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Effects of boron content on the microstructures and mechanical properties of reactive hot-pressed BxC-TiB2-SiC composites
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Effects of boron content on the microstructures and mechanical properties of reactive hot-pressed BxC-TiB2-SiC composites Qianglong Hea, Jingjing Xiea,∗∗, Aiyang Wanga, Chun Liua, Tian Tiana, Lanxin Hua, Chenhong Yib, Zhixiao Zhangc, Hao Wanga, Weimin Wanga,∗, Zhengyi Fua a
The State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan, 430070, China Institute of Fluid Physics, China Academy of Engineering Physics, Mianyang, 621900, China c College of Materials Science and Engineering, Hebei University of Engineering, Handan, 056038, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: BxC-TiB2-SiC Phase transition Microstructures Mechanical properties Twins
BxC-TiB2-SiC ceramic composites were fabricated via reactive hot pressing using TiC, B, and Si as the raw materials. The phase transition process was studied by heating powder mixtures to different temperatures in combination with X-ray diffraction analysis. The stoichiometric ratio between B and C in boron carbide is variable. A series of powder mixtures containing excess boron (0 wt%, 10 wt%, 20 wt%, or 30 wt% B) were sintered, and the microstructures and mechanical properties of the composites were investigated. The results showed that the B6.1C-TiB2-SiC composite prepared from the starting powders with 30 wt% excess boron had the best comprehensive mechanical properties, with a relative density, hardness, bending strength, and fracture toughness of 98.32%, 33.2 GPa, 840 MPa, and 5.22 MPa m1/2, respectively. Excessive boron substitution may cause lattice distortion in boron carbide, and the boron carbide grains in this state may form a large number of twins under the compressive stress generated by the TiB2 grains, which will affect the properties of the composites.
1. Introduction Boron carbide (B4C) has excellent combinatorial features of low density (2.52 g/cm3), high hardness, high melting point, high Young's modulus, good chemical stability, and excellent neutron absorption ability, so it is widely used for lightweight armour and cutting tools and in the nuclear industry. Nevertheless, its disadvantages of relatively low fracture toughness and poor sinterability limit the application of B4C to some extent [1–6]. Many efforts have been made to improve the sinterability and mechanical properties of B4C ceramics. Results have shown that TiB2 and SiC are the ideal choices as additives for a B4C matrix, because the addition of TiB2 and SiC can maximise the retention of the unique properties of B4C [7–14]. The ternary composite B4C-TiB2-SiC has been investigated by some researchers [15–17], and was usually fabricated using reactive sintering through reactions (1) and (2). B4C + 2TiC + 3Si → 2TiB2 + 3SiC
(1)
3B4C + 2Ti3SiC2 → 6TiB2 + 2SiC + 5C
(2)
∗
Reactive sintering provides a way to obtain powder mixtures with fine particle sizes even from coarse precursors [18]. Tiny TiB2 and SiC particles have been obtained through the above approaches. However, B4C still originates from commercial coarse powders. In order for B4C, TiB2, and SiC all to be fine powders, they must be generated in-situ. Moreover, the stoichiometric ratio between B and C is variable in boron carbides, which have a wide composition range of 8.8%–20% C [19,20]. The results show that ZrB2-BN composites can be prepared by reaction sintering of ZrN-B [21,22], inspired by the nitride boronizing process, in this study, we fabricated BxC-TiB2-SiC ceramic composites via reactive hot pressing using TiC, B, and Si as starting powders with different boron contents through reaction (3). Furthermore, the effects of B content on the microstructures and mechanical properties of BxCTiB2-SiC composites were investigated. TiC + (6 - 4x)B + xSi → (1 - x)B4C + TiB2 + xSiC (0 ˂ x ˂ 1)
(3)
2. Experimental procedures Commercially available TiC powder (d50 = 3 μm, Shanghai Aladdin
Corresponding author. Corresponding author. E-mail addresses: [email protected] (J. Xie), [email protected] (W. Wang).
∗∗
https://doi.org/10.1016/j.ceramint.2019.06.214 Received 18 May 2019; Received in revised form 17 June 2019; Accepted 20 June 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Qianglong He, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.06.214
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Table 1 Characteristics of the raw material powders.
Particle size Purity Oxygen content
Table 2 Volume fractions of B4C, TiB2, and SiC with different x values.
TiC
B
Si
3.0 μm > 99% 0.59%
0.9 μm 95–97% 2.91%
1.0 μm > 99.9% 2.09%
x = 0.1 x = 0.2 x = 0.3 x = 0.4 x = 0.5
B4C (vol%)
TiB2 (vol%)
SiC (vol%)
54.3 49.5 44.5 39.2 33.6
42.3 43.4 44.6 45.9 47.2
3.4 7.1 10.9 14.9 19.2
TBS10, TBS20, and TBS30. The powder mixtures were ball milled with ethanol using SiC balls in a polyethylene cup for 2 h. The mixed slurries were dried at 65 °C in a rotating evaporator (R, Shanghai SENCO Technology Co. Ltd., China), and then at 60 °C for 24 h in a vacuum drying chamber (DZF-6050, Shanghai Jinghong Laboratory Equipment Co. Ltd., China). After drying, the obtained powder mixtures were screened using a 200-mesh sieve to minimise powder agglomeration. Then, the TBS10 powder mixture was loaded into a stainless steel die with a diameter of 20 mm and dry pressed under 5 MPa into a disk with a thickness of approximately 3 mm using a tablet press (TY301-1T, Yuyao Tianyu Machinery Equipment Co. Ltd., China). The disks were placed on graphite paper and heated to different temperatures (900 °C, 1000 °C, 1100 °C, 1200 °C, 1300 °C, or 1400 °C) in a hot-pressing furnace (916G-G Press, Thermal Technology Inc., USA) under vacuum. The heating rate was 20 °C/min and the dwelling time was 1 h. After they had cooled naturally to room temperature, the powder disks were removed and crushed into powder using an agate mortar for characterisation. Afterwards, the TBS0, TBS10, TBS20, and TBS30 powder mixtures were poured into a cylindrical graphite die with an inner diameter of 48 mm. In addition, a 0.2-mm-thick graphite foil