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High-performance B4C–TiB2–SiC composites with tuneable properties fabricated by reactive hot pressing
Journal of the European Ceramic Society 39 (2019) 2995–3002
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Original article
High-performance B4C–TiB2–SiC composites with tuneable properties fabricated by reactive hot pressing
T
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Xiaorong Zhanga, Zhixiao Zhangb,d, , Yuming Liub, Aiyang Wangc, Shi Tianc, Weimin Wangc, John Wangd a
College of Mechanical and Equipment Engineering, Hebei University of Engineering, Handan, 056038, China College of Materials Science and Engineering, Hebei University of Engineering, Handan, 056038, China c State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan, 430070, China d Department of Materials Science and Engineering, National University of Singapore, 117574, Singapore b
A R T I C LE I N FO
A B S T R A C T
Keywords: B4C-TiB2-SiC composites Reactive hot pressing Toughening and strengthening Microstructure Mechanical properties
In this work, we systematically studied the effects of powder characteristics (B4C, TiC and Si powders) on the existential form of toughening phases (SiC and TiB2) as well as the overall microstructure and properties of B4C–TiB2–SiC composites fabricated by reactive hot pressing. The particle size of the TiC powder plays a largely determining role in the development of novel toughening phases, the TiB2–SiC composite structure, that are formed in the B4C matrix, while the Si particle size affects the agglomerate level of the SiC phase. The TiB2–SiC composite structure and SiC agglomerates enhance the fracture toughness, but decrease the flexural strength. Both the microstructure and mechanical properties of B4C–TiB2–SiC composites can be effectively tuned by regulating the combinations of the particle sizes of the starting powders. The B4C–TiB2–SiC composites demonstrate flexural strength, fracture toughness and Vickers hardness in the respective range of 567–632 MPa, 5.11–6.38 MPa m1/2, and 34.8–35.6 GPa.
1. Introduction Boron carbide (B4C), as a useful engineering material, has attracted considerable attention for applications in the fields of armors, structure and aerospace engineering, owing to its excellent thermal, chemical and mechanical properties, including high temperature stability, chemical inertness, high corrosion resistance and the combination of ultrahigh hardness and low density, which make it oustanding compared with other structural ceramics [1–4]. However, their broad range of applications are limited by the poor sinterability and intrinsically low fracture toughness, and therefore poor mechanical properties, when they are fabricated by the conventional ceramic process [5–8]. Considerable attempts have thus been made with the novel fabrication and toughening by developing appropriate composite structures. Indeed, one of the most common and feasible approaches is the introduction of various second-phases, such as TiB2, HfB2, ZrB2, SiC and Al2O3, into the B4C matrix to develop a composite-type ceramic, together with enhanced sintering processes, for example, by hot pressing and spark plasma sintering [1,9–14]. As shown in Fig. 1(a), the microstructure of B4C-based composites is generally characterised by the presence of single second phase dispersed in the ceramic matrix. Accordingly,
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several studies have been made on the toughening of ceramic matrix composites, by manipulating the type, morphology and distribution of the second-phase. The toughening effects are also affected by the difference in thermal expansion coefficients between the B4C matrix and second-phase, which can induce strain energy and cracks. Although a few studies are available on the three-phase B4C composites toughened by two types of second-phases [10,15,16], where the microstructure of the three-phased B4C composites is characterized by the two types of second-phases that are segregated from each other and independently dispersed in the B4C matrix, as shown in the diagram of Fig. 1(b). The dispersion state of the isolated and independent second-phases remains unchanged. Therefore, the toughening structure shown in Fig. 1(b) is almost the same as that shown in Fig. 1(a), thereby the over toughening effects are limited. The fracture toughness of B4C-based composites are typically in the range between 3 and 6 MPa.m1/2 [1,6,12,17–19]. To further improve the fracture toughness is rather difficult within the frame of isolated and independent second-phases. It would be thus of considerable interest to innovate and optimize a new toughening structure, in order to further improve the fracture toughness. In one of our previous studies [20], on the basis of reaction (1) below, we had successfully fabricated B4C–TiB2–SiC composites by
Corresponding author at: College of Materials Science and Engineering, Hebei University of Engineering, Handan, 056038, China. E-mail address: [email protected] (Z. Zhang).
https://doi.org/10.1016/j.jeurceramsoc.2019.04.001 Received 15 February 2019; Received in revised form 29 March 2019; Accepted 1 April 2019 Available online 02 April 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.
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Fig. 1. Schematic diagram of the microstructure of the second-phase segregation and independent distribution inside a ceramic composite matrix. (a) Two- and (b) three-phased ceramic composites.
different particle sizes were used as the starting materials. The different powders were labelled by the respective chemical formula and followed by a sequence number (1, 2, or 3), corresponding to the particle size from small to large. Table 1 lists the powder characteristics.
reactive hot pressing, where the starting materials were B4C, TiC and Si powders. (1+x)B4C+2xTiC+3xSi=B4C+2xTiB2+ 3xSiC (0 < x)
(1)
In addition, we had detailed a study on the toughening effect of a composite-type phase, namely TiB2–SiC agglomerates, with sizes of 10–40 μm, which is dispersed in B4C matrix. Fig. 2 shows the TiB2–SiC agglomerates as a second-phase introduced into the B4C matrix. The composite unit of TiB2–SiC agglomerates as a toughening phase is different compared with the common ceramic matrix composites, where the secondary phases are isolated from each other and independently dispersed [20]. In this work, we have conducted a further in-depth study on B4C–TiB2–SiC composite that were fabricated by reactive hot pressing, by looking into the effects of powder characteristics of TiC and Si on the existential forms of TiB2 and SiC (both individual and TiB2–SiC agglomerates) and on the microstructure and mechanical properties of B4C–TiB2–SiC composites thus developed. The respective hardening, strengthening, and toughening mechanisms of the resulting ceramic composites are discussed, in exploring the relationships between microstructure and mechanical properties.
2.2. Fabrications According to reaction (1), the mass loading of TiB2 and SiC (TiB2+SiC) in the B4C composite can be adjusted by regulating the amount of B4C powder added in the starting materials. The mass ratio of TiB2 and SiC is kept at a fixed value of 7:6. Our previous study [20] had shown that the B4C-TiB2-SiC composite containing 20 wt. % (TiB2+SiC) possessed desirable properties (where x = 0.054 in the Reaction (1), corresponding volume fractions of B4C, TiB2, and SiC are 85.77%, 6.45%, and 7.78%, respectively). Therefore, in this work, we have fixed up the mass loading of TiB2+SiC in the B4C composite at 20%. Six different B4C–TiB2–SiC ceramic composites were fabricated using the powder mixtures with different particle sizes. The ceramic composites thus made are referred to as BTαSβ (α represents the sequence number of TiC powder, and β represents the sequence number of Si powder). For example, BT1S1 denotes that the ceramic composite derived from the powder mixture of B4C, TiC-1 and Si-1 powders. Table 2 lists the details of the B4C–TiB2–SiC ceramic composites, the abbreviations and their corresponding starting materials. Each of the powder mixtures was prepared by wet ball milling in ethanol medium for 24 h with SiC balls (Φ 3–10 mm) and polytef vials (500 mL). The wet ball milling was conducted in a planetary ball milling machine at a low rotation speed of 150 rpm. The ball-powder mass ratio was fixed to be 8:1. The resulting slurries were dried and then sieved with a 200-mesh sieve. The powder mixtures thus obtained were sintered by hot pressing at the conditions of 60 MPa, 1950 °C, 60 min dwell time in argon atmosphere. The detailed preparation procedures were described in Ref. [20].
2. Material and methods 2.1. Materials One B4C powder and three types of TiC and Si powders with
2.3. Characterizations The phase components in the ceramic composites were studied using the X-ray diffraction (XRD, X’Pert PRO-PANalytical) with CuKα radiation (40 kV, 40 mA). Their microstructures were characterized using the scanning electron microscopy (SEM, Hitachi S4800), and high-resolution transmission electron microscopy (HRTEM, JEOL JEM2100 F). The density of each ceramic composite was measured using the Archimedes method. Theoretical density of the composite was calculated according to the following equation: Fig. 2. Schematic diagram of the microstructure of the ceramic composites containing regions of composite toughening phases.
