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Fig.
1.
Raw powders used in this study: (a) 5 µm B4C; (b) 14 µm SiC.
| Cite this article as: |
Rong-zhen Liu, Wen-wei Gu, Yu Yang, Yuan Lu, Hong-bin Tan, and Jian-feng Yang, Microstructure and mechanical properties of reaction-bonded B4C–SiC composites, Int. J. Miner. Metall. Mater., 28(2021), No. 11, pp.1828-1835. https://dx.doi.org/10.1007/s12613-020-2207-9
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Reaction-bonded B4C–SiC composites are highly promising materials for numerous advanced technological applications. However, their microstructure evolution mechanism remains unclear. Herein, B4C–SiC composites were fabricated through the Si-melt infiltration process. The influences of the sintering time and the B4C content on the mechanical properties, microstructure, and phase evolution were investigated. X-ray diffraction results showed the presence of SiC, boron silicon, boron silicon carbide, and boron carbide. Scanning electron microscopy results showed that with the increase in the boron carbide content, the Si content decreased and the unreacted B4C amount increased when the sintering temperature reached 1650°C and the sintering time reached 1 h. The unreacted B4C diminished with increasing sintering time and temperature when B4C content was lower than 35wt%. Further microstructure analysis showed a transition area between B4C and Si, with the C content marginally higher than in the Si area. This indicates that after the silicon infiltration, the diffusion mechanism was the primary sintering mechanism of the composites. As the diffusion process progressed, the hardness increased. The maximum values of the Vickers hardness, flexural strength, and fracture toughness of the reaction-bonded B4C–SiC ceramic composite with 12wt% B4C content sintered at 1600°C for 0.5 h were about HV 2400, 330 MPa, and 5.2 MPa·m0.5, respectively.
Reaction-bonded B4C–SiC (RBSBC) materials fabricated through the Si infiltration method feature excellent properties, including outstanding hardness, good wear resistance, excellent chemical stability, and low density [1–2]. These materials have applications in many advanced technological fields, such as armor production and laser mirror, wear parts, and nuclear industry parts manufacture [3–5].
The morphology and microstructure of these RBSBC composites play an essential role in manipulating their properties, and several researchers have investigated the microstructure evolution mechanism during the reaction sintering process [6–9]. Han et al. [10] analyzed the influence of the B4C content in RBSBC composites. They identified that the composite material contained only SiC, B4C, and Si phases. After studying the process of Si infiltration into the B4C body, Hayun et al. [11] proposed that dissolution is the principal mechanism of the reaction-bonding process, wherein Si-containing B4C is deposited on the core B4C, so that a core–shell structure is formed. Furthermore, following reaction sintering, B4C, SiC, and Si reach equilibrium, and the structure within the material is dominated by the ternary phases B12(C,Si,B)3 and Si [12]. Zhou et al. [13] investigated the effect of polycarbosilane addition. The B4C–SiC composite with polycarbosilane exhibited a flexural strength of (319 ± 12) MPa, elastic modulus of (402 ± 18) GPa, and hardness of (17.3 ± 0.2) GPa. The authors found that following Si infiltration, the reaction-bonded boron carbide composites contained four phases: B4C, B12(C,Si,B)3, β-SiC, and residual Si. This agrees well with the research results by Hayun et al. [11]. Xu et al. [14] used the gel casting method to form B4C/C preform and fabricated RBSBC composite through Si infiltration. They also found that the composite phases were composed of B–C–Si ternary phases, Si, SiC, and B4C. Song et al. [15] also found that molten Si could react with the outer layer of the B4C particles to form a B–C–Si ternary phase and small amounts of SiC. Wang et al. [16] found that newly formed SiC precipitated on preexisting SiC particles, leading to more B4C dissolution in the RBSBC fabrication process. The above studies reveal that the RBSBC phase is composed of B4C, SiC, Si, and the ternary phase B–C–Si (B12(C,Si,B)3) and that the RBSBC microstructure evolution process is dominated by the dissolution mechanism.
However, Frage et al. [17] pointed out that the Si and C reaction was influenced by the Si atomic ratio (when the Si atomic ratio was lower than 13at% in a Cu–Si melt, no SiC was formed). This indicates that the reaction between B4C and Si may also be influenced by the relative ratio of the raw material. Recently, Dutto et al. [18] fabricated RBSBC via microwave heating and found that the SiC formation reaction was initially fast and then quickly slowed down. The fast stage was mainly dominated by the dissolution precipitation mechanism; subsequently, the B12(C,B,Si)3 covered the B4C grains and carbon. The silicon should diffuse through the B12(C,B,Si)3 phase to continue the reaction; thus, the reaction speed was retarded. The sintering process was dominated by the diffusion process rather than the dissolution process. This indicates that the material structure and performance can be manipulated by changing the sintering mechanism, which can be realized by altering the material composition and sintering method.