lining was used to prevent direct contact between the powder mixtures and the graphite die. The synthesis process was conducted in a hot-pressing furnace (916G-G Press, Thermal Technology Inc., USA) at a temperature of 1950 °C for 1 h under a uniaxial pressure of 30 MPa. The heating rates were 20 °C/min from room temperature to 1400 °C under vacuum and then, after holding for 30 min, 10 °C/min to 1950 °C in an argon atmosphere. The pressure remained at 5 MPa until the temperature reached 1400 °C, and with the temperature rising to 1950 °C, the pressure rose to 30 MPa at a constant rate. The final bulk density of the sintered sample was measured using the Archimedes method. The three-point bending test with a span of 20 mm and a loading rate of 0.5 mm/min was conducted on the specimens (2 mm × 3 mm × 25 mm) to evaluate the bending strength using a ceramic test system; the single-edge notched beam method with a span of 20 mm and a loading rate of 0.05 mm/min was used on the specimens (2 mm × 4 mm × 25 mm, notch width = 0.18 mm, notch depth = 2 mm) to measure the fracture toughness (KIC) with the same ceramic test system (CMT6503, Ji'nan Meitesi Testing Technology Co. Ltd., China). The strength and toughness values were determined from the results averaged from seven bars, respectively. The Vickers hardness was determined with an applied load of 9.8 N and a dwelling time of 15 s on a polished surface (430SVD, Wolpert, USA), and the result was the average of ten measurements. The Gibbs free energies of the reactions in the sintering process were calculated using the chemical thermodynamics simulation software HSC Chemistry 6. Standard Gibbs free energy was adopted as the thermodynamics basic criterion, although the reactions took place in a vacuum condition, this approximation method can still be used to compare different reactions. The oxygen contents were evaluated using an oxygen and nitrogen analyser (TC600, USA). The phase assemblage and microstructure of heat-treated powders and sintered ceramics were investigated using X-ray diffraction (XRD) (Rigaku Ultima III, Japan), scanning electron microscopy (SEM) (Hitachi3400, Japan), and transmission electron microscopy (TEM) (Talos F200S, USA).
Fig. 1. SEM images of the starting materials: (a) TiC, (b) B, and (c) Si.
Biochemical Technology Co., Ltd., Shanghai, China), amorphous B powder (Dandong Chemical Research Institute Co., Ltd., Dandong, China), and Si powder (Shanghai Aladdin Biochemical Technology Co., Ltd., Shanghai, China) were used as the raw materials. Table 1 shows the characteristics of the raw material powders. Fig. 1 shows the scanning electron microscopy (SEM) images of the starting materials. Table 2 shows the volume fractions of B4C, TiB2, and SiC with different x values. In the current study, we chose an x value of 0.3 for formula (3). In this case, the content of SiC is suitable, and the matrix phases of B4C and TiB2 have similar volume fractions, the mutual grain growth inhibiting effect is more pronounced, and the agglomeration of the similar phase particles is less pronounced [23]. Then, a series of powder mixtures containing excess boron (0 wt%, 10 wt%, 20 wt%, or 30 wt% B) were weighed and marked as TBS0, 2
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Fig. 3. Gibbs free energy changes of reactions (4), (5), and (6).
(5)
SiB6 + 2.5C → 1.5B4C + SiC
(6)
The corresponding thermodynamics data of the reaction equations above are shown in Fig. 3. The Gibbs free energies for all reactions between 800 °C and 1400 °C are negative, indicating that all of these reactions could occur during the process. Between 900 °C and 1000 °C, reactions (4) and (5) have similar Gibbs free energies; this is consistent with the disappearance of Si and the detection of trace TiB2 in this temperature range. The Gibbs free energy of reaction (6) is more negative than of reaction (5) when the temperature is higher than 1200 °C, and this is also consistent with the detection of B4C at 1300 °C. Fig. 4(a) shows the SEM image of the morphology of a powder mixture before heating, and Fig. 4(b) shows the corresponding backscattered electron (BSE) image. The large white particles with smooth surfaces are TiC. The particles with a particle size smaller than 1 μm and a rough surface are amorphous B, which are widely distributed in the powder mixture. In addition, the flake particles approximately 1 μm in size are Si. Fig. 4(c) shows the SEM image of the morphology of the powder mixture heated at 1400 °C for 1 h, and Fig. 4(d) shows the corresponding BSE image. Based on the low atomic weight, the particles shown in grey are B4C and SiC, and most of them have a particle size distribution between 200 nm and 300 nm. Meanwhile, the white particles are TiB2, based on the high atomic weight. The size of the TiB2 particles is approximately 150 nm, and they are clustered together to form micron aggregates, as shown in Fig. 4(e) and (f), which are the enlarged images of the rectangular areas in Fig. 4(c) and (d), respectively. The solid phase reaction in this study is a process controlled by diffusion. B atoms diffused into TiC and adopted TiC with a large particle size as a skeleton, and the generated TiB2 particles with a small particle size agglomerated together. The generated TiB2 particles will hinder any further diffusion of B [24]. This is why there was still a small amount of TiC residue in the powder mixture treated at 1400 °C for 1 h. Fig. 5 shows the reaction process of the TiC, B, and Si powder mixture. When the temperature rose to 1000 °C, the first reaction that occurred was between B and Si to form SiB6. The reaction products of B and Si are generally SiB4 or SiB6 [25,26], and the atomic ratio of B and Si in this reaction was 17.6:1, so the product should be SiB6. As the temperature increased, B continued diffusing to TiC instead of C to generate TiB2. The small TiB2 particles formed from the skeleton of the large TiC particles would still aggregate together. Finally, the C replaced by B would diffuse to the outside and react with B and SiB6 to form boron carbide and silicon carbide.
Fig. 2. XRD patterns of the powder mixtures synthesised at different temperatures.