ρ=
2996
mtotal mB 4C ρB 4C
+
mTiB2 ρTiB2
+
mSiC ρSiC
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Table 1 Characteristics of starting powders used in the present work. Starting powders
Purity [%]
Mean particle size [μm]
Dominating particle size distribution [μm]
Supplier
B4C TiC-1 TiC-2 TiC-3 Si-1 Si-2 Si-3
97.5 99.0 99.0 99.0 99.0 99.0 99.9
3.00 0.05 0.60 3.00 0.15 2.00 20.00
1.00-5.00 0.02-0.08 0.05-2.00 0.10-5.00 0.10-1.00 0.10-4.00 1.00-30.00
Jingangzuan Boron Carbide Co., Ltd., China Aladdin Co., Ltd., China Shanghai Chaowei Nanotechnology Co., Ltd., China Aladdin Co., Ltd., China Shanghai Chaowei Nanotechnology Co., Ltd., China Shanghai Chaowei Nanotechnology Co., Ltd., China Sinopharm Chemical Reagent Co., Ltd., China
Table 2 B4C–TiB2–SiC composite ceramic abbreviations and their corresponding raw materials. B4C-TiB2-SiC composite ceramics
BT1S1 BT1S2 BT1S3 BT2S1 BT3S1 BT3S3
Raw powders B4C (84.31 wt. %)
TiC (9.23 wt. %)
Si (6.46 wt.%)
B4C
TiC-1 TiC-1 TiC-1 TiC-2 TiC-3 TiC-3
Si-1 Si-2 Si-3 Si-1 Si-1 Si-3
The mass percent of the starting powders was calculated according to the reaction (1) (x = 0.054).
where ρ is the theoretical density of the ceramic composite. mtotal is the total mass of the composite. mB4C, mTiB2, and mSiC are the mass of B4C, TiB2, and SiC in the composites, respectively. ρB4C, ρTiB2, and ρSiC are the theoretical density of B4C (2.52 g/cm3), TiB2 (4.52 g/cm3), and SiC (3.20 g/cm3), respectively. The relative density was the quotient of the actual density divided by the theoretical density. The average value and standard deviation of the relative density for each material were obtained based on the measurements of five samples. Three-point bending strength was tested on the specimen bars of 3 mm × 4 mm × 36 mm with the span of 30 mm. Fracture toughness was determined by the single-edge notched beam (SENB) method using the specimen bars of 3 mm × 5 mm × 25 mm with 2.5 mm notch deepth and 0.2 mm notch width. The strength and toughness values for each sample were calculated based on the measurements of five bars. Vickers hardness was determined using a hardness tester (Wolpert 430SVD) with a load of 1 × 9.8 N, applied for 15 s on the polished surface, where the average was calculated from 7 indentations.
Fig. 3. Relative density and Vickers hardness of the ceramic composites fabricated by employing the starting powders of different particle sizes.
3. Results and discussion 3.1. Mechanical properties
Fig. 4. Flexural strength and fracture toughness of the ceramic composites fabricated by employing the starting powders of different particle sizes.
Fig. 3 shows the relative density and Vickers hardness of the B4C–TiB2–SiC ceramic composites derived from different powder mixtures. Regardless of the type of powder mixture used as the starting material, the relative densities of the ceramic composites are similar, ranging from 99.35%–99.62% theoretical, which is nearly at the full density. The high sintered density is due to the high sintering pressure (60 MPa) and the addition of TiC and Si powders in the starting materials, which can effectively promote the sintering of B4C, because of the “in-situ” reactions during the sintering process [21]. The hardness of the ceramic composites shows a slight fluctuation ranging from 34.87 GPa to 35.66 GPa, where the effects of the starting powder sizes on the relative density and hardness are not obvious. Fig. 4 shows the remarkable effects of the particle sizes of the starting materials on the flexural strength and fracture toughness of B4C–TiB2–SiC composites. In accordance with the change in particle size, there is a change trend for strength and toughness. With TiC-1
powder (BT1Sβ, β = 1,2,3), the fracture toughness of the composites increases from 5.11 MPa.m1/2 to 5.63 MPa.m1/2 with the increase in Si particle size, but the flexural strength decreases from 632 MPa to 594 MPa. Similarly, when BTαS1 (α = 1,2,3) is fabricated with the same Si-1 powder as the starting powder, with the increase in particle size of the TiC powder, the fracture toughness of the composites increases from 5.11 MPa.m1/2 to 6.15 MPa.m1/2, but the flexural strength decreases from 632 MPa to 581 MPa. BT1S1 are derived using the fine TiC-1 and Si-1 powders simultaneously, and the highest strength (632 MPa) is obtained, but with the smallest toughness (5.11 MPa.m1/ 2 ). By contrast, BT3S3 is derived using relatively coarse TiC-3 and Si-3 powders simultaneously, and the smallest strength (567 MPa) is obtained, but with the largest toughness (6.38 MPa.m1/2) being observed. For ceramic composites, the microstructure is a crucial factor that 2997
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average size of less than 2 μm being uniformly dispersed in the B4C matrix, when the small-sized TiC-1 is used as the starting powder regardless of Si-1 or Si-3 being selected. However, some large SiC regions with sizes of 5–10 μm are observed (black squares). In addition, the size and number of SiC phase agglomerates in BT1S3 are more than those in BT1S1, thereby indicating that the SiC agglomerate level increases remarkably with the increase in Si particle size. By comparing Fig. 6(a) and (c), when the TiC-3 powder of large particles is used with the same Si-1 powder, the resulting ceramic composite contains a considerable amount of TiB2–SiC agglomerates (marked by black arrows), which are not formed in the composites fabricated using nano-sized TiC-1 powder in the starting materials. In addition, Fig. 6(c) and (d) show that many TiB2–SiC grain agglomerates exist inside the composites (black arrows) when the TiC-3 powder is used in the starting powder, regardless of the Si-1 or Si-3 powder being selected. Meanwhile, the SiC agglomeration phenomena are also observed, and the SiC agglomerate level increases significantly with the increase in the particle size of the Si powders. The morphology of TiB2–SiC agglomerates in the ceramic composites obtained using TiC-3 as the starting powder is almost identical to that of the TiB2–SiC ceramic composite obtained by the reaction hot pressing with B4C, TiC and Si as the starting materials in Ref. [22]. This similarity suggests that the formation of TiB2–SiC agglomerates is mainly caused by the local stoichiometric reaction of B4C, TiC and Si that occurs in the B4C matrix. Therefore, TiB2–SiC agglomerates can be considered tiny TiB2–SiC composite phases, which are introduced into the B4C matrix as a new type of second-phase. From the above discussion, regardless of the TiC powder used as the starting powder, SiC agglomeration phenomena are always observed, and the SiC agglomerate level increases with the increase in Si powder size. Si powder appears to dominate the existential form of the SiC phase. However, regardless of the type of Si powder used as the starting material, the TiB2–SiC agglomerates are formed in the B4C matrix when the large TiC-3 powder is employed. When the TiC-1 powder is chosen, no TiB2–SiC agglomerates can be observed, thereby indicating that the formation of TiB2–SiC agglomerates is mainly determined by the particle size of the TiC powder. SEM images of the fracture surface of BT1S3 are shown in Fig. 7(a), and the corresponding BSE images are shown in Fig. 7(b), to demonstrate the different phases. TiB2 is uniformly dispersed in the matrix without significant level of agglomeration. However, the SiC
Fig. 5. XRD patterns of the ceramic composites derived from powders of different particle sizes.
affects their mechanical properties. Therefore, the microstructure of different composites developed in the present work is examined in detail, in order to study the effects of starting particle size on the flexural strength and fracture toughness.
