- Email: [email protected]
Densification and mechanical properties of boron carbide prepared via spark plasma sintering with cubic boron nitride as an additive
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
Contents lists available at ScienceDirect
Journal of the European Ceramic Society journal homepage: www.elsevier.com/locate/jeurceramsoc
Original article
Densification and mechanical properties of boron carbide prepared via spark plasma sintering with cubic boron nitride as an additive Jinchang Suna, Bo Niua, Lin Renb,*, Jinyong Zhanga, Liwen Leia, Fan Zhanga,* a b
State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan, 430070, PR China School of Science, Wuhan University of Technology, Wuhan, 430070, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: B4C-based ceramic Phase transformation Spark plasma sintering Densification Mechanical properties
Hexagonal boron nitride (h-BN) can reinforce boron carbide (B4C) ceramics, but homogeneous dispersion of hBN is difficult to achieve using conventional methods. Herein, B4C/h-BN composites were manufactured via the transformation of cubic (c-) BN during spark plasma sintering at 1800 °C. The effects of the c-BN content on the microstructure, densification, and mechanical properties of B4C/h-BN composites were evaluated. In situ synthesized h-BN platelets were homogeneously dispersed in the B4C matrix and the growth of B4C grains was effectively suppressed. Moreover, the c-BN to h-BN phase transformation improved the sinterability of B4C. The sample with 5 vol.% c-BN exhibited excellent integrated mechanical properties (hardness of 30.5 GPa, bending strength of 470 MPa, and fracture toughness of 3.84 MPa⋅ m1/2). Higher c-BN contents did not significantly affect the bending strength and fracture toughness but clearly decreased the hardness. The main toughening mechanisms were crack deflection, crack bridging, and pulling out of h-BN.
1. Introduction Boron carbide (B4C) has attracted much attention because of its low density (2.52 g/cm3), high hardness (35 GPa), excellent wear stability, great chemical resistance, and high melting point (2450 °C) [1–3]. Owing to these outstanding properties, B4C ceramics have been used in many structural applications such as aerospace craft, shield materials, and abrasive and polishing media for hard and wear-resistant materials [4,5]. B4C ceramics are also currently used in neutron detectors because of their high neutron absorption cross sections [6]. However, the practical applications of B4C ceramics are severely hindered by numerous drawbacks such as low fracture toughness, low flexural strength, and poor sinterability. Three methods are commonly used to obtain highly dense B4C ceramics: (i) increasing the sintering temperature, (ii) applying a high pressure (≥80 MPa) [7], and (iii) using additives, including Ti, Al, Fe, Ni, Cu, Si, TiO2, Cr2O3, and ZrB2 [8–14]. However, a high temperature leads to excessive growth of the B4C grains, which could reduce the strenght [7], and it is difficult to apply a high pressure in standard sintering equipment. In contrast, additives are widely used in the production of B4C because they can reduce the sintering temperature to
below 2150 °C [15] or enhance the mechanical properties of the manufactured ceramics. In recent years, graphene platelets (GPLs) have received considerable attention as efficient reinforcement materials for B4C ceramics because of their outstanding properties such as being lightweight and having a two-dimensional (2D) high aspect ratio and a high tensile strength [16,17]. However, as reinforcement materials for ceramic matrix composites, the applicability of GPLs is limited, mainly owing to the difficulties in realizing heterogeneous dispersion for its large surface area. Furthermore, GPLs exhibit poor high-temperature structure stability [18] and can react with B4C, weakening the reinforcement ability of GPLs in B4C. Currently, BN platelets, a structural analogue of GPLs, are usually used as a reinforcement material for ceramics. Mettaya Kitiwan [19] used spark plasma sintering method to fabricate Titanium nitride (TiN)/Titanium diboride (TiB2)/hexagonal boron nitride (h-BN) composites at 1700 °C under 100 MPa. A fracture toughness of 4.3 MPa⋅ m1/2 was reported for composites with 15 vol.% of h-BN, which is higher than the 2.8 MPa⋅ m1/2 of the pure TiN/TiB2 composite. Using an HP process, Lee et al. [20] prepared Si3N4/h-BN with a fracture toughness of 9.76 MPa⋅ m1/2 for the composite with 2 wt.% h-BN, corresponding to a 24.7 % increase compared with the value for pure
Abbreviations: (GPL), graphene platelet; (2D), two-dimensional; (HP), hot-pressing; (h-BN), hexagonal boron nitride; (c-BN), cubic boron nitride; (SPS), spark plasma sintering; (XRD), X-ray diffraction; (SEM), scanning electron microscopy; (TEM), transmission electron microscopy; (EDS), energy-dispersive X-ray spectroscopy ⁎ Corresponding authors. E-mail addresses: [email protected] (L. Ren), [email protected] (F. Zhang). https://doi.org/10.1016/j.jeurceramsoc.2019.12.047 Received 31 May 2019; Received in revised form 18 December 2019; Accepted 23 December 2019 0955-2219/ © 2019 Published by Elsevier Ltd.
Please cite this article as: Jinchang Sun, et al., Journal of the European Ceramic Society, https://doi.org/10.1016/j.jeurceramsoc.2019.12.047
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
Germany). A high-strength graphite die with an inner diameter of 50 mm was filled with the homogenized mixed powders and then heated to 1800 °C. To evaluate the effect of c-BN on the densification of B4C ceramics, pure B4C powders and B4C/5 vol.% c-BN composites were also prepared using SPS at 1400–1800 °C. All samples were maintained at the sintering temperature for 10 min with a heating rate of 100 °C/ min. A uniaxial pressure of 50 MPa was applied throughout the sintering cycle. In addition, B4C/5 vol.% h-BN composite and pure c-BN powders were sintered at 1800 °C and 1400–1800 °C, respectively, for 10 min under a pressure of 50 MPa using the same SPS apparatus and heating rate. The Archimedes method was used to determine the densities of the B4C-based ceramics in deionized water. The crystalline phases of the powder mixtures and the as-sintered samples were identified by X-ray diffraction (XRD, Ultima III, Rigaku, Japan) with Cu Kα radiation. The microstructures of the prepared samples were characterized by scanning electron microscopy (SEM, Hitachi 3400, Japan). A thin foil of B4C/h-BN composite was prepared by mechanical thinning using a focused ion beam and analysed by transmission electron microscopy (TEM, JEOL JEM-2010HT, Japan) coupled with energy-dispersive X-ray spectroscopy (EDS). To visualize the grain boundaries clearly, the surfaces of the polished sintered samples were electrically etched in a 1 wt. % NaOH solution with a current density of 0.1 A/cm2 for 10–30 s. The grain sizes of the sintered B4C-based samples were determined from the SEM images by averaging the lengths and widths of at least 100 randomly selected grains. The mechanical properties of the bulk materials containing different quantities of in situ synthesized h-BN were analysed by measuring their hardness, fracture toughness, flexural strength, and elastic modulus. The Vickers hardness (Hv) was determined with a Vickers hardness testing machine (Wolpert-430SV, USA) using an indentation load of 1 kg for a dwell time of 15 s, and the hardness values were calculated from 10 indentations. The flexural strength, elastic modulus, and fracture toughness were measured by three-point bending tests using a ceramic test system (MTS 810, MTS, USA). The flexural strengths (KIC) of specimens with dimensions of 3 mm × 4 mm × 36 mm were measured with a testing span of 30 mm and crosshead speed of 0.5 mm/ min. Specimens with dimensions of 2.5 mm × 5 mm × 25 mm and a notch in the middle (0.15 mm thick and 2.5 mm long) were used for fracture toughness tests. The testing span was 20 mm and the crosshead speed was 0.5 mm/min. All the test samples were ground with a 15 μm diamond grinding wheel. The four edges were chamfered to a depth of 0.12 ± 0.03 mm and an angle of 45° ± 5° to eliminate stress at the edges of the samples. Average values were obtained from the results for at least six specimens. The elastic modulus (E) was characterized with a depth-sensing nano-indentation tester (NHT, CSEM, Switzerland) equipped with a Berkovich indenter, using a peak load of 2 N on the polished surface of each sample. At least five nano-indentations were performed for each sample. In each test, the indenter was driven into the sample surface at a rate of 200 mN/min and then unloaded at the same rate. The load and the indentation depth were simultaneously