3. Results and discussion 3.1. Phase transition process Fig. 2 shows the effect of temperature on the phase compositions of the powder mixtures. The XRD patterns of powder mixtures before heating were almost identical to those of powder mixtures heated to 900 °C, indicating that no reaction occurred at below 900 °C. When the temperature was increased to 1000 °C, it could be observed clearly that the XRD intensity of the Si phase was almost non-existent. The reason should be that Si reacted with B and formed a SiB6 phase, since weak SiB6 peaks could be detected in the XRD pattern. In order to see the peaks of SiB6 in XRD pattern more clearly, the enlarged version was shown in the upper right corner of Fig. 2(a). In addition, traces of TiB2 could be found at this temperature. As the temperature rose to 1100 °C, apparent peaks of TiB2 could be seen, and the peaks of Si disappeared completely. When the temperature was increased to 1200 °C, the XRD intensity of TiC diminished, and the intensity of TiB2 increased significantly. The intensity of TiB2 increased further, and the intensity of TiC decreased further, as the temperature rose to 1300 °C; furthermore, a B4C phase was generated at this temperature. As the temperature continued to rise to 1400 °C, a β-SiC phase was detected, and there was only a small amount of TiC left. The reactions which occurred during the process are shown in Eq. (4)–(6): TiC + 2B → TiB2 + C
Si + 6B → SiB6
(4) 3
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Fig. 4. SEM and BSE images of (a), (b) TBS10 powder mixture before heating and (c), (d) after heating at 1400 °C, and (e), (f) partial enlarged view.
B contents sintered at 1950 °C. It can be observed that as the B content increased, the peak of BxC shifted significantly to a smaller angle. The C atoms in the icosahedron of B11C in the crystal structure of B4C were replaced by B atoms; when the larger B atoms replaced the smaller C atoms, the peak of B4C shifted to the smaller angle [27,28]. As C atom substitution increased, the peak deviation of boron carbide increased. The inter-plane distances (d) and lattice parameters (a and c) in the hexagonal lattice were calculated on the basis of diffraction angles and corresponding lattice planes (hkl) using the equation
3.2. Effects of boron content on the microstructure and mechanical properties Table 3 shows the compositions and oxygen contents of the sintered samples. As the B content increased, the ratio of boron to carbon in boron carbide increased. The samples with 0 wt%, 10 wt%, 20 wt%, and 30 wt% excess boron had theoretical stoichiometries of B4C, B4.7C, B5.4C, and B6.1C. The increases in B content also resulted in increased boron carbide volume fractions in the composites. After the powder mixtures were sintered, the oxygen contents of TBS0 and TBS10 decreased to the greatest extent, while those of TBS20 and TBS30 decreased to lesser extents. The reasons will be divulged later. Fig. 6 shows the XRD patterns of the powder mixtures with different
l2
dhkl = ⎡ 2 (h2 + k 2 + hk ) + 2 ]−1/2 [27]. After the values of a and c c ⎢ 3a ⎣ were calculated, the carbon content in boron carbide (BxC) could be determined by comparing with the results of previous studies [29]. 4
Fig. 5. Schematic diagram of solid phase reaction of the powder mixture. 4
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Table 3 Compositions and oxygen contents of the sintered samples.
TBS0 TBS10 TBS20 TBS30
Stoichiometry of BxC
BxC (vol%)
TiB2 (vol%)
SiC (vol%)
Oxygen content before sintering (wt%)
Oxygen content after sintering (wt%)
B4C B4.7C B5.4C B6.1C
44.5 47.9 50.9 53.6
44.6 41.9 39.5 37.3
10.9 10.2 9.6 9.1
1.70 1.75 1.79 1.84
0.38 0.41 1.26 1.44
Table 4 Mechanical properties of TBS composites.
TBS0 TBS10 TBS20 TBS30
Relative density (%)
Hardness (GPa)
Bending strength (MPa)
Fracture toughness (MPa·m1/2)
99.12 98.96 98.35 98.32
29.6 30.8 32.3 33.2
587 730 771 840
6.63 6.42 5.64 5.22
± ± ± ±
0.64 0.92 1.20 0.77
± ± ± ±
50 47 39 41
± ± ± ±
0.20 0.32 0.18 0.19
appeared as the dark phase and formed a continuous matrix; TiB2 appeared as the white phase and SiC appeared as the grey phase. Clearly, the fracture mode of TiB2 is mainly transgranular fracture, and the fracture modes of BxC and SiC are mainly intergranular fracture. The mixed fracture mode is favourable for the toughness of the composites. The line-interception method was used to measure the grain size of TiB2 in the composites; more than 100 grains were used in each sample. The results showed that the grain sizes of TiB2 in TBS0, TBS10, TBS20, and TBS30 were 2.46 μm, 2.20 μm, 1.97 μm, and 1.85 μm, respectively. Table 4 shows the mechanical properties of the TBS composites; it can be seen that all samples have high relative densities. This is due to the high sintering activity of the mixed powders produced by reaction (3). Fine particles came from the in-situ reaction process, resulting in more surface energy and a larger sintering driving force. Moreover, fine particles could reduce the atomic diffusion distance. The mechanisms above could accelerate the sintering process [30]. It has been reported that excess B contributes to sintering [27], but TBS0 and TBS10 have higher relative densities than TBS20 and TBS30. As mentioned in Table 2, the four powder mixtures have similar oxygen contents. However, after sintering, the oxygen contents of TBS0 and TBS10 were
Fig. 6. XRD patterns of the powder mixtures with different B contents sintered at 1950 °C.