In the present study, B4C and SiC were infiltrated by Si under different temperatures, sintering times, and different B4C contents. The mechanical properties were systematically investigated. The microstructure was analyzed via scanning electron microscopy (SEM), energy-dispersive X-ray spectroscopy (EDS), and transmission electron microscopy (TEM). Besides correlating the batch composition, processing parameters, and mechanical properties, the purpose of this work was to assess the microstructure evolution mechanism of RBSBC composites.
Green SiC (14 μm) was provided by Ningxia Orient Tantalum Industry Corporation, China. Nanocarbon black N330 (30 nm) was provided by Tokai Carbon Corporation, Japan. Phenolic resin 2130, which was adopted as the binder and the additional carbon source, was provided by Xi’an Zhenhua Chemicals, China. Boron carbide (B4C, 5 μm) was provided by Mudan Jiang Jingangzuan Boron Carbide Corporation, China. Commercial Si powder (97% purity, 5 mm) was supplied by Ningxia Orient Tantalum Industry Corporation, China. The morphologies of the SiC and B4C powders are displayed in Fig. 1. The batch composition of the composites is presented in Table 1. The carbon density in our study ranged from 0.77 to 0.62 g·cm−3 was adopted to avoid silicon infiltration failure caused by higher carbon density and to observe the effect caused by modifying the B4C and Si ratios.
| Composite No. | SiC | N330 | Phenolic resin | B4C |
| 1 | 68 | 14 | 11 | 7 |
| 2 | 61 | 12 | 15 | 12 |
| 3 | 53 | 10 | 13 | 25 |
| 4 | 44 | 7 | 14 | 35 |
| 5 | 38 | 6 | 11 | 45 |
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The SiC, N330, and B4C powders and phenolic resin were mixed in a high-density polyethylene jar for 8 h at 200 r/min via ball milling using ethanol solvent. Afterward, the powder was dried. A 5 mm × 50 mm mold was adopted to fabricate the specimens, and the powder was compacted at 48, 70, and 96 MPa to evaluate the influence of compaction. After the molding process, the samples were placed in an oven for 8 h at 120°C. When the resin was fully cured, the samples were transferred into a graphite crucible filled with 5 mm Si powder. A self-made vacuum furnace was used for the silicon infiltration process. The furnace pressure was about 10 Pa and was modified using a mechanical vacuum pump. The samples were infiltrated with molten Si at 1550, 1600, and 1650°C under vacuum for 30 min. A 1 h sintering time was used under 1600°C to investigate the influence of increasing the sintering time.
After the samples were sintered, they were polished with 0.25 mm diamond paste for the following test. The mechanical tests were performed using samples of 3 mm × 4 mm × 50 mm and at a loading speed of 0.05 mm/min. The flexural strength was measured via a three-point bending test at room temperature, performed on a universal testing machine (Instron 1195, Instron Co., England), with a span length of 16 mm. A cut with a depth of about 2.5 mm (denoted as a) had been machined in the samples using an inner edge cutter. The cut width was less than 0.2 mm, and the span was 16 mm (denoted as L). The fracture toughness was measured by the single-edge-notched beam method on specimens with the same span of 16 mm. According to the specimen width (W), thickness (b), and applied force (F), notch depth (a), and span (L), the fracture toughness (KIc) of the material was calculated according to the following formula:
|
KIc=3FL2bW2×√a |
(1) |
The hardness of this material was tested under a 2.94 N load through the Vickers hardness method. The density and porosity values were obtained via the Archimedes method using distilled water as the immersion medium. The microstructure was characterized via SEM (VEGA II XMU, Czech Republic) and EDS (Oxford energy dispersive spectrometer, United Kingdom). Furthermore, high-magnification images were obtained via TEM (JEM-3010, Japan), and the inter-granular phases were analyzed via selected-area electron diffraction (SAED). The phase composition of the composites was identified via X-ray diffraction (XRD, X’Pert PRO, Netherlands). The TEM specimens were prepared by cutting and grinding the sintered specimens to a plate of 20 μm thickness and then subjecting them to ion-beam thinning. Five specimens were used for each test, and average values with error bars were obtained.
The SEM microstructures of composites with different B4C contents and sintered under different times were observed (Fig. 2). No significant decrease in residual Si was observed with the increase in B4C content under sintering conditions of 1600°C and 0.5 h (Figs. 2(a) and 2(c)); however, the B4C volume ratio significantly increased, leading to an increase in hardness. After 1 h of sintering, the B4C particles were dissolved into the Si phases (Figs. 2(c) and 2(d)). Although the B4C volume ratio decreased, the Si volume ratio also decreased dramatically. Unreacted B4C was observed in the microstructure with 35wt% B4C content.