Si3N4. This improvement in performance occurs because h-BN has a high 2D aspect ratio similar to that of GPLs, and the mechanical properties of h-BN are also comparable to those of GPLs. The elastic modulus of h-BN is approximately 700–900 GPa, whereas that of GPLs is approximately 1 TPa [21]. Additionally, h-BN is not able to react with B4C and can possess better high-temperature structure stability. These findings indicate that h-BN has excellent potential for reinforcing B4C ceramics. However, the homogeneous dispersion of h-BN powder into a B4C ceramic matrix is difficult using conventional ball milling. The hBN platelets can easily agglomerate in ceramic matrices, resulting in large flaws. These flaws initiate cracks and deteriorate the mechanical properties [22]. Moreover, because of their poor sinterability, h-BN platelets hinder the interdiffusion and rearrangement of ceramic matrix particles during the sintering process [23]. Zhong et al. [24] reported that this problem could be solved by synthesizing h-BN platelets viain situ chemical reaction of aluminum nitride (AlN) and boric acid (H3BO3) in the ceramic matrix, leading to a homogeneous distribution, a compact structure, and fine h-BN grains. According to a previous report [25], the hexagonal transformation of cubic boron nitrite (c-BN) occurs at 1550–1625 °C. Thus, in this study, instead of using chemical reaction routes, a new in situ route for introducing h-BN into B4C ceramics was developed via the transformation of c-BN to h-BN by spark plasma sintering (SPS). The effects of the c-BN additive on densification and the mechanical properties, such as the fracture toughness and flexural strength, of B4C matrix ceramics were evaluated. 2. Experimental The B4C/c-BN and B4C/5 vol.% h-BN composites were prepared using commercial powders: micron B4C powders (W1.5, Mudanjiang Jingangzuan Boron Carbide, China) and c-BN powders (Funik Ultrahard Materials, China), h-BN powders (Advanced Technology & Materials Co. Ltd., China). According to the manufacturers’ data, the B2O3 content in the raw B4C powder is less than 0.2 wt.%. The average grain sizes of the B4C, c-BN and h-BN powders were 1.6, 0.3 and 4.0 μm (both from suppliers’ data), respectively. Fig. 1 shows SEM images of starting materials. First, different quantities of c-BN (0, 2, 5, 7, and 10 vol.%) and 5 vol.% h-BN were added to the B4C powder. These powder mixtures were homogenized by wet ball milling in ethyl alcohol using agate balls for 24 h on low-speed roller ball mill (GMJ/B, Xianyang JinHong General Machinery Co., LTD, Xianyang, China). The resulting mud was dried in a rotary evaporator at 65 °C and then kept in a vacuum chamber at 60 °C for 24 h. Finally, the obtained powders were sieved through a 200 mesh sieve. The Si impurity contents of the original B4C powders and B4C powders dealt with agate ball measured by inductively coupled plasma-optical emission spectrometer (ICP, Optima 4300DV, PerkinElmer, American) are 0.02 wt.% and 0.04 wt.%, respectively. Bulk samples of in situ B4C/h-BN composites and pure B4C were manufactured using SPS processes (HPD 60/0 furnace, FCT Systeme,
Fig. 1. Secondary-electron SEM image of the raw materials: (a) B4C powder and (b) c-BN powder. 2
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
obtained and used for simulation of the elastic modulus (E) with the Oliver–Pharr method [26]. The R-curve behaviour of the materials was evaluated by the indention strength bending technique. Specimens with dimensions of 3 mm × 4 mm × 36 mm were indented at the centre of the polished surface with a Vickers diamond pyramid indentation under loads of 4.9, 9.8, 29.4, 49, 98, 196, and 294 N. Subsequently, the indentation strength of each specimen was tested using a three-point bending test with a span of 30 mm and crosshead speed of 0.5 mm/min. The average indentation strength was obtained from the results for five measured specimens at each indentation load. According to indentation fracture mechanics, the stress intensity factor of the Vickers indentation crack tip (K) can be expressed by the following formula [27]:
K(c) =K σ + Kr =ΨσA c1/2 + ξ(
E 1/2 -3/2 ) Pc H
(1)
where K σ is the stress intensity factor resulting from the applied stress, Kr is the residual stress intensity factor at the indentation of a crack, KR(c) is the fracture toughness curve, Ψ is a material-independent constant that depends on the crack and specimen shape, which is equal to 1.24 [28], c is the size of the crack, P is a given indentation load, ξ is a non-dimensional geometry constant equal to 0.016 [28], E is the Young’s modulus, H is the hardness, and σA is the applied stress. For a given load (P), failure occurred at an applied stress σA = σf, which also satisfied the ‘common tangent’ according to the following equations [27]:
KR(c) =K(c)
(2)
dKR(c) dK(c) = dc dc
(3)
The KR(c) family could be generated from the σf(P) data after appropriately calibrating the coefficients Ψ and ξ . Finally, the R-curve was determined objectively as the common envelope of the tangency points of KR(c). 3. Results and discussion Fig. 2a shows the XRD patterns of pure c-BN sintered at 1400−1800 °C under a pressure of 50 MPa. Only c-BN and h-BN phases were observed in the sintered samples, and an intermediate metastable rhombohedral or wurtzite phase was not observed. The observed increase in the h-BN peak intensity (Fig. 2a) indicates an increase in the content of h-BN and confirms that the transformation of c-BN to h-BN is enhanced as the sintering temperature is increased. The significant increase of the h-BN peak intensity in the bulk verifies that the transformation of c-BN to h-BN starts at 1500 °C, becomes more significant at 1600 °C, and is complete at 1800 °C. The transformation mechanism and temperature are similar to those previously reported [25,29]. XRD patterns were also obtained to identify the crystalline phases of the B4C specimens
Fig. 3. (a) Relative densities of B4C and in situ B4C/h-BN composites with 5 vol. % c-BN manufactured at different temperatures, and relative density of B4C/hBN composite fabricated by direct introduction of 5 vol.% h-BN sintered at 1800 °C (blue triangle). Microstructures of fracture surfaces of monolithic B4C sintered at (b) 1600 °C and (d) 1800 °C and B4C-based ceramics with 5 vol.% c-BN sintered at (c) 1600 °C and (e) 1800 °C. (f) Sintering temperature profile and shrinkage curves of B4C and in situ B4C/h-BN composites with 5 vol.% c-BN (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
Fig. 2. XRD patterns of (a) pure bulk c-BN sintered at 1400–1800 °C and (b) B4C-based ceramics sintered at 1800 °C with 2–10 vol.% added c-BN. 3
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