Then, with further calculations, the stoichiometric ratios of boron carbide (BxC) in TBS30, TBS20 and TBS10 were determined to be B5.7C, B5.1C and B4.4C, respectively. This result is somewhat inconsistent with the theoretical stoichiometric ratios of BxC for the experiment. The above results were caused by the relatively low purity of the raw boron powder and the diffusion of carbon atoms in the sintering environment into the powder mixtures. The fracture surfaces of the TBS composites are shown in Fig. 7. B4C
Fig. 7. BSE images of fracture surfaces: (a) TBS0, (b) TBS10, (c) TBS20, and (d) TBS30. 5
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(d) show the selected area electron diffraction (SAED) images of TiB2, SiC, and BxC respectively. The three adjacent grains in Fig. 9(a) can be identified as TiB2, SiC, and BxC from Fig. 9(b)-(d). It can be observed that the grain boundaries of BxC-TiB2, BxC-SiC, and TiB2-SiC are narrow, and a contrast between the two sides of the grain boundaries is observed, indicating clean grain boundaries and good interfacial compatibility, which are necessary for a composite to have high flexural strength. In addition, it can be observed that there are some straight lines with consistent orientation in B6.1C grains, and this phenomenon was not observed in TiB2 and SiC grains. Fig. 9(b) and (c) show the electron diffraction patterns of TiB2 and SiC, respectively. The [0‾1 0] zone axis diffraction spots of TiB2 and the [0 0 1] zone axis diffraction spots of SiC were observed. The diffraction spots of TiB2 and SiC are arranged in strict periodicity and no spots deviate from the original structure position. This indicates that the internal structures of TiB2 and SiC grains in the samples are complete and there are almost no surface defects. Fig. 9(d) shows the electron diffraction patterns of B6.1C, which is clearly the electron diffraction pattern of a twin structure, and the twin plane is (101). The quadrilateral formed by white lines and the quadrilateral formed by red lines in the figure have mirror symmetry with (101) as the axis. Therefore, the straight lines with consistent orientation in the B6.1C grains in Fig. 9(a) originated from the twin structure. Fig. 9(e) shows the high-resolution TEM (HRTEM) image of the twin structure in the B6.1C grain; the twin bands have various widths between 0.5 nm and 15 nm. Fig. 9(f) shows the HRTEM images of TiB2, B6.1C, and their interface; there are no obvious deformations or defects at the interface. The reason for this structural feature should be that, during the sintering process for TBS30, excess B would replace the C in the B11C 12-atom icosahedra of nominal stoichiometry B4C, and these substitutions would cause lattice structure distortion in boron carbide. Meanwhile, TiB2 has a higher thermal expansion coefficient, which will cause greater compressive stress on the surrounding grains. Under the action of both compressive stress and lattice distortion, the crystal plane in boron carbide is prone to displacement, so there are high-density twins in the boron carbide grains. Then, the stress is released, and defects do not form easily at the grain boundaries of TiB2 and BxC. Furthermore, a large number of twins in boron carbide means an abundance of subgrain boundaries, which enhances the hardness and strength of the composite. Fig. 9(g) shows the HRTEM image of the interface between TiB2 and SiC. The crystal structure of SiC is complete without obvious defects. However, there is some distortion and collapse around the grain boundary between TiB2 and SiC, which is caused by the failure to release the stress produced by thermal expansion. Therefore, TiB2 is weakly bound to SiC at the grain boundary where microcracks are easily formed. Fig. 10(a) shows the dark-field phase TEM image of intragranular phases in TiB2 grains of TBS30. The size of the intragranular phases is approximately 200 nm, which proved again that the BxC synthesised through the in situ reaction has a very fine particle size. Fig. 10(b) and (c) show that the intragranular phase region is enriched in C, but not obviously enriched in B. This indicates that the boron carbide was engulfed by TiB2 grains at a lower temperature before the formation of a boron-rich phase. Fig. 10(d) shows the bright-field phase TEM image of intragranular phases in TiB2 grains. The size of the intragranular phases is approximately 500 nm. Fig. 10(e) and (f) show that the intragranular phase region is enriched in both C and B, indicating that the boron carbide was engulfed by TiB2 grains at a higher temperature after the formation of a boron-rich phase. These boron carbide intragranular phases may induce the transgranular fracture of TiB2 grains.
significantly lower than those of TBS20 and TBS30. The reasons for this result are as follows: the purity of the raw material B powder was not high, and the C atoms in the sintering environment would diffuse into the powder mixtures and react with B. This created a certain amount of residual C, which would react with the main source of oxygen-B2O3 at approximately 1600 °C which had been confirmed in our previous study [31]. Then, the oxygen contents in TBS0 and TBS10 would decrease. In TBS20 and TBS30, C was consumed by excess B, and B2O3 could be removed and was detrimental for the densification [23]. As the B content increased, the volume fraction of BxC in the composites increased, so the hardness also increased, as shown in Table 3. The bending strength of TBS0 is only 587 MPa, and it was considered that the residual C has a serious effect on the strength. As the B content increased, the volume fraction of BxC increased, causing the grain size of TiB2 to decrease and the dispersion uniformity to improve. According to Griffith's microcrack theory, the size of an inherent crack in a material depends on the grain size to a great extent. In addition, the thermal expansion coefficient of TiB2 is much higher than those of B4C and SiC, so in the BxC-TiB2-SiC system, microcracks are mainly generated at the grain boundary between TiB2 and B4C or SiC. Therefore, a smaller TiB2 grain size leads to a smaller initial crack size and, correspondingly, a higher bending strength of the material [32]. TBS0 had the highest fracture toughness of 6.63 MPa m1/2 because it had the largest grain size and volume fraction of TiB2, which makes crack deflection easier and more frequent. Fig. 8(a) shows the crack growth path of TBS0 and Fig. 8(b) shows the crack growth path of TBS30. It is obvious that numerous crack deflections occurred during the crack propagation process in TBS0. However, crack deflection only occurs in the region where TiB2 grains are concentrated in TBS30. The frequency of crack deflection varies with the TiB2 content, so the toughness of TBS0 (6.63 MPa m1/2) is higher than that of TBS30 (5.22 MPa m1/2). 3.3. Analysis of microstructure Fig. 9(a) shows the dark-field TEM image of TBS30, and Fig. 9(b)-
4. Conclusions BxC-TiB2-SiC ceramic composites were fabricated via reactive hot pressing using TiC, B, and Si as raw materials. The composite powders of B4C, TiB2, and SiC could be obtained by heating at 1400 °C for 1 h. The powder synthesised in situ is nanometre sized and has good
Fig. 8. SEM images of the crack propagations of (a) TBS0 and (b) TBS30. 6
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Fig. 9. (a) TEM image of TBS30; (b)–(d) SAED images of TiB2, SiC, and BxC; (e)–(g) HRTEM images of BxC, BxC-TiB2, and TiB2-SiC.
Fig. 10. (a), (d) TEM image of intragranular phase in TiB2 of TBS30; (b), (e) C element distribution; and (c), (f) B element distribution.
Talent of Higher Learning Institutions of Hebei Province (No. BJ2019054).
sintering activity. In addition, excess B can form boron-rich boron carbide, and the grain contains a large number of twins, which is beneficial to the mechanical properties of the composite. Finally, the TBS30 composite had excellent comprehensive mechanical properties, with a relative density, hardness, bending strength, and fracture toughness of 98.32%, 33.2 GPa, 840 MPa, and 5.22 MPa m1/2, respectively.