3.2. Microstructure Fig. 5 shows the XRD patterns of the composites derived from the different powder mixtures. They all exhibit the diffraction peaks of B4C, TiB2 and SiC, and no other phases can be detected, thus verifying that the ceramic composites are composed of B4C, TiB2 and SiC, regardless of the particle sizes of the starting powders used. Fig. 6 depicts the BSE images of the polished surface for the different B4C–TiB2–SiC ceramic composites, where the three components can be clearly distinguished. The dark grey-, light grey- and medium greycoloured regions represent the B4C matrix, TiB2, and SiC, respectively. By comparing the contrast in BSE images, the following observations can be made. Fig. 6(a) and (b) show that the product TiB2 with an
Fig. 6. BSE images of the polished surfaces of the ceramic composites derived from powders of different particle sizes. (a) BT1S1. (b) BT1S3. (c) BT3S1. (d) BT3S3. The black rectangles point out the SiC agglomerates; the black circles point out the TiB2-SiC agglomerates. 2998
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Fig. 7. SEM and corresponding BSE fracture surface images of BT1S3. (a) SEM and (b) BSE images. Rectangular regions are amplified and shown on the two sides of their original images.
the fracture surface of the composite regions of B4C and TiB2 is rather rugged, which is consistent with the observation presented in Fig. 7. Highly concentrated bumpy regions with size s in the range of around 20 μm are observed in BT3S3. To further study the bumpy regions, the rectangular Ⅱ area is amplified. One can clearly observe that the bumpy region is composed of the interspersed TiB2 and SiC grains with sizes of approximately 1–3 μm. The bumpy region is the TiB2–SiC agglomerates, which is consistent with the regions marked by black arrows presented in Fig. 6. Inside the TiB2–SiC agglomerates, intergranular fracture is formed, because of the significantly different thermal expansion coefficients between SiC and TiB2 [15,25–27]. The fracture mode of the TiB2–SiC agglomerates is similar to those of TiB2–SiC ceramic composites [27], thereby further indicating that the toughening phases are the fragments of TiB2–SiC composites, which are added to the B4C matrix as the second-phase. Based on the fracture surface discussed above, one can see that the grain size of SiC within the TiB2–SiC agglomerates and TiB2 grains (including those individual, inside the TiB2 agglomerates and TiB2–SiC agglomerates) is less than 5 μm, because of the intergranular fracture
agglomerate regions can be observed. From the magnified images of rectangles I and II, the regions of single B4C matrix, the composite regions of B4C and SiC, and the inner regions of SiC aggl*-*omerate are all rather smooth, thus indicating that an obvious transgranular fracture had occurred in the three types of regions. However, the composite regions of B4C and TiB2 are rather rugged, thus illustrating an obvious intergranular fracture. The above observed phenomena show that the introduction of TiB2 can change the fracture mode of the B4C matrix from transgranular to intergranular fracture, but the introduction of SiC (including SiC agglomerates) has little impact on the fracture mode change of the B4C matrix. These experimental results are consistent with the conclusions of the relevant references, which have shown that monolithic B4C ceramics and B4C–SiC composites show a transgranular fracture [15,16,23,24] and that B4C–TiB2 ceramic composites show an intergranular fracture [17,18]. Fig. 8 shows the SEM and BSE images of the BT3S3 fracture surface, where Fig. 8(a) indicates that the fracture surface in the B4C matrix regions, SiC agglomerate and the boundary between B4C and SiC is smooth (as shown in the amplified image of rectangular Ⅰ). However,
Fig. 8. SEM and corresponding BSE fracture surface images of BT3S3. (a) SEM and (b) BSE images. Rectangular regions are amplified and shown on the two sides of their original images. 2999
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Fig. 9. (a) TEM images of a selected region in BT3S3. (b) EDS analysis for the image (a).
that had occurred in TiB2–SiC and B4C–TiB2 composites. However, the grain size in the SiC agglomerate regions remains unclear, based on the fracture surface, because, due to the transgranular fracture mode of the SiC phase [28,29], the fracture surface of the SiC agglomerate regions is smooth, regardless of whether the agglomerate regions consist of numerous small SiC grains or an abnormally large grain. To determine the constitute forms of the SiC agglomerate and SiC grain sizes, TEM study with EDS analysis is conducted. Fig. 9(a) shows the TEM image for a region in BT3S3 with similar contrast. The region contains six grains with the sizes of around 3–5 μm. EDS analysis of this region indicates that the main component elements are Si and C (Fig. 9(b)). Therefore, this region is a part of the SiC grain agglomerate. The above analysis illustrates that the SiC agglomerate regions consist of numerous small SiC grains, and the size range of these SiC grains is approximately 3–5 μm.
composites. Site I shows that TiB2 grains can induce the crack deflection along the grain boundary, because of the mismatch in thermal expansion coefficients of B4C and TiB2 (4.5 × 10−6 K-1 and 8.1 × 10-6 K-1 [21], respectively) that leads to a large residual stress at the B4C–TiB2 interfaces [17]. The prolonged crack propagation path helps improve toughness, but the resistance on crack propagation is weak, due to the limited sizes of TiB2, thus leading to limited toughening effect. SiC grains or grain agglomerates cannot change the crack propagation direction, as the thermal expansion coefficients of B4C and SiC (4.5 × 10−6 K-1 and 4.35 × 10-6 K-1 [21], respectively) are similar, as shown in sites II and IV. However, SiC can continue to consume much energy at the crack tip, because the toughness of SiC is higher than that of B4C, thereby contributing to the enhanced toughness of the B4C matrix [30]. On the basis of site III, the crack can be significantly deflected along the interior boundary of TiB2 and SiC inside the TiB2–SiC agglomerates due to the large difference in the thermal expansion coefficient between TiB2 and SiC [17]. Accordingly, TiB2–SiC agglomerates can consume a significant energy during the crack extension by extending the crack growth path and changing the crack propagation direction. Among these second-phases, TiB2–SiC agglomerates therefore play a dominating role in improving the fracture toughness of B4C–TiB2–SiC ceramic composites, which has also been observed in our previous work [21]. BT1Sβ (β = 1,2,3) employed the nano-sized TiC-1 as the starting powder. SiC grain agglomerates are formed inside the composites, and the SiC agglomerate level increases with the increase in Si powder size. The presence of large SiC regions reduces the flexural strength, because of the undesirable stress effect around the large regions [12] but enhance the fracture toughness, because of the high toughness of SiC compared with that of B4C [30]. Therefore, the flexural strength of BT1Sβ decreases and fracture toughness increases with the increase in Si particle size in the starting materials. In this case, hardly any TiB2–SiC agglomerates, which can increase the fracture toughness of the composites, are formed. Therefore, the maximum fracture toughness (5.63 MPa.m1/2) of BT1Sβ is not high. BTαS1 (α = 1,2,3) uses Si-1 as the starting material. TiB2–SiC agglomerates are formed with the increase in TiC particle size. These agglomerates can significantly enhance fracture toughness due to their fracture modes, but decrease flexural strength due to the stress effect around the large regions [20]. Therefore, the flexural strength of BTαS1 decreases but fracture toughness increases with the increase in TiC particle size. When fine Si1 and TiC-1 powders are used at the same time in the starting materials, the SiC agglomerate level is mildest, and no TiB2–SiC agglomerates are formed inside the composites. Accordingly, BT1S1 exhibits the
3.3. Relationships between microstructure and mechanical properties Based on the above microstructure analyses, the B4C–TiB2–SiC ceramic composites shows significant differences when the particle size of the starting powders is varied. The features and change trends of mechanical properties of these composites are related to their diverse microstructure parameters. Therefore, the relationships between the microstructure and mechanical properties are discussed as follows. Our previous study [21] has shown that the different existent forms and types of TiB2 and SiC can present different toughening effects on the B4C–TiB2–SiC composites, because these different existent forms lead to diverse abilities of absorbing energy during crack propagation. Fig. 10 presents a schematic of the crack deflection inside the
Fig. 10. Schematic diagram of the crack propagation in B4C-TiB2-SiC ceramic composites containing TiB2-SiC grain agglomerates. 3000
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maximum flexural strength (632 MPa) and minimum fracture toughness (5.11 MPa.m1/2). By contrast, BT3S3 exhibits the minimum flexural strength (567 MPa) and maximum fracture toughness (638 MPa.m1/2), because of the abundant SiC agglomerates being present, and TiB2–SiC agglomerates that coexist inside the composites when coarse Si-3 and TiC-3 powders are used simultaneously. Mechanical properties of dense B4C composites are typically in the following ranges: Vickers hardness of 21–34 GPa, fracture toughness of 2.2–5.6 MPa m1/2 and flexural strength of 350–550 MPa [1,6,13,31–34]. In this work, by employing the combinations of starting powders with different particle sizes, the Vickers hardness, fracture toughness, flexural strength of B4C–TiB2–SiC composites have been tuned in the range from 34.8 GPa to 35.6 GPa, 567 MPa to 632 MPa and 5.11 MPa m1/2 to 6.38 MPa m1/2, respectively. The excellent hardness values of the composites obtained in this work are attributed to the combined contributions of high density (above 99%), which is the determining parameter for high hardness, and the high hardness of the second-phases (SiC and TiB2), which does not significantly reduce the hardness of B4C matrix. The overall superior fracture toughness and flexural strength of the composites obtained in this work are attributed to two main reasons. The first one is the coexistence of the TiB2, SiC agglomerates and the TiB2-SiC agglomerates, which enhance fracture toughness. Another reason is the desired interfacial compatibility between the matrix and the second-phases, which have benefitted from the in-situ reaction during the sintering process. However, the contributions from TiB2, SiC agglomerates and TiB2–SiC agglomerates on the fracture toughness and flexural strength are of significant differences. Therefore, the flexural strength and fracture toughness of the ceramic composites with different microstructures show rather large differences, but in regular change trends.