The surface morphology of the sample with 7wt% B4C sintered at 1600°C for 1 h showed the presence of cracks (Fig. 3), which could not be eliminated by further polishing. Since these cracks only appeared at the higher sintering temperatures (above 1650°C) or after extended sintering to 1 h, they were not generated during the preparation of the green body via drying or molding. Therefore, the crack occurrence is attributed to the volume expansion due to the reactions between B4C and Si. The mismatch caused by the excessive volume expansion leads to an increase in the number of internal cracks in the material and thus a decrease in the material strength.
X-ray diffraction was used to assess the phase changes with increasing sintering time (Fig. 4). With the increase in sintering time, the diffraction peak intensity for (111) of Si (28.34°) decreased, indicating that the amount of residual Si decreased. The peak intensity of SiC, B8C, and the B–C–Si ternary phase at ~36° increased with the increase in sintering time, indicating that the reactant contents generated by the reaction between B4C and Si increased. These results agree well with those of the hardness test and SEM analysis.
High-resolution transmission electron microscopy (HRTEM) was used to study the interface between the B4C and Si and thus clarify the formation mechanism of SiC grains (Fig. 5). The B4C particles were distributed around the SiC particles, and a clear boundary occurred between Si and B4C (Figs. 5(a) and 5(b)). The HRTEM image of the B4C particle showed a distortion area with cell lattice parameters (d) of 0.460 and 0.272 nm (Fig. 5(c)). The SAED pattern of the B4C crystal showed a clear orthorhombic pattern. This is consistent with the orthorhombic cell lattice parameters of 0.460, 0.272, and 0.231 nm, corresponding to the (021), (1321), and (1300) planes of the B8C crystal (Fig. 5(d)). The B : C atomic ratio agrees well with the EDS results. The C ratio in the raw powder was about two times of the sintered sample, indicating a loss of C during reaction sintering process.
The EDS results of the interface between the B4C rim and Si showed a clear reduction in the Si content in the transition area from B4C to Si (Fig. 6). The transition area width was about ~304 nm. According to the EDS results, the atomic ratio of C in both the Si and the B4C areas was ~16.15at%, whereas that in the transition area was ~18.11at%. Thus, the transition area exhibited a 2% increase in the C content. The Si content in the B4C area was ~2.73at%. This result agrees with that by Wang et al. [19]. Thus, due to the stoichiometric ratio of B4C not fixed, ~2.73at% of the Si atoms replaced the C atoms in the B4C lattice, increasing the C content in the transition area. The C detected in the Si area may be strongly related to the diffusion process. Although the C solubility in Si at 1600°C was only ~0.02%, due to the C diffusion process, the C could continuously dissolve into Si [20]. The C dissolution process in liquid Si is exothermic and can elevate the reaction zone temperature [21]; thus, the dissolution process is promoted due to higher temperature of reaction zone. Therefore, the Si melt with supersaturated C could be obtained with the C diffusion process. Newly formed SiC can also be obtained with the diffusion process. Additionally, with the generation of the transition area, the B4C dissolution was hindered, and the sintering process was dominated by the diffusion process due to the differences in B and C contents between the transition area and the Si area.
According to the above discussion, a mechanism to explain the existence of the transition area and the cause of the diffusion dominating the reaction process is proposed: When the infiltration process begins, dissolution and precipitation occur as the Si wets the B4C surface. Newly formed SiC will precipitate on original SiC particles, as previously observed [16], and a transition area is also generated. Following the formation of the transition area, the C and B atoms will diffuse under the concentration gradient until phase equilibrium is reached. Therefore, although newly formed SiC will preferentially precipitate on the original SiC, due to the existence of the transition area, the subsequent sintering process will be dominated by the diffusion mechanism.
The influence of pressure on the preform density was assessed (Fig. 7). With the increase in pressure, the density increased, reaching ~1.84 g·cm−3 at 96 MPa. When the molding pressure reached 96 MPa, the green body density slightly increased compared with that at 70 MPa, and the possibility for Si infiltration failure was higher. Therefore, 70 MPa was adopted as the molding pressure in the further experiments.