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Acknowledgment The authors acknowledge the financial support from the National Key Research and Development Plan of China (2017YFB0310400), the National Natural Science Foundation of China (5167020705), the State Key Laboratory of Advanced Technology for Materials Synthesis and Processing (No. 2019-KF-15), and the Program for the Youth Top-notch 7
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Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Effects of boron content on the microstructures and mechanical properties of reactive hot-pressed BxC-TiB2-SiC composites Qianglong Hea, Jingjing Xiea,∗∗, Aiyang Wanga, Chun Liua, Tian Tiana, Lanxin Hua, Chenhong Yib, Zhixiao Zhangc, Hao Wanga, Weimin Wanga,∗, Zhengyi Fua a
The State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan, 430070, China Institute of Fluid Physics, China Academy of Engineering Physics, Mianyang, 621900, China c College of Materials Science and Engineering, Hebei University of Engineering, Handan, 056038, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: BxC-TiB2-SiC Phase transition Microstructures Mechanical properties Twins
BxC-TiB2-SiC ceramic composites were fabricated via reactive hot pressing using TiC, B, and Si as the raw materials. The phase transition process was studied by heating powder mixtures to different temperatures in combination with X-ray diffraction analysis. The stoichiometric ratio between B and C in boron carbide is variable. A series of powder mixtures containing excess boron (0 wt%, 10 wt%, 20 wt%, or 30 wt% B) were sintered, and the microstructures and mechanical properties of the composites were investigated. The results showed that the B6.1C-TiB2-SiC composite prepared from the starting powders with 30 wt% excess boron had the best comprehensive mechanical properties, with a relative density, hardness, bending strength, and fracture toughness of 98.32%, 33.2 GPa, 840 MPa, and 5.22 MPa m1/2, respectively. Excessive boron substitution may cause lattice distortion in boron carbide, and the boron carbide grains in this state may form a large number of twins under the compressive stress generated by the TiB2 grains, which will affect the properties of the composites.
1. Introduction Boron carbide (B4C) has excellent combinatorial features of low density (2.52 g/cm3), high hardness, high melting point, high Young's modulus, good chemical stability, and excellent neutron absorption ability, so it is widely used for lightweight armour and cutting tools and in the nuclear industry. Nevertheless, its disadvantages of relatively low fracture toughness and poor sinterability limit the application of B4C to some extent [1–6]. Many efforts have been made to improve the sinterability and mechanical properties of B4C ceramics. Results have shown that TiB2 and SiC are the ideal choices as additives for a B4C matrix, because the addition of TiB2 and SiC can maximise the retention of the unique properties of B4C [7–14]. The ternary composite B4C-TiB2-SiC has been investigated by some researchers [15–17], and was usually fabricated using reactive sintering through reactions (1) and (2). B4C + 2TiC + 3Si → 2TiB2 + 3SiC
(1)
3B4C + 2Ti3SiC2 → 6TiB2 + 2SiC + 5C
(2)
∗
Reactive sintering provides a way to obtain powder mixtures with fine particle sizes even from coarse precursors [18]. Tiny TiB2 and SiC particles have been obtained through the above approaches. However, B4C still originates from commercial coarse powders. In order for B4C, TiB2, and SiC all to be fine powders, they must be generated in-situ. Moreover, the stoichiometric ratio between B and C is variable in boron carbides, which have a wide composition range of 8.8%–20% C [19,20]. The results show that ZrB2-BN composites can be prepared by reaction sintering of ZrN-B [21,22], inspired by the nitride boronizing process, in this study, we fabricated BxC-TiB2-SiC ceramic composites via reactive hot pressing using TiC, B, and Si as starting powders with different boron contents through reaction (3). Furthermore, the effects of B content on the microstructures and mechanical properties of BxCTiB2-SiC composites were investigated. TiC + (6 - 4x)B + xSi → (1 - x)B4C + TiB2 + xSiC (0 ˂ x ˂ 1)
(3)
2. Experimental procedures Commercially available TiC powder (d50 = 3 μm, Shanghai Aladdin
Corresponding author. Corresponding author. E-mail addresses: [email protected] (J. Xie), [email protected] (W. Wang).
∗∗
https://doi.org/10.1016/j.ceramint.2019.06.214 Received 18 May 2019; Received in revised form 17 June 2019; Accepted 20 June 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Qianglong He, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.06.214
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Table 1 Characteristics of the raw material powders.
Particle size Purity Oxygen content
Table 2 Volume fractions of B4C, TiB2, and SiC with different x values.
TiC
B
Si
3.0 μm > 99% 0.59%
0.9 μm 95–97% 2.91%
1.0 μm > 99.9% 2.09%
x = 0.1 x = 0.2 x = 0.3 x = 0.4 x = 0.5
B4C (vol%)
TiB2 (vol%)
SiC (vol%)
54.3 49.5 44.5 39.2 33.6
42.3 43.4 44.6 45.9 47.2
3.4 7.1 10.9 14.9 19.2
TBS10, TBS20, and TBS30. The powder mixtures were ball milled with ethanol using SiC balls in a polyethylene cup for 2 h. The mixed slurries were dried at 65 °C in a rotating evaporator (R, Shanghai SENCO Technology Co. Ltd., China), and then at 60 °C for 24 h in a vacuum drying chamber (DZF-6050, Shanghai Jinghong Laboratory Equipment Co. Ltd., China). After drying, the obtained powder mixtures were screened using a 200-mesh sieve to minimise powder agglomeration. Then, the TBS10 powder mixture was loaded into a stainless steel die with a diameter of 20 mm and dry pressed under 5 MPa into a disk with a thickness of approximately 3 mm using a tablet press (TY301-1T, Yuyao Tianyu Machinery Equipment Co. Ltd., China). The disks were placed on graphite paper and heated to different temperatures (900 °C, 1000 °C, 1100 °C, 1200 °C, 1300 °C, or 1400 °C) in a hot-pressing furnace (916G-G Press, Thermal Technology Inc., USA) under vacuum. The heating rate was 20 °C/min and the dwelling time was 1 h. After they had cooled naturally to room temperature, the powder disks were removed and crushed into powder using an agate mortar for characterisation. Afterwards, the TBS0, TBS10, TBS20, and TBS30 powder mixtures were poured into a cylindrical graphite die with an inner diameter of 48 mm. In addition, a 0.2-mm-thick graphite foil lining was used to prevent direct contact between the powder mixtures and the graphite die. The synthesis process was conducted in a hot-pressing furnace (916G-G Press, Thermal Technology Inc., USA) at a temperature of 1950 °C for 1 h under a uniaxial pressure of 30 MPa. The heating rates were 20 °C/min from room temperature to 1400 °C under vacuum and then, after holding for 30 min, 10 °C/min to 1950 °C in an argon atmosphere. The pressure remained at 5 MPa until the temperature reached 1400 °C, and with the temperature rising to 1950 °C, the pressure rose to 30 MPa at a constant rate. The final bulk density of the sintered sample was measured using the Archimedes method. The three-point bending test with a span of 20 mm and a loading rate of 0.5 mm/min was conducted on the specimens (2 mm × 3 mm × 25 mm) to evaluate the bending strength using a ceramic test system; the single-edge notched beam method with a span of 20 mm and a loading rate of 0.05 mm/min was used on the specimens (2 mm × 4 mm × 25 mm, notch width = 0.18 mm, notch depth = 2 mm) to measure the fracture toughness (KIC) with the same ceramic test system (CMT6503, Ji'nan Meitesi Testing Technology Co. Ltd., China). The strength and toughness values were determined from the results averaged from seven bars, respectively. The Vickers hardness was determined with an applied load of 9.8 N and a dwelling time of 15 s on a polished surface (430SVD, Wolpert, USA), and the result was the average of ten measurements. The Gibbs free energies of the reactions in the sintering process were calculated using the chemical thermodynamics simulation software HSC Chemistry 6. Standard Gibbs free energy was adopted as the thermodynamics basic criterion, although the reactions took place in a vacuum condition, this approximation method can still be used to compare different reactions. The oxygen contents were evaluated using an oxygen and nitrogen analyser (TC600, USA). The phase assemblage and microstructure of heat-treated powders and sintered ceramics were investigated using X-ray diffraction (XRD) (Rigaku Ultima III, Japan), scanning electron microscopy (SEM) (Hitachi3400, Japan), and transmission electron microscopy (TEM) (Talos F200S, USA).