Acknowledgments This work was financially supported by the National Natural Science Foundation of China (Nos. 51702080 and 61701164), the State Key Laboratory of Advanced Technology for Materials Synthesis and Processing (Wuhan University of Technology) (No. 2019-KF-15), and the Program for the Youth Top-notch Talent of Higher Learning Institutions of Hebei Province (Nos. BJ2018039 and BJ2019054). John Wang thanks the support of the National Research Foundation Singapore (NRF-CRP17-2017-01), for research conducted at the National University of Singapore. References [1] A.K. Suri, C. Subramanian, J.K. Sonber, T.S.R.C. Murthy, Synthesis and consolidation of boron carbide: a review, Int. Mater. Rev. 55 (2010) 4–40. [2] Y. Zhu, H. Cheng, Y. Wang, R. An, Effects of carbon and silicon on microstructure and mechanical properties of pressureless sintered B4C/TiB2 composites, J. Alloys. Compd. 772 (2019) 537–545. [3] I. Bogomol, H. Borodianska, T. Zhao, T. Nishimura, Y. Sakka, P. Loboda, O. 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Zhang, Y. Ding, Microstructure and mechanical properties of B4C–TiB2 composites prepared by reaction hot pressing using Ti3SiC2 as additive, Ceram. Int. 42 (2016) 650–656. [10] A. Li, Y. Zhen, Q. Yin, L. Ma, Y. Yin, Microstructure and properties of (SiC, TiB2)/ B4C composites by reaction hot pressing, Ceram. Int. 32 (2006) 849–856. [11] A. Goldstein, Y. Geffen, A. Goldenberg, Boron carbide–zirconium boride in situ composites by the reactive pressureless sintering of boron carbide–zirconia mixtures, J. Am. Ceram. Soc. 84 (2001) 642–644. [12] B.M. Moshtaghioun, A.L. Ortiz, D. Gómez-García, A. Domínguez-Rodríguez, Toughening of super-hard ultra-fine grained B4C densified by spark-plasma sintering via SiC addition, J. Eur. Ceram. Soc. 33 (2013) 1395–1401. [13] H.W. Kim, Y.H. Koh, H.E. Kim, Densification and mechanical properties of B4C with Al2O3 as a sintering aid, J. Am. Ceram. Soc. 83 (2000) 2863–2865. [14] Z. Liu, D. Wang, J. Li, Q. Huang, S. Ran, Densification of high-strength B4C–TiB2 composites fabricated by pulsed electric current sintering of TiC–B mixture, Scripta. Mater. 135 (2017) 15–18. [15] Q. He, A. Wang, C. Liu, W. Wang, H. Wang, Z. Fu, Microstructures and mechanical properties of B4C-TiB2 -SiC composites fabricated by ball milling and hot pressing, J. Eur. Ceram. Soc. 38 (2018) 2832–2840. [16] Q. Wen, Y. Tan, Z. Zhong, H. Zhang, X. Zhou, High toughness and electrical discharge machinable B4C-TiB2-SiC composites fabricated at low sintering temperature, Mat. Sci. Eeg. A. 701 (2017) 338–343. [17] Z. Liu, X. Deng, J. Li, Y. Sun, S. Ran, Effects of B4C particle size on the microstructures and mechanical properties of hot-pressed B4C–TiB2 composites, Ceram. Int. 44 (2018) 21415–21420. [18] S. Yamada, K. Hirao, Y. Yamauchi, S. Kanzaki, High strength B4C–TiB2 composites fabricated by reaction hot-pressing, J. Eur. Ceram. Soc. 23 (2003) 1123–11130. [19] J. Sun, C. Liu, C. Duan, Effect of Al and TiO2 on sinterability and mechanical properties of boron carbide, Mat. Sci. Eeg. A. 509 (2009) 89–93. [20] X. Zhang, Z. Zhang, W. Wang, J. Shan, H. Che, J. Mu, G. Wang, Microstructure and mechanical properties of B4C-TiB2-SiC composites toughened by composite structural toughening phases, J. Am. Ceram. Soc. 100 (2017) 3099–3107. [21] X. Zhang, Z. Zhang, W. Wang, X. Zhang, J. Mu, G. Wang, Z. Fu, Preparation of B4C composites toughened by TiB2-SiC agglomerates, J. Eur. Ceram. Soc. 37 (2017) 865–869. [22] X. Zhang, Z. Zhang, Y. Liu, Y. Wang, X. Fu, Y. Bai, G. Wang, J. Mu, X. Zhang, H. Che, Effects of particle size of raw materials on the characteristics of TiB2–SiC composites fabricated from B4C, TiC and Si powders, Ceram. Int. 45 (2019) 978–984. [23] K. Sairam, J.K. Sonber, T.S.R.C. Murthy, C. Subramanian, R.K. Fotedar, P. Nanekar, R.C. Hubli, Influence of spark plasma sintering parameters on densification and mechanical properties of boron carbide, Int. J. Refract. Met. Hard Mater. 42 (2014) 185–192. [24] X.L. Shi, F.M. Xu, Y. Tan, L. Wang, Mechanical properties of hot-pressed B4C-SiC composites, in: N.P. Bansal, J.P. Singh, J. Lamon, S.R. Choi, M.M. Mahmoud (Eds.), Processing and Properties of Advanced Ceramics and Composites II, John Wiley & Sons, Inc., Hoboken, 2010, pp. 189–195.
4. Conclusions The particle sizes of TiC and Si starting powders have a large effect on the microstructure and mechanical properties of B4C–TiB2–SiC ceramic composites fabricated by reaction hot pressure. The TiC particle size can largely determine whether the TiB2–SiC composite regions (TiB2–SiC agglomerates) can be formed inside the B4C composite matrix. The Si particle size affect the agglomerate level of SiC phases inside the ceramic composites. The TiB2–SiC agglomerates and SiC agglomerates can enhance the fracture toughness, but decrease the flexural strength of B4C composites, and the presence of the TiB2–SiC agglomerates are crucial for improving the fracture toughness. With an increase in Si particle size, the SiC agglomerate level increases, thereby leading to a decrease in flexural strength of the ceramic composites but the increase in fracture toughness. However, with the increase in TiC particle size, the TiB2–SiC agglomerates are formed, thus resulting in a decrease in the flexural strength of the composites but a significant increase in fracture toughness. When fine Si and TiC powders are used at the same time in the starting materials, the SiC agglomerate level is low, and no TiB2–SiC agglomerates are formed. Therefore, the ceramic composite thus obtained demonstrates a maximum flexural strength and minimum fracture toughness. By contrast, when coarse Si and TiC powders are used simultaneously used in the starting materials, considerable level of SiC grain agglomerates and TiB2–SiC agglomerates coexist. Therefore, the ceramic composites thus developed exhibit a minimum flexural strength and maximum fracture toughness. The Vickers hardness, fracture toughness, flexural strength of the B4C–TiB2–SiC ceramic composites are in the range from 34.8 GPa to 35.6 GPa, 567 MPa to 632 MPa and 5.11 MPa m1/2 to 6.38 MPa m1/2, respectively, which are compared favourably with those reported previously. Both the microstructure and mechanical properties of B4C–TiB2–SiC ceramic composites can be effectively tuned by regulating the combinations of the particle sizes of the starting powders.
