growths are noted to increase and develop into a three-dimensional interpenetrating network of particles, thus, leading to further increase in the degree of densification in 70B4C-20TiB2 and 70B4C-6.67TiC-13.33TiB2. The pores are observed to open and get interconnected to facilitate the gas discharge in the sintering process. Thus, the microstructure of 70B4C-20TiB2 developed more obviously, along with further densification to 98.86%. Similar phenomena is observed to occur for Group 2 materials, as shown in Fig. 6. The sintering necks of the three samples grew, with 70B4C-20TiC and 60B4C-30TiB2 exhibiting a higher degree of growth and fewer pores. Hence, the densification of 70B4C-20TiC and 60B4C-30TiB2 is relatively higher in Group 2, with a relative density of 97.63% and 94.91%, respectively, as shown in Fig. 9a. It can be seen in Fig. 6d that the TiB2 particles are finer and more evenly distributed than those in Fig. 6b and Fig. 6 f. The fine microstructure leads to high mechanical properties due to the Hall-Petch effect, whereas the evenly distributed particles can effectively toughen the composite ceramics, thus, resulting in optimal mechanical properties for the 60B4C-30TiB2 composite. Comparing the SE images with the corresponding BSE images, the fracture mode of B4C is noted to be mainly transgranular, whereas it is mainly intergranular for TiB2. The mixed fracture mode is favorable for the toughness of the B4C- TiB2 composites, as also reported in literature [8, 9].
Figure 7 shows the crack growth path of the B4C–TiB2 composite. It is obvious that numerous crack deflections occurred during the crack propagation process, and the large angle crack deflection (blue dotted bordered rectangle) occurred in the two-phase junction region. Also, the crack microbridging (red dotted bordered rectangle) was observed, and the results indicated that it could improve the fracture toughness of the ceramics. Most of the cracks were observed to grow around the microcrack and pores (white dotted bordered rectangle). Besides, the grains deflected the cracks (black dotted bordered rectangle), thus, the fracture mode changed from the transgranular fracture to intergranular fracture, followed by changing back to the transgranular fracture after the crack moved around the grain. The phenomenon of crack branching can also be noted in the upper left corner of Fig. 7. As the main crack leading to the fracture expands, the evenly distributed microcracks due to the different thermal expansion coefficients of the ceramic matrix and second phase will promote the bifurcation of the main crack, thus, making the path of the main crack to become tortuous and uneven. It increases the surface energy during the expansion process, thus, hindering the rapid crack expansion and enhancing the toughness of the material, which are the fundamental reasons for the microcrack toughening. Both microcrack and crack deflection toughening result from the residual stresses caused by the mismatch in the thermal expansion coefficient of TiB2 and B4C. Figure 8 shows the phenomenon of the particle pullout, and a hexagonal pit 0.5 µm wide and 1 µm long can be observed. The particle pullout is a typical feature of the intergranular fracture, which can consume a large amount of crack energy, thus, improving the toughness of the composites. Overall, thhe crack deflection, particle pullout, crack microbridging, microcrack toughening and mixed fracture mode are effective means to improve the ceramic toughness. On the other hand, the thermal expansion coefficient difference between particle and ceramic matrix is the main factor determining the toughening effect.
3.3 Mechanical properties
The relative density, Rockwell hardness, bending strength and fracture toughness of Group 1 and Group 2 are shown in Fig. 9. It is observed that the comprehensive mechanical properties of Group 2 are better than Group 1, indicating that the composites containing 30 wt% TiB2 exhibit superior mechanical performance than the composites containing 20 wt% TiB2. This is attributed to an increase in the thermal expansion coefficient of the composite with the amount of TiB2. Specifically, Fig. 9a shows the relative densities of the composites. Compared to the other two components in Group 1, the maximum value is noted to be 98.86% for 70B4C-20TiB2. It indicates that the sintering of the B4C-20 wt% TiB2 composite at 2130 ℃ leads to high density in the case of direct addition of the B4C powder rather than the addition of TiC or TiC and B4C mixed powders. In short, the effect of direct addition of TiB2 to the composite ceramics to improve the relative density of the ceramics is observed to the most effective strategy. Similar to Group 1, the ceramic sintered with the 70B4C-20TiC components exhibited the highest densification in Group 2, with a value of 97.63%. Thus, in terms of the relative density of the B4C-30 wt%TiB2 component, direct addition of the TiC powder is more effective as compared to the addition of B4C or TiC and B4C mixed powders. In other words, TiB2 generated by the in-situ reaction of TiC in the B4C-30 wt%TiB2 system improved the sinterability of the composite ceramics effectively. Figure 9b shows the Rockwell hardness of the Group 1 and Group 2 sintered samples. 70B4C-20TiB2 and 60B4C-30TiB2 exhibited the highest hardness of 85.4HRA and 86.7HRA in each group, respectively. This indicates that the TiB2 particles play a key role in the hardness of the B4C-TiB2 composite ceramics. Moreover, compared to the reaction of TiC with B4C to form TiB2, it is found that TiB2, as the second phase directly added to the composites, can improve the ceramic hardness more effectively. Hence, the hardness of the composites sintered with the pre-designed B4C-TiC-TiB2 component is higher than the composites sintered with the B4C-TiC component.