Fig. 1. SEM images of the starting materials: (a) TiC, (b) B, and (c) Si.
Biochemical Technology Co., Ltd., Shanghai, China), amorphous B powder (Dandong Chemical Research Institute Co., Ltd., Dandong, China), and Si powder (Shanghai Aladdin Biochemical Technology Co., Ltd., Shanghai, China) were used as the raw materials. Table 1 shows the characteristics of the raw material powders. Fig. 1 shows the scanning electron microscopy (SEM) images of the starting materials. Table 2 shows the volume fractions of B4C, TiB2, and SiC with different x values. In the current study, we chose an x value of 0.3 for formula (3). In this case, the content of SiC is suitable, and the matrix phases of B4C and TiB2 have similar volume fractions, the mutual grain growth inhibiting effect is more pronounced, and the agglomeration of the similar phase particles is less pronounced [23]. Then, a series of powder mixtures containing excess boron (0 wt%, 10 wt%, 20 wt%, or 30 wt% B) were weighed and marked as TBS0, 2
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Fig. 3. Gibbs free energy changes of reactions (4), (5), and (6).
(5)
SiB6 + 2.5C → 1.5B4C + SiC
(6)
The corresponding thermodynamics data of the reaction equations above are shown in Fig. 3. The Gibbs free energies for all reactions between 800 °C and 1400 °C are negative, indicating that all of these reactions could occur during the process. Between 900 °C and 1000 °C, reactions (4) and (5) have similar Gibbs free energies; this is consistent with the disappearance of Si and the detection of trace TiB2 in this temperature range. The Gibbs free energy of reaction (6) is more negative than of reaction (5) when the temperature is higher than 1200 °C, and this is also consistent with the detection of B4C at 1300 °C. Fig. 4(a) shows the SEM image of the morphology of a powder mixture before heating, and Fig. 4(b) shows the corresponding backscattered electron (BSE) image. The large white particles with smooth surfaces are TiC. The particles with a particle size smaller than 1 μm and a rough surface are amorphous B, which are widely distributed in the powder mixture. In addition, the flake particles approximately 1 μm in size are Si. Fig. 4(c) shows the SEM image of the morphology of the powder mixture heated at 1400 °C for 1 h, and Fig. 4(d) shows the corresponding BSE image. Based on the low atomic weight, the particles shown in grey are B4C and SiC, and most of them have a particle size distribution between 200 nm and 300 nm. Meanwhile, the white particles are TiB2, based on the high atomic weight. The size of the TiB2 particles is approximately 150 nm, and they are clustered together to form micron aggregates, as shown in Fig. 4(e) and (f), which are the enlarged images of the rectangular areas in Fig. 4(c) and (d), respectively. The solid phase reaction in this study is a process controlled by diffusion. B atoms diffused into TiC and adopted TiC with a large particle size as a skeleton, and the generated TiB2 particles with a small particle size agglomerated together. The generated TiB2 particles will hinder any further diffusion of B [24]. This is why there was still a small amount of TiC residue in the powder mixture treated at 1400 °C for 1 h. Fig. 5 shows the reaction process of the TiC, B, and Si powder mixture. When the temperature rose to 1000 °C, the first reaction that occurred was between B and Si to form SiB6. The reaction products of B and Si are generally SiB4 or SiB6 [25,26], and the atomic ratio of B and Si in this reaction was 17.6:1, so the product should be SiB6. As the temperature increased, B continued diffusing to TiC instead of C to generate TiB2. The small TiB2 particles formed from the skeleton of the large TiC particles would still aggregate together. Finally, the C replaced by B would diffuse to the outside and react with B and SiB6 to form boron carbide and silicon carbide.
Fig. 2. XRD patterns of the powder mixtures synthesised at different temperatures.