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[25] D.S. King, W.G. Fahrenholtz, G.E. Hilmas, Silicon carbide–titanium diboride ceramic composites, J. Eur. Ceram. Soc. 33 (2013) 2943–2951. [26] D.S. King, W.G. Fahrenholtz, G.E. Hilmas, Microstructural effects on the mechanical properties of SiC-15 vol%TiB2 particulate-reinforced ceramic composites, J. Am. Ceram. Soc. 96 (2013) 577–583. [27] Z. Zhang, C. Xu, X. Du, Z. Li, J. Wang, W. Xing, Y. Sheng, W. Wang, Z. Fu, Synthesis mechanism and mechanical properties of TiB2–SiC composites fabricated with the B4C–TiC–Si system by reactive hot pressing, J. Alloys Compd. 619 (2015) 26–30. [28] Z.H. Zhang, F.C. Wang, J. Luo, S.K. Lee, L. Wang, Processing and characterization of fine-grained monolithic SiC ceramic synthesized by spark plasma sintering, Mat. Sci. Eeg. A. 527 (2010) 2099–2103. [29] D. Bucevac, S. Boskovic, B. Matovic, V. Krstic, Toughening of SiC matrix with in-situ created TiB2 particles, Ceram. Int. 36 (2010) 2181–2188. [30] Z. Zhang, X. Du, Z. Li, W. Wang, J. Zhang, Z. Fu, Microstructures and mechanical
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properties of B4C–SiC intergranular/intragranular nanocomposite ceramics fabricated from B4C, Si, and graphite powders, J. Eur. Ceram. Soc. 34 (2014) 2153–2161. J. Sun, C. Liu, R. Wang, Low pressure hot pressing of B4C matrix ceramic composites improved by Al2O3 and TiC additives, Mat. Sci. Eeg. A. 519 (2009) 27–31. S.S. Rehman, W. Ji, Z. Fu, W. Wang, H. Wang, M. Asif, J. Zhang, In situ synthesis and sintering of B4C/ZrB2 composites from B4C and ZrH2 mixtures by spark plasma sintering, J. Eur. Ceram. Soc. 35 (2015) 1139–1145. M. Saeedi Heydari, H.R. Baharvandi, K. Dolatkhah, Effect of TiO2 nanoparticles on the pressureless sintering of B4C–TiB2 nanocomposites, Int. J. Refract. Met. Hard Mater. 51 (2015) 6–13. A. Moradkhani, H. Baharvandi, M.M. Mohammadi Samani, Mechanical properties and microstructure of B4C–NanoTiB2–Fe/Ni composites under different sintering temperatures, Mat. Sci. Eeg. A. 665 (2016) 141–153.
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Journal of the European Ceramic Society journal homepage: www.elsevier.com/locate/jeurceramsoc
Original article
High-performance B4C–TiB2–SiC composites with tuneable properties fabricated by reactive hot pressing
T
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Xiaorong Zhanga, Zhixiao Zhangb,d, , Yuming Liub, Aiyang Wangc, Shi Tianc, Weimin Wangc, John Wangd a
College of Mechanical and Equipment Engineering, Hebei University of Engineering, Handan, 056038, China College of Materials Science and Engineering, Hebei University of Engineering, Handan, 056038, China c State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan, 430070, China d Department of Materials Science and Engineering, National University of Singapore, 117574, Singapore b
A R T I C LE I N FO
A B S T R A C T
Keywords: B4C-TiB2-SiC composites Reactive hot pressing Toughening and strengthening Microstructure Mechanical properties
In this work, we systematically studied the effects of powder characteristics (B4C, TiC and Si powders) on the existential form of toughening phases (SiC and TiB2) as well as the overall microstructure and properties of B4C–TiB2–SiC composites fabricated by reactive hot pressing. The particle size of the TiC powder plays a largely determining role in the development of novel toughening phases, the TiB2–SiC composite structure, that are formed in the B4C matrix, while the Si particle size affects the agglomerate level of the SiC phase. The TiB2–SiC composite structure and SiC agglomerates enhance the fracture toughness, but decrease the flexural strength. Both the microstructure and mechanical properties of B4C–TiB2–SiC composites can be effectively tuned by regulating the combinations of the particle sizes of the starting powders. The B4C–TiB2–SiC composites demonstrate flexural strength, fracture toughness and Vickers hardness in the respective range of 567–632 MPa, 5.11–6.38 MPa m1/2, and 34.8–35.6 GPa.
1. Introduction Boron carbide (B4C), as a useful engineering material, has attracted considerable attention for applications in the fields of armors, structure and aerospace engineering, owing to its excellent thermal, chemical and mechanical properties, including high temperature stability, chemical inertness, high corrosion resistance and the combination of ultrahigh hardness and low density, which make it oustanding compared with other structural ceramics [1–4]. However, their broad range of applications are limited by the poor sinterability and intrinsically low fracture toughness, and therefore poor mechanical properties, when they are fabricated by the conventional ceramic process [5–8]. Considerable attempts have thus been made with the novel fabrication and toughening by developing appropriate composite structures. Indeed, one of the most common and feasible approaches is the introduction of various second-phases, such as TiB2, HfB2, ZrB2, SiC and Al2O3, into the B4C matrix to develop a composite-type ceramic, together with enhanced sintering processes, for example, by hot pressing and spark plasma sintering [1,9–14]. As shown in Fig. 1(a), the microstructure of B4C-based composites is generally characterised by the presence of single second phase dispersed in the ceramic matrix. Accordingly,
⁎
several studies have been made on the toughening of ceramic matrix composites, by manipulating the type, morphology and distribution of the second-phase. The toughening effects are also affected by the difference in thermal expansion coefficients between the B4C matrix and second-phase, which can induce strain energy and cracks. Although a few studies are available on the three-phase B4C composites toughened by two types of second-phases [10,15,16], where the microstructure of the three-phased B4C composites is characterized by the two types of second-phases that are segregated from each other and independently dispersed in the B4C matrix, as shown in the diagram of Fig. 1(b). The dispersion state of the isolated and independent second-phases remains unchanged. Therefore, the toughening structure shown in Fig. 1(b) is almost the same as that shown in Fig. 1(a), thereby the over toughening effects are limited. The fracture toughness of B4C-based composites are typically in the range between 3 and 6 MPa.m1/2 [1,6,12,17–19]. To further improve the fracture toughness is rather difficult within the frame of isolated and independent second-phases. It would be thus of considerable interest to innovate and optimize a new toughening structure, in order to further improve the fracture toughness. In one of our previous studies [20], on the basis of reaction (1) below, we had successfully fabricated B4C–TiB2–SiC composites by
Corresponding author at: College of Materials Science and Engineering, Hebei University of Engineering, Handan, 056038, China. E-mail address: [email protected] (Z. Zhang).
https://doi.org/10.1016/j.jeurceramsoc.2019.04.001 Received 15 February 2019; Received in revised form 29 March 2019; Accepted 1 April 2019 Available online 02 April 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.