prepared with different c-BN volume fractions. All the samples were almost completely densified, except pure B4C, and the densities of in situ B4C/h-BN ceramics did not vary significantly. The hardness of the pure B4C ceramic was 35.6 GPa, similar to a previously reported result [30] and 7 GPa higher than that of hot-pressed B4C [6]. Generally, the hardness of a composite mainly depends on the relative density, grain characteristics, intrinsic hardness, and content of a second phase in the matrix. As shown in Table 1, the hardness of the B4C/h-BN composites is lower than that of the pure B4C ceramic, showing a significant decrease as the content of in situ synthesized h-BN in the B4C matrix increases, even though the B4C/h-BN composites are obviously denser than pure B4C. Thus, the h-BN phase plays a more important role in the hardness of the B4C-based ceramics than density or grain refinement. Because of its large content of in situ synthesized h-BN and homogeneous dispersion, the B4C/h-BN composite with 10 vol.% c-BN has the lowest hardness of only 26.5 GPa, which is much lower than the intrinsic hardness of B4C. As shown in Table 1, the flexural strength and fracture toughness of the B4C ceramic are 354 MPa and 2.42 MPa⋅ m1/2, respectively, which are similar to the values previously reported for pure B4C ceramics [6,30]. Both the flexural strength and fracture toughness of the B4C/h-BN composites increased constantly with the addition of c-BN. However, several reports [22,23,31,32] demonstrated that the introduction of h-BN in ceramics could arouse the deterioration of flexural strength, which derived from internal residual pores behaving as crack-initiation points. These internal pores from agglomerate and low relative density could cause the formation of submicronsize notch. In this study, the transformation of c-BN to h-BN utilized as a new in situ route for introducing h-BN into B4C ceramics, not only makes h-BN uniformly dispersed in B4C, but also effectively increases the relative density of B4C/h-BN composite ceramics comparing pure B4C. Therefore, the addition of homogeneously dispersed h-BN into B4C ceramic (shown in Fig. 5) in this report enables to increase the flexural strength, which has also been demonstrated by previous works [20,33]. According to the data in Table 1, the optimal mechanical properties were achieved with the addition of 10 vol.% c-BN, with the flexural strength and fracture toughness reaching approximately 494 MPa and 4.2 MPa m1/2, respectively. Thus, the in situ synthesized h-BN significantly improved the mechanical properties of the B4C ceramics. However, considering the hardness, flexural strength, and fracture toughness of the ceramic simultaneously, the sample with 5 vol.% c-BN as the additive had the optimal integrated mechanical properties with a hardness of 30.5 GPa, a bending strength of 470 MPa, and a fracture toughness of 3.84 MPa⋅ m1/2. Fig. 4a shows a log plot of the fracture strength against the indentation load for the in situ B4C/h-BN ceramics fabricated with 5 vol.% c-BN. The intersection point, logP*, divides the plot into two regions. In the left region, where the samples have a low indentation load, the fracture strength is mainly controlled by the microstructure and is nearly the same that of the samples with no indentation load. In the right region, the fracture strength of the samples with high indention loads is dominated by external flaws. According to the indentation fracture mechanics, the relation of the fracture strength (σf) and the indentation load (P) can be described as following formula [34]:
with varying amounts of added c-BN after sintering at 1800 °C, as shown in Fig. 2b. Only the B4C phase was detected in the pure B4C ceramics. However, for the B4C-based ceramics with 2, 5, 7, and 10 vol. % c-BN, both B4C and h-BN phases were detected in the sintered specimens. This finding indicates that the c-BN phase in the composites is completely transformed into the h-BN phase. Fig. 3a shows not only the relative densities of the monolithic B4C ceramics and the B4C-based ceramics with 5 vol.% c-BN sintered at 1400–1800 °C, but also the relative densities of the B4C-based ceramics with 5 vol.% h-BN sintered at 1800 °C (blue triangle). The relative densities of the B4C/h-BN ceramics (93.1 %) with direct introduction of 5 vol.% h-BN is lower than that of both monolithic B4C and B4C/h-BN ceramic fabricated with 5 vol.% c-BN sintered at 1800 °C, which testifies that the direct introduction of h-BN does not contribute to the densification of B4C. Furthermore, the relative densities of the composites with 5 vol.% c-BN are obviously higher than those of the pure B4C samples when sintered at 1500–1800 °C. The relative densities of the B4C ceramics with 5 vol.% c-BN prepared at 1500, 1600, 1700, and 1800 °C are 70.5 %, 92.3 %, 97.2 %, and 98.8 % respectively, whereas those of the pure B4C ceramics are 64.0 %, 76.2 %, 87.3 %, and 94.2 %, respectively. However, the relative densities of the composite with 5 vol.% c-BN and the pure B4C ceramic are similar when sintered at 1400 °C, which is consistent with the XRD results (Fig. 2a), which revealed that the transformation of c-BN to h-BN only started at 1500 °C and was completed by 1800 °C. Thus, these findings indicate that the transformation of c-BN to h-BN is beneficial for the densification of B4C. Fig. 3b–e show typical SEM images of the fracture surface of pure B4C and the in situ B4C/h-BN composites with 5 vol.% c-BN sintered at 1600 and 1800 °C. As shown in the SEM images, most of the pores are eliminated from the B4C/h-BN composites, as there are significantly fewer pores present in the fracture surfaces of the B4C/h-BN composites than in those of the pure B4C ceramics sintered at same temperature, confirming the changes in relative densities observed in Fig. 3a. When the sintering temperature reached 1800 °C, in situ bulk B4C/h-BN with 5 vol.% c-BN nearly achieved full densification, whereas pure B4C does not. The results in Fig. 3a are also confirmed by the sintering temperature profile and the shrinkage curves of B4C and the B4C composites with 5 vol.% c-BN in Fig. 3f. Greater shrinkage is observed for the B4C composites with 5 vol.% c-BN than for B4C between 1400 and 1800 °C. The observation of complete densification for the B4C composites with 5 vol.% c-BN at 1800 °C but not for B4C indicates that the c-BN to h-BN phase transformation contributed to the densification of B4C. To eliminate the effect of silica on the densification of these composites, ICP is employed to measure the Si impurity contents of the original B4C powders and B4C powders dealt with agate ball. The results of original one and the dealt one are 0.02 wt.% and 0.04 wt.%, respectively, which indicated nearly no difference in Si content after milling with agate ball and no effect of silica on the densification of these composites. According to the previous report [29], phase transformation of c-BN to hBN is a way of decreasing the total free energy, which indicating the phase transformation is an exothermic process. Therefore the attribution of the phase transformation to the densification of B4C may originate from released energy. The more detailed mechanism may exist and needs to be studied by more future research. Table 1 shows the density, elastic modulus, Vickers hardness, fracture toughness, and flexural strength of the B4C/h-BN ceramics
σf ∝P−β
(4)
β is equal to 1/3 for the material without resistance behavior, while the
Table 1 Density and mechanical properties of the B4C-based ceramics fabricated with different c-BN contents (vol.%) at 1800 °C. Sample B4C B4C-2 vol.% c-BN B4C-5 vol.% c-BN B4C-7 vol.% c-BN B4C-10 vol.% c-BN
Theoretical density (g/cm3)
Relative density (%)
Elastic modulus (GPa)
Hardness (GPa)
Flexural strength (MPa)
Fracture toughness (MPa·m1/2)
2.520 2.513 2.503 2.496 2.487
94.30 99.72 98.88 98.64 98.35
427 ± 16 417 ± 4 401 ± 7 388 ± 9 326 ± 13
35.6 ± 0.6 33.2 ± 0.3 30.5 ± 0.7 29.1 ± 0.5 25.6 ± 0.4
354 ± 34 467 ± 25 470 ± 20 488 ± 17 494 ± 10
2.42 ± 0.70 3.13 ± 0.46 3.84 ± 0.55 4.10 ± 0.34 4.20 ± 0.27
4
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
Fig. 4. (a) Effect of the indentation load on the fracture strength of pure B4C and the in situ B4C/h-BN composites with 5 vol.% c-BN. (b) Toughness curve family diagram for the in situ B4C/h-BN composites with 5 vol.% c-BN. The R-curve behaviour is revealed by the envelope of the tangency points. (c) Load-displacement curves for five indentation tests on in situ B4C/h-BN composites with 5 vol.% c-BN. (d) Typical load-displacement curves for indentation tests on in situ B4C/h-BN composites with 0, 2, 5, 7, and 10 vol.% c-BN.