The flexural strengths of the materials at different sintering temperatures (1550, 1600, and 1650°C) were compared (Fig. 8). The strength of the samples first increased with the increase in the B4C content and then decreased, with the maximum strength recorded at 25wt% B4C. At this B4C content, the strength of the sample sintered at 1650°C was ~220 MPa, while that of the sample sintered at 1550°C was ~350 MPa. Due to the effect of silicon infiltration, pores in the green body will be filled up by molten silicon. The porosity of this composite was very low. In this study, the open porosity of sintered material ranged from 0.4% to 0.8%. Near-completely dense samples were obtained through the silicon infiltration process. Therefore, the influence of porosity on strength can be neglected. The difference in strengths at various temperatures was mainly due to the increase in microcracks caused by the volume expansion generated by the reaction of B4C and Si and the grain growth with increasing sintering temperature.
The fracture toughness of the materials first increased with increasing B4C content and then gradually decreased (Fig. 9). At 12wt% B4C, the fracture toughness was the highest, with an average fracture toughness of ~5.25 MPa·m0.5. This trend is similar to that for the influence of B4C content on composite strength; however, the maximum strength was obtained at 25wt% but not 12wt% B4C. The fracture of the RBSBC composites is attributed to the interfacial bonds between the grains. In the reaction-bonded materials, these bonds are strongly correlated to the formation of β-SiC during silicon infiltration. So et al. [22] reported that both the strength and fracture toughness were enhanced because more energy was consumed when cracks propagated through the grains of the composites. Therefore, when the B4C content is lower than 12wt%, with the increase in the B4C content, both the fracture toughness and strength increase. When the B4C content is greater than 12wt%, the internal residual stress caused by the reaction between B4C and Si limits the fracture toughness and strength. As demonstrated in Fig. 8, the strength value obtained at 25wt% B4C was larger than that obtained at 12wt% B4C. The B4C amount has a greater impact on the composite mechanical properties than the residual stress does [23]; therefore, the flexural strength can still increase with the increase in the B4C content until the benefits from the B4C cannot offset the negative effects caused by the increase in internal stress. This explains why the maximum strength was obtained at 25wt% but not 12wt%. This also agrees well with the results shown in Fig. 2.
With the increase in the B4C content, the composite hardness increased (Fig. 10), and it reached HV ~2537 at 25wt% B4C and sintering conditions of 1600°C and 0.5 h. So et al. [24] reported similar trends in SiC–B4C fabricated through hot pressing. In their study, when the holding time was increased to 1 h, the hardness increased to HV ~2600.
The hardness of the composites can be directly calculated by the rule-of-mixtures based on the material phase hardness. To simplify the analysis process herein, the reaction between B4C and Si was neglected at the beginning of the calculation, and the hardness of the composites (Hc) was calculated as follows:
|
Hc=VSi⋅HSi+VSiC⋅HSiC+VB4C⋅HB4C |
(2) |
where VSi, VSiC, and
In this study, B4C–SiC composites with different B4C contents were fabricated through the Si-melt infiltration process in a vacuum. The composites were mainly composed of SiC, B4C, B-rich silicon, and B12(C,Si,B)3. When the sintering time was 0.5 h, unreacted B4C still existed in the samples. When the sintering time was 1 h, the samples exhibited surface cracks. The unreacted B4C will be fully consumed by residual Si when the B4C content is 12wt% but can still be observed when B4C content reaches 35wt%. The reaction mechanisms during the Si infiltration process involve dissolution and precipitation, as well as diffusion. Microstructure and phase analyses showed the presence of a transition area with C content difference. This transition area hindered the dissolution process; thus, although the newly formed SiC preferentially deposited on the original SiC, the sintering process was dominated by the diffusion mechanism. The fracture toughness was maximum (5.2 MPa·m0.5) at 12wt% B4C, and the flexural strength was maximum (360 MPa) at 25wt% B4C. The B4C content showed more influence on mechanical properties than the residual stress did. The hardness exhibited a positive correlation with the B4C content and the sintering time. The increase in hardness is attributed to (1) the effect of the increase in the B4C content and (2) a decrease in the amount of residual Si due to the reaction between B4C and Si.
This work was financially supported by the National Natural Science Foundation of China (No. 51875222), the China Postdoctoral Science Foundation (No. 2017M622426), the First Class Special Funding for Postdoctoral Scientific Research of Hubei Province, China (No. 2017-G3), and the Opening Fund of State key laboratory for Environment-friendly Energy Materials (No. 17kffk 12). The authors are grateful to the State Key Laboratory of Materials Processing and Die & Mould Technology for the mechanical property tests and the Analysis and Testing Center of Huazhong University of Science and Technology for the XRD and SEM tests.
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Figures(10) / Tables(1)
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| Composite No. | SiC | N330 | Phenolic resin | B4C |
| 1 | 68 | 14 | 11 | 7 |
| 2 | 61 | 12 | 15 | 12 |
| 3 | 53 | 10 | 13 | 25 |
| 4 | 44 | 7 | 14 | 35 |
| 5 | 38 | 6 | 11 | 45 |
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