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Fig. 1. Schematic diagram of the microstructure of the second-phase segregation and independent distribution inside a ceramic composite matrix. (a) Two- and (b) three-phased ceramic composites.
different particle sizes were used as the starting materials. The different powders were labelled by the respective chemical formula and followed by a sequence number (1, 2, or 3), corresponding to the particle size from small to large. Table 1 lists the powder characteristics.
reactive hot pressing, where the starting materials were B4C, TiC and Si powders. (1+x)B4C+2xTiC+3xSi=B4C+2xTiB2+ 3xSiC (0 < x)
(1)
In addition, we had detailed a study on the toughening effect of a composite-type phase, namely TiB2–SiC agglomerates, with sizes of 10–40 μm, which is dispersed in B4C matrix. Fig. 2 shows the TiB2–SiC agglomerates as a second-phase introduced into the B4C matrix. The composite unit of TiB2–SiC agglomerates as a toughening phase is different compared with the common ceramic matrix composites, where the secondary phases are isolated from each other and independently dispersed [20]. In this work, we have conducted a further in-depth study on B4C–TiB2–SiC composite that were fabricated by reactive hot pressing, by looking into the effects of powder characteristics of TiC and Si on the existential forms of TiB2 and SiC (both individual and TiB2–SiC agglomerates) and on the microstructure and mechanical properties of B4C–TiB2–SiC composites thus developed. The respective hardening, strengthening, and toughening mechanisms of the resulting ceramic composites are discussed, in exploring the relationships between microstructure and mechanical properties.
2.2. Fabrications According to reaction (1), the mass loading of TiB2 and SiC (TiB2+SiC) in the B4C composite can be adjusted by regulating the amount of B4C powder added in the starting materials. The mass ratio of TiB2 and SiC is kept at a fixed value of 7:6. Our previous study [20] had shown that the B4C-TiB2-SiC composite containing 20 wt. % (TiB2+SiC) possessed desirable properties (where x = 0.054 in the Reaction (1), corresponding volume fractions of B4C, TiB2, and SiC are 85.77%, 6.45%, and 7.78%, respectively). Therefore, in this work, we have fixed up the mass loading of TiB2+SiC in the B4C composite at 20%. Six different B4C–TiB2–SiC ceramic composites were fabricated using the powder mixtures with different particle sizes. The ceramic composites thus made are referred to as BTαSβ (α represents the sequence number of TiC powder, and β represents the sequence number of Si powder). For example, BT1S1 denotes that the ceramic composite derived from the powder mixture of B4C, TiC-1 and Si-1 powders. Table 2 lists the details of the B4C–TiB2–SiC ceramic composites, the abbreviations and their corresponding starting materials. Each of the powder mixtures was prepared by wet ball milling in ethanol medium for 24 h with SiC balls (Φ 3–10 mm) and polytef vials (500 mL). The wet ball milling was conducted in a planetary ball milling machine at a low rotation speed of 150 rpm. The ball-powder mass ratio was fixed to be 8:1. The resulting slurries were dried and then sieved with a 200-mesh sieve. The powder mixtures thus obtained were sintered by hot pressing at the conditions of 60 MPa, 1950 °C, 60 min dwell time in argon atmosphere. The detailed preparation procedures were described in Ref. [20].
2. Material and methods 2.1. Materials One B4C powder and three types of TiC and Si powders with
2.3. Characterizations The phase components in the ceramic composites were studied using the X-ray diffraction (XRD, X’Pert PRO-PANalytical) with CuKα radiation (40 kV, 40 mA). Their microstructures were characterized using the scanning electron microscopy (SEM, Hitachi S4800), and high-resolution transmission electron microscopy (HRTEM, JEOL JEM2100 F). The density of each ceramic composite was measured using the Archimedes method. Theoretical density of the composite was calculated according to the following equation: Fig. 2. Schematic diagram of the microstructure of the ceramic composites containing regions of composite toughening phases.
ρ=
2996
mtotal mB 4C ρB 4C
+
mTiB2 ρTiB2
+
mSiC ρSiC
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Table 1 Characteristics of starting powders used in the present work. Starting powders
Purity [%]
Mean particle size [μm]
Dominating particle size distribution [μm]
Supplier
B4C TiC-1 TiC-2 TiC-3 Si-1 Si-2 Si-3
97.5 99.0 99.0 99.0 99.0 99.0 99.9
3.00 0.05 0.60 3.00 0.15 2.00 20.00
1.00-5.00 0.02-0.08 0.05-2.00 0.10-5.00 0.10-1.00 0.10-4.00 1.00-30.00
Jingangzuan Boron Carbide Co., Ltd., China Aladdin Co., Ltd., China Shanghai Chaowei Nanotechnology Co., Ltd., China Aladdin Co., Ltd., China Shanghai Chaowei Nanotechnology Co., Ltd., China Shanghai Chaowei Nanotechnology Co., Ltd., China Sinopharm Chemical Reagent Co., Ltd., China
Table 2 B4C–TiB2–SiC composite ceramic abbreviations and their corresponding raw materials. B4C-TiB2-SiC composite ceramics
BT1S1 BT1S2 BT1S3 BT2S1 BT3S1 BT3S3
Raw powders B4C (84.31 wt. %)
TiC (9.23 wt. %)
Si (6.46 wt.%)
B4C
TiC-1 TiC-1 TiC-1 TiC-2 TiC-3 TiC-3
Si-1 Si-2 Si-3 Si-1 Si-1 Si-3
The mass percent of the starting powders was calculated according to the reaction (1) (x = 0.054).
where ρ is the theoretical density of the ceramic composite. mtotal is the total mass of the composite. mB4C, mTiB2, and mSiC are the mass of B4C, TiB2, and SiC in the composites, respectively. ρB4C, ρTiB2, and ρSiC are the theoretical density of B4C (2.52 g/cm3), TiB2 (4.52 g/cm3), and SiC (3.20 g/cm3), respectively. The relative density was the quotient of the actual density divided by the theoretical density. The average value and standard deviation of the relative density for each material were obtained based on the measurements of five samples. Three-point bending strength was tested on the specimen bars of 3 mm × 4 mm × 36 mm with the span of 30 mm. Fracture toughness was determined by the single-edge notched beam (SENB) method using the specimen bars of 3 mm × 5 mm × 25 mm with 2.5 mm notch deepth and 0.2 mm notch width. The strength and toughness values for each sample were calculated based on the measurements of five bars. Vickers hardness was determined using a hardness tester (Wolpert 430SVD) with a load of 1 × 9.8 N, applied for 15 s on the polished surface, where the average was calculated from 7 indentations.
Fig. 3. Relative density and Vickers hardness of the ceramic composites fabricated by employing the starting powders of different particle sizes.
3. Results and discussion 3.1. Mechanical properties
Fig. 4. Flexural strength and fracture toughness of the ceramic composites fabricated by employing the starting powders of different particle sizes.