The bending strength and fracture toughness of the materials are shown in Fig. 9c and Fig. 9d, respectively. Compared with the monolithic B4C ceramics, the toughness of the B4C-TiB2 ceramic composites was greatly improved due to the presence of TiB2. The difference in the bending strength is noted to be consistent with the change in the relative density. Generally, low densification is one of the reasons for the reduction in the bending strength, as the stress concentration tends to occur near the pores, which after a critical value leads to the material failure. The maximum bending strength among the Group 1 materials is noted for 70B4C-20TiB2, whereas the maximum bending strength in Group 2 is observed for 70B4C-20TiC, with values of 308.1 MPa and 291.9 MPa, respectively. The maximum fracture toughness of the B4C–TiB2 composites in this study was observed to be 5.38 MPa·m1/2, which is more than twice the values reported for the monolithic B4C ceramics (2.4–2.6 MPa·m1/2)[7], rarely observed for the pressureless sintered B4C ceramics. Moreover, the fracture toughness of 70B4C-20TiB2 and 60B4C-30TiB2 is much higher than the other two in the group, and the material with the lowest fracture toughness among the samples has higher toughness than the monolithic B4C ceramics. It indicates that the addition of the TiB2 particles can effectively improve the fracture toughness of the B4C–TiB2 composite ceramics, and the effect of the direct addition of TiB2 particles is the most prominent. Comparing the mechanical properties of the samples, it can be observed that the two-phase system has superior properties than the three-phase system, with optimal comprehensive properties achieved by the direct addition of TiB2. Also, the addition of TiB2 or the incorporation of TiC to the B4C ceramics to form TiB2 as the second phase represent two completely different principles. The results proved that the two methods of forming TiB2 did not promote each other and only one method played a major role during the synthesis. Moreover, the direct addition of TiB2 exhibited a more important role than the use of TiC to form TiB2. However, this phenomenon is dependent of the reaction conditions as the particle size, sintering temperature, sintering holding time and other conditions have an impact on the experimental results.
The effect of TiB2 addition methods on the microstructure and mechanical properties of the pressureless sintered B4C-TiB2 composites at 2130 ℃ was investigated. The results indicate that the mechanical properties of the two-phase system are better than the three-phase system. Moreover, the direct addition of the TiB2 powder is noted to be better than the addition of TiC or TiB2 and TiC mixed powders in terms of the comprehensive mechanical properties. The B4C-TiB2 composites with 30 wt% TiB2 added directly exhibit optimal performance, with relative density, hardness, bending strength and fracture toughness values of 94.81%, 86.71HRA, 277.6 MPa, and 5.38 MPa·m1/2, respectively. The toughening mechanisms are observed to be microcrack deflection, particle pullout and crack microbridging.
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Table 1
Composition and nomenclature of the powder mixture.
|
Designation |
B4C/ wt.% |
TiC/ wt.% |
TiB2/wt.% |
C/ wt.% |
Si/ wt.% |
|
|
Sample1
|
80B4C-10TiC |
80 |
10 |
0 |
7 |
3 |
|
70B4C-20TiB2 |
70 |
0 |
20 |
7 |
3 |
|
|
70B4C-6.67TiC-13.33TiB2 |
70 |
6.67 |
13.33 |
7 |
3 |
|
|
Sample2
|
70B4C-20TiC |
70 |
20 |
0 |
7 |
3 |
|
60B4C-30TiB2 |
60 |
0 |
30 |
7 |
3 |
|
|
60B4C-10TiC-20TiB2 |
60 |
10 |
20 |
7 |
3 |
|