β is samller than 1/3 for the material with resistance behavior. And the relation between the fracture strength (σf) and the indentation load (P) in the high indentation load region (Fig. 4a) can be expressed using the following formula:
logσf = 2.688 − 0.192 logP
ceramics. In Fig. 5, the dark grey phase corresponds to B4C, whereas the bright flakes correspond to h-BN. Thus, these images also reveal that the h-BN phase is homogeneously dispersed in the B4C matrix without severe aggregation of the h-BN flakes. The fracture surface of pure B4C is very flat, indicating a typical transgranular fracture (Fig. 5a). In contrast, the fracture surfaces of the B4C/h-BN composites are not as flat as that of pure B4C, which could be attributed to the effect of h-BN on the crack propagation path. Generally, microcracks propagate along the weak interface between h-BN and B4C grains, and deflect when meeting h-BN flakes and hard B4C. Thus, the h-BN flakes formed through the phase transformation of c-BN are essential for enhancing the mechanical properties of B4C-based ceramics. A TEM analysis was performed on the interface between B4C and the h-BN in the B4C/h-BN composites, as shown in Fig. 6. The EDS analysis (Fig. 6a, inset) indicates that particles B and A are h-BN and B4C, respectively. The HRTEM image of the boundary in Fig. 6b shows a clean and narrow grain boundary between the B4C and h-BN grains, suggesting strong bonding between the B4C and h-BN grains. Because the boundary strength between the matrix and the second phase significantly affects the mechanical properties of a composite, the B4C/hBN composites achieve enhanced mechanical properties. To reveal the grain boundaries and morphologies of the B4C-based ceramics, electrolytic etching was performed on the polished surface. Representative SEM images of the etched surfaces of the pure B4C ceramic and the B4C/h-BN composite with 2 vol.% c-BN are shown in Fig. 7. The B4C-based composite with added c-BN (Fig. 7b) showed a finer and less faceted grain structure than the pure B4C ceramic (Fig. 7a), with mean grain sizes of approximately 3.2 and 2.2 μm, respectively. When raw c-BN is added to the B4C matrix, the in situ synthesized h-BN grains are located at B4C boundaries and efficiently
(5)
The fit of the data in Fig. 4a to this formula gives a slope of -0.192, which is larger than -1/3, indicating that the in situ B4C/h-BN ceramics fabricated with 5 vol.% c-BN exhibit rising R-curve behaviour. The calculated families of KR(c) curves and R-curves for the in situ B4C/h-BN ceramics are shown in Fig. 4b. The envelope of the tangency points reveals rising R-curve behaviour, which suggests that the fracture toughness will increase with an increase in the crack length. In addition, Fig. 4c shows the load and displacement data recorded during five nano-indentation tests, and good repeatability is observed. Typical load-displacement curves for the B4C-based ceramics, obtained using a depth-sensing nano-indentation tester with a peak load of 2 N, are shown in Fig. 4d. Based on the Oliver–Pharr method [26], the nanoindentation tests were used to determine the elastic modulus of each B4C-based ceramic (Table 1). A high elastic modulus of approximately 400 GPa and degradation were observed at higher in situ h-BN contents. Furthermore, higher permanent deformation and higher elastoplasticity were also observed at higher in situ h-BN contents, which is consistent with the hardness results in Table 1 and reveals that the B4C/h-BN ceramics exhibit increased elastoplasticity. The R-curve behaviour and the load-displacement curves both confirm that in situ synthesized h-BN enhances the mechanical properties of the B4C ceramics. As shown by the SEM images of the fracture surfaces of the pure B4C and B4C/h-BN ceramics (Fig. 5), pores were hardly observed in the B4C/ h-BN composites, whereas some pores were observed in the pure B4C 5
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
Fig. 5. SEM images of the fracture surfaces of B4C/h-BN ceramics fabricated with (a) 0, (b) 2, (c) 5, (d) 7, and (e) 10 vol.% c-BN.
inhibit the growth of B4C grains. Generally, finer grains will significantly improve the mechanical properties of ceramics. Fig. 8 shows various microstructural evidence of the toughening by h-BN platelets. Homogeneously dispersed h-BN platelets absorb the crack propagation energy via different toughening mechanisms such as
pulling out of h-BN, crack bridging, crack deflection, and crack branching. These mechanisms are similar for all filler-reinforced ceramics, but differ in their frequency and effectiveness of load transfer between the matrix and the filler. The h-BN platelets are insert between the matrix grains because of the stress applied by neighbouring B4C
Fig. 6. TEM images of the morphology and grain boundary structure between B4C and h-BN. Inset: EDS analysis identifying particles A and B as B4C and h-BN, respectively. 6
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
Fig. 7. SEM images of the etched surfaces of (a) pure B4C and (b) B4C/h-BN with 2 vol.% c-BN.
Such an interaction decreases the stress intensity at the crack top and even at the temporary crack top. Thus, it is believed that such a crack deflection mechanism can help to increase the toughness of B4C-based ceramics by creating a curving path. Another toughness mechanism, crack branching, can disperse and decrease the stress at the crack top, as observed in Fig. 8b. This mechanism can also significantly contribute to improving fracture toughness of B4C-based ceramics.
grains. Moreover, in situ synthesized h-BN platelets have a high 2D aspect ratio and are homogeneously dispersed in the B4C matrix. Thus, the contact area between h-BN and the B4C matrix is large and is expected to be greater than that between fibres and a matrix. Thus, the energy required to pull out an h-BN platelet is expected to be greater than that required to pull out a fibre or h-BN nanotube. The pulling out of h-BN platelets is believed to be an effective toughening mechanism for the composites, as shown in Fig. 8a, owing to the friction resistance at the interface between the h-BN platelets and the B4C matrix grains, which significantly dissipates the crack propagation energy. Additionally, h-BN platelets could also behave as bridging ligaments in B4C-based ceramics. As h-BN is too rigid to be broken by internal cracks in B4C-based ceramics, h-BN can bridge the cracks and absorb the crack propagation energy, as shown in Fig. 8c. Crack deflection and a change in the direction of crack propagation are shown in Fig. 8d when the crack meets h-BN platelets. The deflection length and deflection angle range from 0.5 μm to several microns and from 20° to 90°, respectively.
4. Conclusion In this study, B4C/h-BN composites prepared using the phase transformation of c-BN to h-BN were characterized. The SEM results and the density analysis revealed that the densification behaviour of B4Cbased ceramics was improved by the energy released by the transformation of c-BN to h-BN, and almost completely densified B4C/h-BN composites were fabricated by SPS. The results showed that in situ synthesized h-BN can be homogeneously dispersed in the B4C matrix
Fig. 8. Toughening mechanisms for B4C/h-BN composites: (a) pulling out of h-BN platelets, (b) crack branding, (c) crack bridging, and (d) crack deflection. 7
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
microstructure. Moreover, finer B4C grains were obtained in the composite owing to the inhibition of B4C grain growth by in situ synthesized of h-BN at the grain boundaries. The flexural strength and fracture toughness of the B4C ceramics reinforced by in situ synthesized h-BN were significantly higher than those of pure B4C. The sample prepared with 5 vol.% c-BN as an additive exhibited excellent integrated mechanical properties with a hardness of 30.5 GPa, a bending strength of 470 MPa, and a fracture toughness of 3.84 MPa⋅ m1/2. Increasing the cBN content further decreased the hardness but did not significantly improve the bending strength and fracture toughness. Various toughening mechanisms, such as pulling out of h-BN, crack deflection, crack branching, and crack bridging, were responsible for the enhanced mechanical properties of the B4C/h-BN composites. This study showed that the in situ h-BN from c-BN was effective to reinforce B4C ceramics as a structural analogue of GPLs. The exothermal phase transformation could be a new method to improve the densification process of the ceramics and simultaneously offer reinforcement material for ceramic matrix.