Fig. 3 shows the relative density and Vickers hardness of the B4C–TiB2–SiC ceramic composites derived from different powder mixtures. Regardless of the type of powder mixture used as the starting material, the relative densities of the ceramic composites are similar, ranging from 99.35%–99.62% theoretical, which is nearly at the full density. The high sintered density is due to the high sintering pressure (60 MPa) and the addition of TiC and Si powders in the starting materials, which can effectively promote the sintering of B4C, because of the “in-situ” reactions during the sintering process [21]. The hardness of the ceramic composites shows a slight fluctuation ranging from 34.87 GPa to 35.66 GPa, where the effects of the starting powder sizes on the relative density and hardness are not obvious. Fig. 4 shows the remarkable effects of the particle sizes of the starting materials on the flexural strength and fracture toughness of B4C–TiB2–SiC composites. In accordance with the change in particle size, there is a change trend for strength and toughness. With TiC-1
powder (BT1Sβ, β = 1,2,3), the fracture toughness of the composites increases from 5.11 MPa.m1/2 to 5.63 MPa.m1/2 with the increase in Si particle size, but the flexural strength decreases from 632 MPa to 594 MPa. Similarly, when BTαS1 (α = 1,2,3) is fabricated with the same Si-1 powder as the starting powder, with the increase in particle size of the TiC powder, the fracture toughness of the composites increases from 5.11 MPa.m1/2 to 6.15 MPa.m1/2, but the flexural strength decreases from 632 MPa to 581 MPa. BT1S1 are derived using the fine TiC-1 and Si-1 powders simultaneously, and the highest strength (632 MPa) is obtained, but with the smallest toughness (5.11 MPa.m1/ 2 ). By contrast, BT3S3 is derived using relatively coarse TiC-3 and Si-3 powders simultaneously, and the smallest strength (567 MPa) is obtained, but with the largest toughness (6.38 MPa.m1/2) being observed. For ceramic composites, the microstructure is a crucial factor that 2997
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average size of less than 2 μm being uniformly dispersed in the B4C matrix, when the small-sized TiC-1 is used as the starting powder regardless of Si-1 or Si-3 being selected. However, some large SiC regions with sizes of 5–10 μm are observed (black squares). In addition, the size and number of SiC phase agglomerates in BT1S3 are more than those in BT1S1, thereby indicating that the SiC agglomerate level increases remarkably with the increase in Si particle size. By comparing Fig. 6(a) and (c), when the TiC-3 powder of large particles is used with the same Si-1 powder, the resulting ceramic composite contains a considerable amount of TiB2–SiC agglomerates (marked by black arrows), which are not formed in the composites fabricated using nano-sized TiC-1 powder in the starting materials. In addition, Fig. 6(c) and (d) show that many TiB2–SiC grain agglomerates exist inside the composites (black arrows) when the TiC-3 powder is used in the starting powder, regardless of the Si-1 or Si-3 powder being selected. Meanwhile, the SiC agglomeration phenomena are also observed, and the SiC agglomerate level increases significantly with the increase in the particle size of the Si powders. The morphology of TiB2–SiC agglomerates in the ceramic composites obtained using TiC-3 as the starting powder is almost identical to that of the TiB2–SiC ceramic composite obtained by the reaction hot pressing with B4C, TiC and Si as the starting materials in Ref. [22]. This similarity suggests that the formation of TiB2–SiC agglomerates is mainly caused by the local stoichiometric reaction of B4C, TiC and Si that occurs in the B4C matrix. Therefore, TiB2–SiC agglomerates can be considered tiny TiB2–SiC composite phases, which are introduced into the B4C matrix as a new type of second-phase. From the above discussion, regardless of the TiC powder used as the starting powder, SiC agglomeration phenomena are always observed, and the SiC agglomerate level increases with the increase in Si powder size. Si powder appears to dominate the existential form of the SiC phase. However, regardless of the type of Si powder used as the starting material, the TiB2–SiC agglomerates are formed in the B4C matrix when the large TiC-3 powder is employed. When the TiC-1 powder is chosen, no TiB2–SiC agglomerates can be observed, thereby indicating that the formation of TiB2–SiC agglomerates is mainly determined by the particle size of the TiC powder. SEM images of the fracture surface of BT1S3 are shown in Fig. 7(a), and the corresponding BSE images are shown in Fig. 7(b), to demonstrate the different phases. TiB2 is uniformly dispersed in the matrix without significant level of agglomeration. However, the SiC
Fig. 5. XRD patterns of the ceramic composites derived from powders of different particle sizes.
affects their mechanical properties. Therefore, the microstructure of different composites developed in the present work is examined in detail, in order to study the effects of starting particle size on the flexural strength and fracture toughness.
3.2. Microstructure Fig. 5 shows the XRD patterns of the composites derived from the different powder mixtures. They all exhibit the diffraction peaks of B4C, TiB2 and SiC, and no other phases can be detected, thus verifying that the ceramic composites are composed of B4C, TiB2 and SiC, regardless of the particle sizes of the starting powders used. Fig. 6 depicts the BSE images of the polished surface for the different B4C–TiB2–SiC ceramic composites, where the three components can be clearly distinguished. The dark grey-, light grey- and medium greycoloured regions represent the B4C matrix, TiB2, and SiC, respectively. By comparing the contrast in BSE images, the following observations can be made. Fig. 6(a) and (b) show that the product TiB2 with an
Fig. 6. BSE images of the polished surfaces of the ceramic composites derived from powders of different particle sizes. (a) BT1S1. (b) BT1S3. (c) BT3S1. (d) BT3S3. The black rectangles point out the SiC agglomerates; the black circles point out the TiB2-SiC agglomerates. 2998
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Fig. 7. SEM and corresponding BSE fracture surface images of BT1S3. (a) SEM and (b) BSE images. Rectangular regions are amplified and shown on the two sides of their original images.
the fracture surface of the composite regions of B4C and TiB2 is rather rugged, which is consistent with the observation presented in Fig. 7. Highly concentrated bumpy regions with size s in the range of around 20 μm are observed in BT3S3. To further study the bumpy regions, the rectangular Ⅱ area is amplified. One can clearly observe that the bumpy region is composed of the interspersed TiB2 and SiC grains with sizes of approximately 1–3 μm. The bumpy region is the TiB2–SiC agglomerates, which is consistent with the regions marked by black arrows presented in Fig. 6. Inside the TiB2–SiC agglomerates, intergranular fracture is formed, because of the significantly different thermal expansion coefficients between SiC and TiB2 [15,25–27]. The fracture mode of the TiB2–SiC agglomerates is similar to those of TiB2–SiC ceramic composites [27], thereby further indicating that the toughening phases are the fragments of TiB2–SiC composites, which are added to the B4C matrix as the second-phase. Based on the fracture surface discussed above, one can see that the grain size of SiC within the TiB2–SiC agglomerates and TiB2 grains (including those individual, inside the TiB2 agglomerates and TiB2–SiC agglomerates) is less than 5 μm, because of the intergranular fracture
agglomerate regions can be observed. From the magnified images of rectangles I and II, the regions of single B4C matrix, the composite regions of B4C and SiC, and the inner regions of SiC aggl*-*omerate are all rather smooth, thus indicating that an obvious transgranular fracture had occurred in the three types of regions. However, the composite regions of B4C and TiB2 are rather rugged, thus illustrating an obvious intergranular fracture. The above observed phenomena show that the introduction of TiB2 can change the fracture mode of the B4C matrix from transgranular to intergranular fracture, but the introduction of SiC (including SiC agglomerates) has little impact on the fracture mode change of the B4C matrix. These experimental results are consistent with the conclusions of the relevant references, which have shown that monolithic B4C ceramics and B4C–SiC composites show a transgranular fracture [15,16,23,24] and that B4C–TiB2 ceramic composites show an intergranular fracture [17,18]. Fig. 8 shows the SEM and BSE images of the BT3S3 fracture surface, where Fig. 8(a) indicates that the fracture surface in the B4C matrix regions, SiC agglomerate and the boundary between B4C and SiC is smooth (as shown in the amplified image of rectangular Ⅰ). However,
Fig. 8. SEM and corresponding BSE fracture surface images of BT3S3. (a) SEM and (b) BSE images. Rectangular regions are amplified and shown on the two sides of their original images. 2999
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Fig. 9. (a) TEM images of a selected region in BT3S3. (b) EDS analysis for the image (a).
that had occurred in TiB2–SiC and B4C–TiB2 composites. However, the grain size in the SiC agglomerate regions remains unclear, based on the fracture surface, because, due to the transgranular fracture mode of the SiC phase [28,29], the fracture surface of the SiC agglomerate regions is smooth, regardless of whether the agglomerate regions consist of numerous small SiC grains or an abnormally large grain. To determine the constitute forms of the SiC agglomerate and SiC grain sizes, TEM study with EDS analysis is conducted. Fig. 9(a) shows the TEM image for a region in BT3S3 with similar contrast. The region contains six grains with the sizes of around 3–5 μm. EDS analysis of this region indicates that the main component elements are Si and C (Fig. 9(b)). Therefore, this region is a part of the SiC grain agglomerate. The above analysis illustrates that the SiC agglomerate regions consist of numerous small SiC grains, and the size range of these SiC grains is approximately 3–5 μm.