3. Results and discussion 3.1. Phase transition process Fig. 2 shows the effect of temperature on the phase compositions of the powder mixtures. The XRD patterns of powder mixtures before heating were almost identical to those of powder mixtures heated to 900 °C, indicating that no reaction occurred at below 900 °C. When the temperature was increased to 1000 °C, it could be observed clearly that the XRD intensity of the Si phase was almost non-existent. The reason should be that Si reacted with B and formed a SiB6 phase, since weak SiB6 peaks could be detected in the XRD pattern. In order to see the peaks of SiB6 in XRD pattern more clearly, the enlarged version was shown in the upper right corner of Fig. 2(a). In addition, traces of TiB2 could be found at this temperature. As the temperature rose to 1100 °C, apparent peaks of TiB2 could be seen, and the peaks of Si disappeared completely. When the temperature was increased to 1200 °C, the XRD intensity of TiC diminished, and the intensity of TiB2 increased significantly. The intensity of TiB2 increased further, and the intensity of TiC decreased further, as the temperature rose to 1300 °C; furthermore, a B4C phase was generated at this temperature. As the temperature continued to rise to 1400 °C, a β-SiC phase was detected, and there was only a small amount of TiC left. The reactions which occurred during the process are shown in Eq. (4)–(6): TiC + 2B → TiB2 + C
Si + 6B → SiB6
(4) 3
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Fig. 4. SEM and BSE images of (a), (b) TBS10 powder mixture before heating and (c), (d) after heating at 1400 °C, and (e), (f) partial enlarged view.
B contents sintered at 1950 °C. It can be observed that as the B content increased, the peak of BxC shifted significantly to a smaller angle. The C atoms in the icosahedron of B11C in the crystal structure of B4C were replaced by B atoms; when the larger B atoms replaced the smaller C atoms, the peak of B4C shifted to the smaller angle [27,28]. As C atom substitution increased, the peak deviation of boron carbide increased. The inter-plane distances (d) and lattice parameters (a and c) in the hexagonal lattice were calculated on the basis of diffraction angles and corresponding lattice planes (hkl) using the equation
3.2. Effects of boron content on the microstructure and mechanical properties Table 3 shows the compositions and oxygen contents of the sintered samples. As the B content increased, the ratio of boron to carbon in boron carbide increased. The samples with 0 wt%, 10 wt%, 20 wt%, and 30 wt% excess boron had theoretical stoichiometries of B4C, B4.7C, B5.4C, and B6.1C. The increases in B content also resulted in increased boron carbide volume fractions in the composites. After the powder mixtures were sintered, the oxygen contents of TBS0 and TBS10 decreased to the greatest extent, while those of TBS20 and TBS30 decreased to lesser extents. The reasons will be divulged later. Fig. 6 shows the XRD patterns of the powder mixtures with different
l2
dhkl = ⎡ 2 (h2 + k 2 + hk ) + 2 ]−1/2 [27]. After the values of a and c c ⎢ 3a ⎣ were calculated, the carbon content in boron carbide (BxC) could be determined by comparing with the results of previous studies [29]. 4
Fig. 5. Schematic diagram of solid phase reaction of the powder mixture. 4
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Table 3 Compositions and oxygen contents of the sintered samples.
TBS0 TBS10 TBS20 TBS30
Stoichiometry of BxC
BxC (vol%)
TiB2 (vol%)
SiC (vol%)
Oxygen content before sintering (wt%)
Oxygen content after sintering (wt%)
B4C B4.7C B5.4C B6.1C
44.5 47.9 50.9 53.6
44.6 41.9 39.5 37.3
10.9 10.2 9.6 9.1
1.70 1.75 1.79 1.84
0.38 0.41 1.26 1.44
Table 4 Mechanical properties of TBS composites.
TBS0 TBS10 TBS20 TBS30
Relative density (%)
Hardness (GPa)
Bending strength (MPa)
Fracture toughness (MPa·m1/2)
99.12 98.96 98.35 98.32
29.6 30.8 32.3 33.2
587 730 771 840
6.63 6.42 5.64 5.22
± ± ± ±
0.64 0.92 1.20 0.77
± ± ± ±
50 47 39 41
± ± ± ±
0.20 0.32 0.18 0.19
appeared as the dark phase and formed a continuous matrix; TiB2 appeared as the white phase and SiC appeared as the grey phase. Clearly, the fracture mode of TiB2 is mainly transgranular fracture, and the fracture modes of BxC and SiC are mainly intergranular fracture. The mixed fracture mode is favourable for the toughness of the composites. The line-interception method was used to measure the grain size of TiB2 in the composites; more than 100 grains were used in each sample. The results showed that the grain sizes of TiB2 in TBS0, TBS10, TBS20, and TBS30 were 2.46 μm, 2.20 μm, 1.97 μm, and 1.85 μm, respectively. Table 4 shows the mechanical properties of the TBS composites; it can be seen that all samples have high relative densities. This is due to the high sintering activity of the mixed powders produced by reaction (3). Fine particles came from the in-situ reaction process, resulting in more surface energy and a larger sintering driving force. Moreover, fine particles could reduce the atomic diffusion distance. The mechanisms above could accelerate the sintering process [30]. It has been reported that excess B contributes to sintering [27], but TBS0 and TBS10 have higher relative densities than TBS20 and TBS30. As mentioned in Table 2, the four powder mixtures have similar oxygen contents. However, after sintering, the oxygen contents of TBS0 and TBS10 were
Fig. 6. XRD patterns of the powder mixtures with different B contents sintered at 1950 °C.