Mater. 17 (1999) 27–32, https://doi.org/10.1016/S0263-4368(98)00052-3. [12] H. Zhao, K. Hiragushi, Y. Mizota, Densification and mechanical properties of mullite/TiO2-coated B4C composites, J. Eur. Ceram. Soc. 23 (2003) 1485–1490, https://doi.org/10.1016/S0955-2219(02)00307-2. [13] W.-M. Guo, L.-X. Wu, Y. You, H.-T. Lin, G.-J. Zhang, Three-step reactive hot pressing of B4C–ZrB2 ceramics, J. Eur. Ceram. Soc. 36 (2016) 951–957, https://doi. org/10.1016/j.jeurceramsoc.2015.11.022. [14] L. Levin, N. Frage, M.P. Dariel, A novel approach for the preparation of B4C-based cermets, Int. J. Refract. Met. Hard Mater. 18 (2000) 131–135, https://doi.org/10. 1016/S0263-4368(00)00012-3. [15] M.W. Chen, J.W. McCauley, J.C. LaSalvia, K.J. Hemker, Microstructural characterization of commercial hot-pressed boron carbide ceramics, J. Am. Ceram. Soc. 88 (2005) 1935–1942, https://doi.org/10.1111/j.1551-2916.2005.00346.x. [16] A. Kovalčíková, R. Sedlák, P. Rutkowski, J. Dusza, Mechanical properties of boron carbide+graphene platelet composites, Ceram. Int. 42 (2016) 2094–2098, https:// doi.org/10.1016/j.ceramint.2015.09.139. [17] R. Sedlák, A. Kovalčíková, E. Múdra, P. Rutkowski, A. Dubiel, V. Girman, R. Bystrický, J. Dusza, Boron carbide/graphene platelet ceramics with improved fracture toughness and electrical conductivity, J. Eur. Ceram. Soc. 37 (2017) 3773–3780, https://doi.org/10.1016/j.jeurceramsoc.2017.04.061. [18] M. Wojtoniszak, X. Chen, R.J. Kalenczuk, A. Wajda, J. Łapczuk, M. Kurzewski, M. Drozdzik, P.K. Chu, E. Borowiak-Palen, Synthesis, dispersion, and cytocompatibility of graphene oxide and reduced graphene oxide, Colloids Surf., B 89 (2012) 79–85, https://doi.org/10.1016/j.colsurfb.2011.08.026. [19] M. Kitiwan, A. Iton, T. Goto, Spark plasma sintering of TiN–TiB2–hBN composites and their properties, Ceram. Int. 41 (2015) 4498–4503, https://doi.org/10.1016/j. ceramint.2014.11.144. [20] B. Lee, D. Lee, J.H. Lee, H.J. Ryu, S.H. Hong, Enhancement of toughness and wear resistance in boron nitride nanoplatelet (BNNP) reinforced Si3N4 nanocomposites, Sci. Rep. 6 (2016) 27609, https://doi.org/10.1038/srep27609. [21] D. Lee, S.H. Song, J. Hwang, S.H. Jin, K.H. Park, B.H. Kim, S.H. Hong, S. Jeon, Enhanced mechanical properties of epoxy nanocomposites by mixing noncovalently functionalized boron nitride nanoflakes, Small 9 (2013) 2602–2610, https://doi. org/10.1002/smll.201203214. [22] W.S. Cho, Z.H. Piao, K.J. Lee, Y.C. Yoo, J.H. Lee, M.W. Cho, Y.C. Hong, K. Park, W.S. Hwang, Microstructure and mechanical properties of AlN–hBN based machinable ceramics prepared by pressureless sintering, J. Eur. Ceram. Soc. 27 (2007) 1425–1430, https://doi.org/10.1016/j.jeurceramsoc.2006.05.044. [23] X. Li, Y. Gao, W. Pan, X. Wang, L. Song, Z. Zhong, S. Wu, Fabrication and characterization of B4C-based ceramic composites with different mass fractions of hexagonal boron nitride, Ceram. Int. 41 (2015) 27–36, https://doi.org/10.1016/j. ceramint.2014.07.042. [24] B. Zhong, G.L. Zhao, X.X. Huang, L. Xia, X.H. Tang, S.C. Zhang, G.W. Wen, Microstructure and mechanical properties of ZTA/BN machinable ceramics fabricated by reactive hot pressing, J. Eur. Ceram. Soc. 35 (2015) 641–649, https://doi. org/10.1016/j.jeurceramsoc.2014.09.002. [25] J.C. Garrett, I. Sigalas, A.K. Wolfrum, M. Herrmann, Effect of cubic boron nitride grain size in the reinforcing of α-Sialon ceramics sintered via SPS, J. Eur. Ceram. Soc. 35 (2015) 451–462, https://doi.org/10.1016/j.jeurceramsoc.2014.09.036. [26] F. Khodabakhshi, A.P. Gerlich, D. Verma, M. Haghshenas, Nano-indentation behavior of layered ultra-fine grained AA8006 aluminum alloy and AA8006-B4C nanostructured nanocomposite produced by accumulative fold forging process, Mater. Sci. Eng., A 744 (2019) 120–136, https://doi.org/10.1016/j.msea.2018.12.013. [27] S. Li, C. Wei, W. Wang, P. Wang, Y. Liu, S. Li, G. Wen, Fracture toughness and Rcurve behavior of laminated ZrB2–SiC/SiCw ceramic, J. Alloys. Compd. 784 (2019) 96–101, https://doi.org/10.1016/j.jallcom.2018.12.304. [28] T. Jiang, Z. Jin, J. Yang, G. Qiao, Mechanical property and R-curve behavior of the B4C/BN ceramics composites, Mater. Sci. Eng., A 494 (2008) 203–216, https://doi. org/10.1016/j.msea.2008.04.047. [29] I.A. Petrusha, Features of a cBN-to-graphite-like BN phase transformation under pressure, Diamond Relat. Mater. 9 (2000) 1487–1493, https://doi.org/10.1016/ S0925-9635(00)00273-9. [30] Y. Tan, H. Zhang, S. Peng, Electrically conductive graphene nanoplatelet/boron carbide composites with high hardness and toughness, Scr. Mater. 114 (2016) 98–102, https://doi.org/10.1016/j.scriptamat.2015.12.008. [31] A. Kovalčíková, J. Balko, C. Balázsi, P. Hvizdoš, J. Dusza, Influence of hBN content on mechanical and tribological properties of Si3N4/BN ceramic composites, J. Eur. Ceram. Soc. 34 (2014) 3319–3328, https://doi.org/10.1016/j.jeurceramsoc.2014. 02.021. [32] Y.K. Seo, Y.W. Kim, K.J. Kim, W.S. Seo, Electrically conductive SiC-BN composites, J. Eur. Ceram. Soc. 36 (2016) 3879–3887, https://doi.org/10.1016/j.jeurceramsoc. 2016.06.040. [33] T. Jiang, Z. Jin, J. Yang, G. Qiao, Investigation on the preparation and machinability of the B4C/BN nanocomposites by hot-pressing process, J. Mater. Process. Tech. 209 (2009) 561–571, https://doi.org/10.1016/j.jmatprotec.2008.02.026. [34] L.M. Braun, S.J. Bennison, B.R. Lawn, Objective evaluation of short-crack toughness curves using indentation flaws: case study on alumina-based ceramics, J. Am. Ceram. Soc. 75 (1992) 3049–3057, https://doi.org/10.1111/j.1151-2916.1992. tb04385.x.
Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements This work was financially supported by the National Natural Science Foundation of China (51502220, 51521001, 51672197), the Ministry of Science and Technology of the People’s Republic of China (2015DFR50650), the Self-determined and Innovative Research Funds of WUT (2017II17XZ), and the Open Project Program of the Key Laboratory of Inorganic Functional Materials and Devices, Chinese Academy of Sciences (Grant No. KLIFMD201606). References [1] B. Niu, F. Zhang, J. Zhang, W. Ji, W. Wang, Z. Fu, Ultra-fast densification of boron carbide by flash spark plasma sintering, Scr. Mater. 116 (2016) 127–130, https:// doi.org/10.1016/j.scriptamat.2016.02.012. [2] K.H. Kim, J.H. Chae, J.S. Park, J.P. Ahn, K.B. Shim, Sintering behavior and mechanical properties of B4C ceramics fabricated by spark plasma sintering, J. Ceram. Process. Res. 10 (2009) 716–720. [3] 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, https://doi.org/ 10.1179/095066009X12506721665211. [4] V. Domnich, S. Reynaud, R.A. Haber, M. Chhowalla, Boron carbide: structure, properties, and stability under stress, J. Am. Ceram. Soc. 94 (2011) 3605–3628, https://doi.org/10.1111/j.1551-2916.2011.04865.x. [5] T.K. Roy, C. Subramanian, A.K. Suri, Pressureless sintering of boron carbide, Ceram. Int. 32 (2006) 227–233, https://doi.org/10.1016/j.ceramint.2005.02.008. [6] Y. Tan, H. Luo, H. Zhang, S. Peng, Graphene nanoplatelet reinforced boron carbide composites with high electrical and thermal conductivity, J. Eur. Ceram. Soc. 36 (2016) 2679–2687, https://doi.org/10.1016/j.jeurceramsoc.2016.04.036. [7] W. Ji, S.S. Rehman, W. Wang, H. Wang, Y. Wang, J. Zhang, F. Zhang, Z. Fu, Sintering boron carbide ceramics without grain growth by plastic deformation as the dominant densification mechanism, Sci. Rep. 5 (2015) 15827, https://doi.org/ 10.1038/srep15827. [8] M. Mashhadi, E. Taheri-Nassaj, V.M. Sglavo, H. Sarpoolaky, N. Ehsani, Effect of Al addition on pressureless sintering of B4C, Ceram. Int. 35 (2009) 831–837, https:// doi.org/10.1016/j.ceramint.2008.03.003. [9] F. Ye, Z. Hou, H. Zhang, L. Liu, Densification and mechanical properties of spark plasma sintered B4C with Si as a sintering aid, J. Am. Ceram. Soc. 93 (2010) 2956–2959, https://doi.org/10.1111/j.1551-2916.2010.03931.x. [10] V.B. Fedorus, G.N. Makarenko, S.P. Gordienko, É.V. Marek, I.I. Timofeeva, Interaction between boron carbide and chromium oxide, Powder Metall. Met. Ceram. 34 (1996) 637–639, https://doi.org/10.1007/BF00559492. [11] P. Rogl, H. Bittermann, Ternary metal boron carbides, Int. J. Refract. Met. Hard
8
Contents lists available at ScienceDirect
Journal of the European Ceramic Society journal homepage: www.elsevier.com/locate/jeurceramsoc
Original article
Densification and mechanical properties of boron carbide prepared via spark plasma sintering with cubic boron nitride as an additive Jinchang Suna, Bo Niua, Lin Renb,*, Jinyong Zhanga, Liwen Leia, Fan Zhanga,* a b
State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan, 430070, PR China School of Science, Wuhan University of Technology, Wuhan, 430070, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: B4C-based ceramic Phase transformation Spark plasma sintering Densification Mechanical properties
Hexagonal boron nitride (h-BN) can reinforce boron carbide (B4C) ceramics, but homogeneous dispersion of hBN is difficult to achieve using conventional methods. Herein, B4C/h-BN composites were manufactured via the transformation of cubic (c-) BN during spark plasma sintering at 1800 °C. The effects of the c-BN content on the microstructure, densification, and mechanical properties of B4C/h-BN composites were evaluated. In situ synthesized h-BN platelets were homogeneously dispersed in the B4C matrix and the growth of B4C grains was effectively suppressed. Moreover, the c-BN to h-BN phase transformation improved the sinterability of B4C. The sample with 5 vol.% c-BN exhibited excellent integrated mechanical properties (hardness of 30.5 GPa, bending strength of 470 MPa, and fracture toughness of 3.84 MPa⋅ m1/2). Higher c-BN contents did not significantly affect the bending strength and fracture toughness but clearly decreased the hardness. The main toughening mechanisms were crack deflection, crack bridging, and pulling out of h-BN.
1. Introduction Boron carbide (B4C) has attracted much attention because of its low density (2.52 g/cm3), high hardness (35 GPa), excellent wear stability, great chemical resistance, and high melting point (2450 °C) [1–3]. Owing to these outstanding properties, B4C ceramics have been used in many structural applications such as aerospace craft, shield materials, and abrasive and polishing media for hard and wear-resistant materials [4,5]. B4C ceramics are also currently used in neutron detectors because of their high neutron absorption cross sections [6]. However, the practical applications of B4C ceramics are severely hindered by numerous drawbacks such as low fracture toughness, low flexural strength, and poor sinterability. Three methods are commonly used to obtain highly dense B4C ceramics: (i) increasing the sintering temperature, (ii) applying a high pressure (≥80 MPa) [7], and (iii) using additives, including Ti, Al, Fe, Ni, Cu, Si, TiO2, Cr2O3, and ZrB2 [8–14]. However, a high temperature leads to excessive growth of the B4C grains, which could reduce the strenght [7], and it is difficult to apply a high pressure in standard sintering equipment. In contrast, additives are widely used in the production of B4C because they can reduce the sintering temperature to
below 2150 °C [15] or enhance the mechanical properties of the manufactured ceramics. In recent years, graphene platelets (GPLs) have received considerable attention as efficient reinforcement materials for B4C ceramics because of their outstanding properties such as being lightweight and having a two-dimensional (2D) high aspect ratio and a high tensile strength [16,17]. However, as reinforcement materials for ceramic matrix composites, the applicability of GPLs is limited, mainly owing to the difficulties in realizing heterogeneous dispersion for its large surface area. Furthermore, GPLs exhibit poor high-temperature structure stability [18] and can react with B4C, weakening the reinforcement ability of GPLs in B4C. Currently, BN platelets, a structural analogue of GPLs, are usually used as a reinforcement material for ceramics. Mettaya Kitiwan [19] used spark plasma sintering method to fabricate Titanium nitride (TiN)/Titanium diboride (TiB2)/hexagonal boron nitride (h-BN) composites at 1700 °C under 100 MPa. A fracture toughness of 4.3 MPa⋅ m1/2 was reported for composites with 15 vol.% of h-BN, which is higher than the 2.8 MPa⋅ m1/2 of the pure TiN/TiB2 composite. Using an HP process, Lee et al. [20] prepared Si3N4/h-BN with a fracture toughness of 9.76 MPa⋅ m1/2 for the composite with 2 wt.% h-BN, corresponding to a 24.7 % increase compared with the value for pure
Abbreviations: (GPL), graphene platelet; (2D), two-dimensional; (HP), hot-pressing; (h-BN), hexagonal boron nitride; (c-BN), cubic boron nitride; (SPS), spark plasma sintering; (XRD), X-ray diffraction; (SEM), scanning electron microscopy; (TEM), transmission electron microscopy; (EDS), energy-dispersive X-ray spectroscopy ⁎ Corresponding authors. E-mail addresses: [email protected] (L. Ren), [email protected] (F. Zhang). https://doi.org/10.1016/j.jeurceramsoc.2019.12.047 Received 31 May 2019; Received in revised form 18 December 2019; Accepted 23 December 2019 0955-2219/ © 2019 Published by Elsevier Ltd.
Please cite this article as: Jinchang Sun, et al., Journal of the European Ceramic Society, https://doi.org/10.1016/j.jeurceramsoc.2019.12.047
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