composites. Site I shows that TiB2 grains can induce the crack deflection along the grain boundary, because of the mismatch in thermal expansion coefficients of B4C and TiB2 (4.5 × 10−6 K-1 and 8.1 × 10-6 K-1 [21], respectively) that leads to a large residual stress at the B4C–TiB2 interfaces [17]. The prolonged crack propagation path helps improve toughness, but the resistance on crack propagation is weak, due to the limited sizes of TiB2, thus leading to limited toughening effect. SiC grains or grain agglomerates cannot change the crack propagation direction, as the thermal expansion coefficients of B4C and SiC (4.5 × 10−6 K-1 and 4.35 × 10-6 K-1 [21], respectively) are similar, as shown in sites II and IV. However, SiC can continue to consume much energy at the crack tip, because the toughness of SiC is higher than that of B4C, thereby contributing to the enhanced toughness of the B4C matrix [30]. On the basis of site III, the crack can be significantly deflected along the interior boundary of TiB2 and SiC inside the TiB2–SiC agglomerates due to the large difference in the thermal expansion coefficient between TiB2 and SiC [17]. Accordingly, TiB2–SiC agglomerates can consume a significant energy during the crack extension by extending the crack growth path and changing the crack propagation direction. Among these second-phases, TiB2–SiC agglomerates therefore play a dominating role in improving the fracture toughness of B4C–TiB2–SiC ceramic composites, which has also been observed in our previous work [21]. BT1Sβ (β = 1,2,3) employed the nano-sized TiC-1 as the starting powder. SiC grain agglomerates are formed inside the composites, and the SiC agglomerate level increases with the increase in Si powder size. The presence of large SiC regions reduces the flexural strength, because of the undesirable stress effect around the large regions [12] but enhance the fracture toughness, because of the high toughness of SiC compared with that of B4C [30]. Therefore, the flexural strength of BT1Sβ decreases and fracture toughness increases with the increase in Si particle size in the starting materials. In this case, hardly any TiB2–SiC agglomerates, which can increase the fracture toughness of the composites, are formed. Therefore, the maximum fracture toughness (5.63 MPa.m1/2) of BT1Sβ is not high. BTαS1 (α = 1,2,3) uses Si-1 as the starting material. TiB2–SiC agglomerates are formed with the increase in TiC particle size. These agglomerates can significantly enhance fracture toughness due to their fracture modes, but decrease flexural strength due to the stress effect around the large regions [20]. Therefore, the flexural strength of BTαS1 decreases but fracture toughness increases with the increase in TiC particle size. When fine Si1 and TiC-1 powders are used at the same time in the starting materials, the SiC agglomerate level is mildest, and no TiB2–SiC agglomerates are formed inside the composites. Accordingly, BT1S1 exhibits the
3.3. Relationships between microstructure and mechanical properties Based on the above microstructure analyses, the B4C–TiB2–SiC ceramic composites shows significant differences when the particle size of the starting powders is varied. The features and change trends of mechanical properties of these composites are related to their diverse microstructure parameters. Therefore, the relationships between the microstructure and mechanical properties are discussed as follows. Our previous study [21] has shown that the different existent forms and types of TiB2 and SiC can present different toughening effects on the B4C–TiB2–SiC composites, because these different existent forms lead to diverse abilities of absorbing energy during crack propagation. Fig. 10 presents a schematic of the crack deflection inside the
Fig. 10. Schematic diagram of the crack propagation in B4C-TiB2-SiC ceramic composites containing TiB2-SiC grain agglomerates. 3000
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maximum flexural strength (632 MPa) and minimum fracture toughness (5.11 MPa.m1/2). By contrast, BT3S3 exhibits the minimum flexural strength (567 MPa) and maximum fracture toughness (638 MPa.m1/2), because of the abundant SiC agglomerates being present, and TiB2–SiC agglomerates that coexist inside the composites when coarse Si-3 and TiC-3 powders are used simultaneously. Mechanical properties of dense B4C composites are typically in the following ranges: Vickers hardness of 21–34 GPa, fracture toughness of 2.2–5.6 MPa m1/2 and flexural strength of 350–550 MPa [1,6,13,31–34]. In this work, by employing the combinations of starting powders with different particle sizes, the Vickers hardness, fracture toughness, flexural strength of B4C–TiB2–SiC composites have been tuned in the range from 34.8 GPa to 35.6 GPa, 567 MPa to 632 MPa and 5.11 MPa m1/2 to 6.38 MPa m1/2, respectively. The excellent hardness values of the composites obtained in this work are attributed to the combined contributions of high density (above 99%), which is the determining parameter for high hardness, and the high hardness of the second-phases (SiC and TiB2), which does not significantly reduce the hardness of B4C matrix. The overall superior fracture toughness and flexural strength of the composites obtained in this work are attributed to two main reasons. The first one is the coexistence of the TiB2, SiC agglomerates and the TiB2-SiC agglomerates, which enhance fracture toughness. Another reason is the desired interfacial compatibility between the matrix and the second-phases, which have benefitted from the in-situ reaction during the sintering process. However, the contributions from TiB2, SiC agglomerates and TiB2–SiC agglomerates on the fracture toughness and flexural strength are of significant differences. Therefore, the flexural strength and fracture toughness of the ceramic composites with different microstructures show rather large differences, but in regular change trends.
Acknowledgments This work was financially supported by the National Natural Science Foundation of China (Nos. 51702080 and 61701164), the State Key Laboratory of Advanced Technology for Materials Synthesis and Processing (Wuhan University of Technology) (No. 2019-KF-15), and the Program for the Youth Top-notch Talent of Higher Learning Institutions of Hebei Province (Nos. BJ2018039 and BJ2019054). John Wang thanks the support of the National Research Foundation Singapore (NRF-CRP17-2017-01), for research conducted at the National University of Singapore. References [1] A.K. Suri, C. Subramanian, J.K. Sonber, T.S.R.C. Murthy, Synthesis and consolidation of boron carbide: a review, Int. Mater. Rev. 55 (2010) 4–40. [2] Y. Zhu, H. Cheng, Y. Wang, R. An, Effects of carbon and silicon on microstructure and mechanical properties of pressureless sintered B4C/TiB2 composites, J. Alloys. Compd. 772 (2019) 537–545. [3] I. Bogomol, H. Borodianska, T. Zhao, T. Nishimura, Y. Sakka, P. Loboda, O. 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4. Conclusions The particle sizes of TiC and Si starting powders have a large effect on the microstructure and mechanical properties of B4C–TiB2–SiC ceramic composites fabricated by reaction hot pressure. The TiC particle size can largely determine whether the TiB2–SiC composite regions (TiB2–SiC agglomerates) can be formed inside the B4C composite matrix. The Si particle size affect the agglomerate level of SiC phases inside the ceramic composites. The TiB2–SiC agglomerates and SiC agglomerates can enhance the fracture toughness, but decrease the flexural strength of B4C composites, and the presence of the TiB2–SiC agglomerates are crucial for improving the fracture toughness. With an increase in Si particle size, the SiC agglomerate level increases, thereby leading to a decrease in flexural strength of the ceramic composites but the increase in fracture toughness. However, with the increase in TiC particle size, the TiB2–SiC agglomerates are formed, thus resulting in a decrease in the flexural strength of the composites but a significant increase in fracture toughness. When fine Si and TiC powders are used at the same time in the starting materials, the SiC agglomerate level is low, and no TiB2–SiC agglomerates are formed. Therefore, the ceramic composite thus obtained demonstrates a maximum flexural strength and minimum fracture toughness. By contrast, when coarse Si and TiC powders are used simultaneously used in the starting materials, considerable level of SiC grain agglomerates and TiB2–SiC agglomerates coexist. Therefore, the ceramic composites thus developed exhibit a minimum flexural strength and maximum fracture toughness. The Vickers hardness, fracture toughness, flexural strength of the B4C–TiB2–SiC ceramic composites are in the range from 34.8 GPa to 35.6 GPa, 567 MPa to 632 MPa and 5.11 MPa m1/2 to 6.38 MPa m1/2, respectively, which are compared favourably with those reported previously. Both the microstructure and mechanical properties of B4C–TiB2–SiC ceramic composites can be effectively tuned by regulating the combinations of the particle sizes of the starting powders.
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