Then, with further calculations, the stoichiometric ratios of boron carbide (BxC) in TBS30, TBS20 and TBS10 were determined to be B5.7C, B5.1C and B4.4C, respectively. This result is somewhat inconsistent with the theoretical stoichiometric ratios of BxC for the experiment. The above results were caused by the relatively low purity of the raw boron powder and the diffusion of carbon atoms in the sintering environment into the powder mixtures. The fracture surfaces of the TBS composites are shown in Fig. 7. B4C
Fig. 7. BSE images of fracture surfaces: (a) TBS0, (b) TBS10, (c) TBS20, and (d) TBS30. 5
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(d) show the selected area electron diffraction (SAED) images of TiB2, SiC, and BxC respectively. The three adjacent grains in Fig. 9(a) can be identified as TiB2, SiC, and BxC from Fig. 9(b)-(d). It can be observed that the grain boundaries of BxC-TiB2, BxC-SiC, and TiB2-SiC are narrow, and a contrast between the two sides of the grain boundaries is observed, indicating clean grain boundaries and good interfacial compatibility, which are necessary for a composite to have high flexural strength. In addition, it can be observed that there are some straight lines with consistent orientation in B6.1C grains, and this phenomenon was not observed in TiB2 and SiC grains. Fig. 9(b) and (c) show the electron diffraction patterns of TiB2 and SiC, respectively. The [0‾1 0] zone axis diffraction spots of TiB2 and the [0 0 1] zone axis diffraction spots of SiC were observed. The diffraction spots of TiB2 and SiC are arranged in strict periodicity and no spots deviate from the original structure position. This indicates that the internal structures of TiB2 and SiC grains in the samples are complete and there are almost no surface defects. Fig. 9(d) shows the electron diffraction patterns of B6.1C, which is clearly the electron diffraction pattern of a twin structure, and the twin plane is (101). The quadrilateral formed by white lines and the quadrilateral formed by red lines in the figure have mirror symmetry with (101) as the axis. Therefore, the straight lines with consistent orientation in the B6.1C grains in Fig. 9(a) originated from the twin structure. Fig. 9(e) shows the high-resolution TEM (HRTEM) image of the twin structure in the B6.1C grain; the twin bands have various widths between 0.5 nm and 15 nm. Fig. 9(f) shows the HRTEM images of TiB2, B6.1C, and their interface; there are no obvious deformations or defects at the interface. The reason for this structural feature should be that, during the sintering process for TBS30, excess B would replace the C in the B11C 12-atom icosahedra of nominal stoichiometry B4C, and these substitutions would cause lattice structure distortion in boron carbide. Meanwhile, TiB2 has a higher thermal expansion coefficient, which will cause greater compressive stress on the surrounding grains. Under the action of both compressive stress and lattice distortion, the crystal plane in boron carbide is prone to displacement, so there are high-density twins in the boron carbide grains. Then, the stress is released, and defects do not form easily at the grain boundaries of TiB2 and BxC. Furthermore, a large number of twins in boron carbide means an abundance of subgrain boundaries, which enhances the hardness and strength of the composite. Fig. 9(g) shows the HRTEM image of the interface between TiB2 and SiC. The crystal structure of SiC is complete without obvious defects. However, there is some distortion and collapse around the grain boundary between TiB2 and SiC, which is caused by the failure to release the stress produced by thermal expansion. Therefore, TiB2 is weakly bound to SiC at the grain boundary where microcracks are easily formed. Fig. 10(a) shows the dark-field phase TEM image of intragranular phases in TiB2 grains of TBS30. The size of the intragranular phases is approximately 200 nm, which proved again that the BxC synthesised through the in situ reaction has a very fine particle size. Fig. 10(b) and (c) show that the intragranular phase region is enriched in C, but not obviously enriched in B. This indicates that the boron carbide was engulfed by TiB2 grains at a lower temperature before the formation of a boron-rich phase. Fig. 10(d) shows the bright-field phase TEM image of intragranular phases in TiB2 grains. The size of the intragranular phases is approximately 500 nm. Fig. 10(e) and (f) show that the intragranular phase region is enriched in both C and B, indicating that the boron carbide was engulfed by TiB2 grains at a higher temperature after the formation of a boron-rich phase. These boron carbide intragranular phases may induce the transgranular fracture of TiB2 grains.
significantly lower than those of TBS20 and TBS30. The reasons for this result are as follows: the purity of the raw material B powder was not high, and the C atoms in the sintering environment would diffuse into the powder mixtures and react with B. This created a certain amount of residual C, which would react with the main source of oxygen-B2O3 at approximately 1600 °C which had been confirmed in our previous study [31]. Then, the oxygen contents in TBS0 and TBS10 would decrease. In TBS20 and TBS30, C was consumed by excess B, and B2O3 could be removed and was detrimental for the densification [23]. As the B content increased, the volume fraction of BxC in the composites increased, so the hardness also increased, as shown in Table 3. The bending strength of TBS0 is only 587 MPa, and it was considered that the residual C has a serious effect on the strength. As the B content increased, the volume fraction of BxC increased, causing the grain size of TiB2 to decrease and the dispersion uniformity to improve. According to Griffith's microcrack theory, the size of an inherent crack in a material depends on the grain size to a great extent. In addition, the thermal expansion coefficient of TiB2 is much higher than those of B4C and SiC, so in the BxC-TiB2-SiC system, microcracks are mainly generated at the grain boundary between TiB2 and B4C or SiC. Therefore, a smaller TiB2 grain size leads to a smaller initial crack size and, correspondingly, a higher bending strength of the material [32]. TBS0 had the highest fracture toughness of 6.63 MPa m1/2 because it had the largest grain size and volume fraction of TiB2, which makes crack deflection easier and more frequent. Fig. 8(a) shows the crack growth path of TBS0 and Fig. 8(b) shows the crack growth path of TBS30. It is obvious that numerous crack deflections occurred during the crack propagation process in TBS0. However, crack deflection only occurs in the region where TiB2 grains are concentrated in TBS30. The frequency of crack deflection varies with the TiB2 content, so the toughness of TBS0 (6.63 MPa m1/2) is higher than that of TBS30 (5.22 MPa m1/2). 3.3. Analysis of microstructure Fig. 9(a) shows the dark-field TEM image of TBS30, and Fig. 9(b)-
4. Conclusions BxC-TiB2-SiC ceramic composites were fabricated via reactive hot pressing using TiC, B, and Si as raw materials. The composite powders of B4C, TiB2, and SiC could be obtained by heating at 1400 °C for 1 h. The powder synthesised in situ is nanometre sized and has good
Fig. 8. SEM images of the crack propagations of (a) TBS0 and (b) TBS30. 6
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Fig. 9. (a) TEM image of TBS30; (b)–(d) SAED images of TiB2, SiC, and BxC; (e)–(g) HRTEM images of BxC, BxC-TiB2, and TiB2-SiC.
Fig. 10. (a), (d) TEM image of intragranular phase in TiB2 of TBS30; (b), (e) C element distribution; and (c), (f) B element distribution.
Talent of Higher Learning Institutions of Hebei Province (No. BJ2019054).
sintering activity. In addition, excess B can form boron-rich boron carbide, and the grain contains a large number of twins, which is beneficial to the mechanical properties of the composite. Finally, the TBS30 composite had excellent comprehensive mechanical properties, with a relative density, hardness, bending strength, and fracture toughness of 98.32%, 33.2 GPa, 840 MPa, and 5.22 MPa m1/2, respectively.
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Acknowledgment The authors acknowledge the financial support from the National Key Research and Development Plan of China (2017YFB0310400), the National Natural Science Foundation of China (5167020705), the State Key Laboratory of Advanced Technology for Materials Synthesis and Processing (No. 2019-KF-15), and the Program for the Youth Top-notch 7
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