Germany). A high-strength graphite die with an inner diameter of 50 mm was filled with the homogenized mixed powders and then heated to 1800 °C. To evaluate the effect of c-BN on the densification of B4C ceramics, pure B4C powders and B4C/5 vol.% c-BN composites were also prepared using SPS at 1400–1800 °C. All samples were maintained at the sintering temperature for 10 min with a heating rate of 100 °C/ min. A uniaxial pressure of 50 MPa was applied throughout the sintering cycle. In addition, B4C/5 vol.% h-BN composite and pure c-BN powders were sintered at 1800 °C and 1400–1800 °C, respectively, for 10 min under a pressure of 50 MPa using the same SPS apparatus and heating rate. The Archimedes method was used to determine the densities of the B4C-based ceramics in deionized water. The crystalline phases of the powder mixtures and the as-sintered samples were identified by X-ray diffraction (XRD, Ultima III, Rigaku, Japan) with Cu Kα radiation. The microstructures of the prepared samples were characterized by scanning electron microscopy (SEM, Hitachi 3400, Japan). A thin foil of B4C/h-BN composite was prepared by mechanical thinning using a focused ion beam and analysed by transmission electron microscopy (TEM, JEOL JEM-2010HT, Japan) coupled with energy-dispersive X-ray spectroscopy (EDS). To visualize the grain boundaries clearly, the surfaces of the polished sintered samples were electrically etched in a 1 wt. % NaOH solution with a current density of 0.1 A/cm2 for 10–30 s. The grain sizes of the sintered B4C-based samples were determined from the SEM images by averaging the lengths and widths of at least 100 randomly selected grains. The mechanical properties of the bulk materials containing different quantities of in situ synthesized h-BN were analysed by measuring their hardness, fracture toughness, flexural strength, and elastic modulus. The Vickers hardness (Hv) was determined with a Vickers hardness testing machine (Wolpert-430SV, USA) using an indentation load of 1 kg for a dwell time of 15 s, and the hardness values were calculated from 10 indentations. The flexural strength, elastic modulus, and fracture toughness were measured by three-point bending tests using a ceramic test system (MTS 810, MTS, USA). The flexural strengths (KIC) of specimens with dimensions of 3 mm × 4 mm × 36 mm were measured with a testing span of 30 mm and crosshead speed of 0.5 mm/ min. Specimens with dimensions of 2.5 mm × 5 mm × 25 mm and a notch in the middle (0.15 mm thick and 2.5 mm long) were used for fracture toughness tests. The testing span was 20 mm and the crosshead speed was 0.5 mm/min. All the test samples were ground with a 15 μm diamond grinding wheel. The four edges were chamfered to a depth of 0.12 ± 0.03 mm and an angle of 45° ± 5° to eliminate stress at the edges of the samples. Average values were obtained from the results for at least six specimens. The elastic modulus (E) was characterized with a depth-sensing nano-indentation tester (NHT, CSEM, Switzerland) equipped with a Berkovich indenter, using a peak load of 2 N on the polished surface of each sample. At least five nano-indentations were performed for each sample. In each test, the indenter was driven into the sample surface at a rate of 200 mN/min and then unloaded at the same rate. The load and the indentation depth were simultaneously
Si3N4. This improvement in performance occurs because h-BN has a high 2D aspect ratio similar to that of GPLs, and the mechanical properties of h-BN are also comparable to those of GPLs. The elastic modulus of h-BN is approximately 700–900 GPa, whereas that of GPLs is approximately 1 TPa [21]. Additionally, h-BN is not able to react with B4C and can possess better high-temperature structure stability. These findings indicate that h-BN has excellent potential for reinforcing B4C ceramics. However, the homogeneous dispersion of h-BN powder into a B4C ceramic matrix is difficult using conventional ball milling. The hBN platelets can easily agglomerate in ceramic matrices, resulting in large flaws. These flaws initiate cracks and deteriorate the mechanical properties [22]. Moreover, because of their poor sinterability, h-BN platelets hinder the interdiffusion and rearrangement of ceramic matrix particles during the sintering process [23]. Zhong et al. [24] reported that this problem could be solved by synthesizing h-BN platelets viain situ chemical reaction of aluminum nitride (AlN) and boric acid (H3BO3) in the ceramic matrix, leading to a homogeneous distribution, a compact structure, and fine h-BN grains. According to a previous report [25], the hexagonal transformation of cubic boron nitrite (c-BN) occurs at 1550–1625 °C. Thus, in this study, instead of using chemical reaction routes, a new in situ route for introducing h-BN into B4C ceramics was developed via the transformation of c-BN to h-BN by spark plasma sintering (SPS). The effects of the c-BN additive on densification and the mechanical properties, such as the fracture toughness and flexural strength, of B4C matrix ceramics were evaluated. 2. Experimental The B4C/c-BN and B4C/5 vol.% h-BN composites were prepared using commercial powders: micron B4C powders (W1.5, Mudanjiang Jingangzuan Boron Carbide, China) and c-BN powders (Funik Ultrahard Materials, China), h-BN powders (Advanced Technology & Materials Co. Ltd., China). According to the manufacturers’ data, the B2O3 content in the raw B4C powder is less than 0.2 wt.%. The average grain sizes of the B4C, c-BN and h-BN powders were 1.6, 0.3 and 4.0 μm (both from suppliers’ data), respectively. Fig. 1 shows SEM images of starting materials. First, different quantities of c-BN (0, 2, 5, 7, and 10 vol.%) and 5 vol.% h-BN were added to the B4C powder. These powder mixtures were homogenized by wet ball milling in ethyl alcohol using agate balls for 24 h on low-speed roller ball mill (GMJ/B, Xianyang JinHong General Machinery Co., LTD, Xianyang, China). The resulting mud was dried in a rotary evaporator at 65 °C and then kept in a vacuum chamber at 60 °C for 24 h. Finally, the obtained powders were sieved through a 200 mesh sieve. The Si impurity contents of the original B4C powders and B4C powders dealt with agate ball measured by inductively coupled plasma-optical emission spectrometer (ICP, Optima 4300DV, PerkinElmer, American) are 0.02 wt.% and 0.04 wt.%, respectively. Bulk samples of in situ B4C/h-BN composites and pure B4C were manufactured using SPS processes (HPD 60/0 furnace, FCT Systeme,
Fig. 1. Secondary-electron SEM image of the raw materials: (a) B4C powder and (b) c-BN powder. 2
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
obtained and used for simulation of the elastic modulus (E) with the Oliver–Pharr method [26]. The R-curve behaviour of the materials was evaluated by the indention strength bending technique. Specimens with dimensions of 3 mm × 4 mm × 36 mm were indented at the centre of the polished surface with a Vickers diamond pyramid indentation under loads of 4.9, 9.8, 29.4, 49, 98, 196, and 294 N. Subsequently, the indentation strength of each specimen was tested using a three-point bending test with a span of 30 mm and crosshead speed of 0.5 mm/min. The average indentation strength was obtained from the results for five measured specimens at each indentation load. According to indentation fracture mechanics, the stress intensity factor of the Vickers indentation crack tip (K) can be expressed by the following formula [27]:
K(c) =K σ + Kr =ΨσA c1/2 + ξ(
E 1/2 -3/2 ) Pc H
(1)
where K σ is the stress intensity factor resulting from the applied stress, Kr is the residual stress intensity factor at the indentation of a crack, KR(c) is the fracture toughness curve, Ψ is a material-independent constant that depends on the crack and specimen shape, which is equal to 1.24 [28], c is the size of the crack, P is a given indentation load, ξ is a non-dimensional geometry constant equal to 0.016 [28], E is the Young’s modulus, H is the hardness, and σA is the applied stress. For a given load (P), failure occurred at an applied stress σA = σf, which also satisfied the ‘common tangent’ according to the following equations [27]:
KR(c) =K(c)
(2)
dKR(c) dK(c) = dc dc
(3)
The KR(c) family could be generated from the σf(P) data after appropriately calibrating the coefficients Ψ and ξ . Finally, the R-curve was determined objectively as the common envelope of the tangency points of KR(c). 3. Results and discussion Fig. 2a shows the XRD patterns of pure c-BN sintered at 1400−1800 °C under a pressure of 50 MPa. Only c-BN and h-BN phases were observed in the sintered samples, and an intermediate metastable rhombohedral or wurtzite phase was not observed. The observed increase in the h-BN peak intensity (Fig. 2a) indicates an increase in the content of h-BN and confirms that the transformation of c-BN to h-BN is enhanced as the sintering temperature is increased. The significant increase of the h-BN peak intensity in the bulk verifies that the transformation of c-BN to h-BN starts at 1500 °C, becomes more significant at 1600 °C, and is complete at 1800 °C. The transformation mechanism and temperature are similar to those previously reported [25,29]. XRD patterns were also obtained to identify the crystalline phases of the B4C specimens
Fig. 3. (a) Relative densities of B4C and in situ B4C/h-BN composites with 5 vol. % c-BN manufactured at different temperatures, and relative density of B4C/hBN composite fabricated by direct introduction of 5 vol.% h-BN sintered at 1800 °C (blue triangle). Microstructures of fracture surfaces of monolithic B4C sintered at (b) 1600 °C and (d) 1800 °C and B4C-based ceramics with 5 vol.% c-BN sintered at (c) 1600 °C and (e) 1800 °C. (f) Sintering temperature profile and shrinkage curves of B4C and in situ B4C/h-BN composites with 5 vol.% c-BN (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
Fig. 2. XRD patterns of (a) pure bulk c-BN sintered at 1400–1800 °C and (b) B4C-based ceramics sintered at 1800 °C with 2–10 vol.% added c-BN. 3
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
prepared with different c-BN volume fractions. All the samples were almost completely densified, except pure B4C, and the densities of in situ B4C/h-BN ceramics did not vary significantly. The hardness of the pure B4C ceramic was 35.6 GPa, similar to a previously reported result [30] and 7 GPa higher than that of hot-pressed B4C [6]. Generally, the hardness of a composite mainly depends on the relative density, grain characteristics, intrinsic hardness, and content of a second phase in the matrix. As shown in Table 1, the hardness of the B4C/h-BN composites is lower than that of the pure B4C ceramic, showing a significant decrease as the content of in situ synthesized h-BN in the B4C matrix increases, even though the B4C/h-BN composites are obviously denser than pure B4C. Thus, the h-BN phase plays a more important role in the hardness of the B4C-based ceramics than density or grain refinement. Because of its large content of in situ synthesized h-BN and homogeneous dispersion, the B4C/h-BN composite with 10 vol.% c-BN has the lowest hardness of only 26.5 GPa, which is much lower than the intrinsic hardness of B4C. As shown in Table 1, the flexural strength and fracture toughness of the B4C ceramic are 354 MPa and 2.42 MPa⋅ m1/2, respectively, which are similar to the values previously reported for pure B4C ceramics [6,30]. Both the flexural strength and fracture toughness of the B4C/h-BN composites increased constantly with the addition of c-BN. However, several reports [22,23,31,32] demonstrated that the introduction of h-BN in ceramics could arouse the deterioration of flexural strength, which derived from internal residual pores behaving as crack-initiation points. These internal pores from agglomerate and low relative density could cause the formation of submicronsize notch. In this study, the transformation of c-BN to h-BN utilized as a new in situ route for introducing h-BN into B4C ceramics, not only makes h-BN uniformly dispersed in B4C, but also effectively increases the relative density of B4C/h-BN composite ceramics comparing pure B4C. Therefore, the addition of homogeneously dispersed h-BN into B4C ceramic (shown in Fig. 5) in this report enables to increase the flexural strength, which has also been demonstrated by previous works [20,33]. According to the data in Table 1, the optimal mechanical properties were achieved with the addition of 10 vol.% c-BN, with the flexural strength and fracture toughness reaching approximately 494 MPa and 4.2 MPa m1/2, respectively. Thus, the in situ synthesized h-BN significantly improved the mechanical properties of the B4C ceramics. However, considering the hardness, flexural strength, and fracture toughness of the ceramic simultaneously, the sample with 5 vol.% c-BN as the additive had the optimal integrated mechanical properties with a hardness of 30.5 GPa, a bending strength of 470 MPa, and a fracture toughness of 3.84 MPa⋅ m1/2. Fig. 4a shows a log plot of the fracture strength against the indentation load for the in situ B4C/h-BN ceramics fabricated with 5 vol.% c-BN. The intersection point, logP*, divides the plot into two regions. In the left region, where the samples have a low indentation load, the fracture strength is mainly controlled by the microstructure and is nearly the same that of the samples with no indentation load. In the right region, the fracture strength of the samples with high indention loads is dominated by external flaws. According to the indentation fracture mechanics, the relation of the fracture strength (σf) and the indentation load (P) can be described as following formula [34]:
with varying amounts of added c-BN after sintering at 1800 °C, as shown in Fig. 2b. Only the B4C phase was detected in the pure B4C ceramics. However, for the B4C-based ceramics with 2, 5, 7, and 10 vol. % c-BN, both B4C and h-BN phases were detected in the sintered specimens. This finding indicates that the c-BN phase in the composites is completely transformed into the h-BN phase. Fig. 3a shows not only the relative densities of the monolithic B4C ceramics and the B4C-based ceramics with 5 vol.% c-BN sintered at 1400–1800 °C, but also the relative densities of the B4C-based ceramics with 5 vol.% h-BN sintered at 1800 °C (blue triangle). The relative densities of the B4C/h-BN ceramics (93.1 %) with direct introduction of 5 vol.% h-BN is lower than that of both monolithic B4C and B4C/h-BN ceramic fabricated with 5 vol.% c-BN sintered at 1800 °C, which testifies that the direct introduction of h-BN does not contribute to the densification of B4C. Furthermore, the relative densities of the composites with 5 vol.% c-BN are obviously higher than those of the pure B4C samples when sintered at 1500–1800 °C. The relative densities of the B4C ceramics with 5 vol.% c-BN prepared at 1500, 1600, 1700, and 1800 °C are 70.5 %, 92.3 %, 97.2 %, and 98.8 % respectively, whereas those of the pure B4C ceramics are 64.0 %, 76.2 %, 87.3 %, and 94.2 %, respectively. However, the relative densities of the composite with 5 vol.% c-BN and the pure B4C ceramic are similar when sintered at 1400 °C, which is consistent with the XRD results (Fig. 2a), which revealed that the transformation of c-BN to h-BN only started at 1500 °C and was completed by 1800 °C. Thus, these findings indicate that the transformation of c-BN to h-BN is beneficial for the densification of B4C. Fig. 3b–e show typical SEM images of the fracture surface of pure B4C and the in situ B4C/h-BN composites with 5 vol.% c-BN sintered at 1600 and 1800 °C. As shown in the SEM images, most of the pores are eliminated from the B4C/h-BN composites, as there are significantly fewer pores present in the fracture surfaces of the B4C/h-BN composites than in those of the pure B4C ceramics sintered at same temperature, confirming the changes in relative densities observed in Fig. 3a. When the sintering temperature reached 1800 °C, in situ bulk B4C/h-BN with 5 vol.% c-BN nearly achieved full densification, whereas pure B4C does not. The results in Fig. 3a are also confirmed by the sintering temperature profile and the shrinkage curves of B4C and the B4C composites with 5 vol.% c-BN in Fig. 3f. Greater shrinkage is observed for the B4C composites with 5 vol.% c-BN than for B4C between 1400 and 1800 °C. The observation of complete densification for the B4C composites with 5 vol.% c-BN at 1800 °C but not for B4C indicates that the c-BN to h-BN phase transformation contributed to the densification of B4C. To eliminate the effect of silica on the densification of these composites, ICP is employed to measure the Si impurity contents of the original B4C powders and B4C powders dealt with agate ball. The results of original one and the dealt one are 0.02 wt.% and 0.04 wt.%, respectively, which indicated nearly no difference in Si content after milling with agate ball and no effect of silica on the densification of these composites. According to the previous report [29], phase transformation of c-BN to hBN is a way of decreasing the total free energy, which indicating the phase transformation is an exothermic process. Therefore the attribution of the phase transformation to the densification of B4C may originate from released energy. The more detailed mechanism may exist and needs to be studied by more future research. Table 1 shows the density, elastic modulus, Vickers hardness, fracture toughness, and flexural strength of the B4C/h-BN ceramics
σf ∝P−β
(4)
β is equal to 1/3 for the material without resistance behavior, while the
Table 1 Density and mechanical properties of the B4C-based ceramics fabricated with different c-BN contents (vol.%) at 1800 °C. Sample B4C B4C-2 vol.% c-BN B4C-5 vol.% c-BN B4C-7 vol.% c-BN B4C-10 vol.% c-BN
Theoretical density (g/cm3)
Relative density (%)
Elastic modulus (GPa)
Hardness (GPa)
Flexural strength (MPa)
Fracture toughness (MPa·m1/2)
2.520 2.513 2.503 2.496 2.487
94.30 99.72 98.88 98.64 98.35
427 ± 16 417 ± 4 401 ± 7 388 ± 9 326 ± 13
35.6 ± 0.6 33.2 ± 0.3 30.5 ± 0.7 29.1 ± 0.5 25.6 ± 0.4
354 ± 34 467 ± 25 470 ± 20 488 ± 17 494 ± 10
2.42 ± 0.70 3.13 ± 0.46 3.84 ± 0.55 4.10 ± 0.34 4.20 ± 0.27
4
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
Fig. 4. (a) Effect of the indentation load on the fracture strength of pure B4C and the in situ B4C/h-BN composites with 5 vol.% c-BN. (b) Toughness curve family diagram for the in situ B4C/h-BN composites with 5 vol.% c-BN. The R-curve behaviour is revealed by the envelope of the tangency points. (c) Load-displacement curves for five indentation tests on in situ B4C/h-BN composites with 5 vol.% c-BN. (d) Typical load-displacement curves for indentation tests on in situ B4C/h-BN composites with 0, 2, 5, 7, and 10 vol.% c-BN.
β is samller than 1/3 for the material with resistance behavior. And the relation between the fracture strength (σf) and the indentation load (P) in the high indentation load region (Fig. 4a) can be expressed using the following formula:
logσf = 2.688 − 0.192 logP
ceramics. In Fig. 5, the dark grey phase corresponds to B4C, whereas the bright flakes correspond to h-BN. Thus, these images also reveal that the h-BN phase is homogeneously dispersed in the B4C matrix without severe aggregation of the h-BN flakes. The fracture surface of pure B4C is very flat, indicating a typical transgranular fracture (Fig. 5a). In contrast, the fracture surfaces of the B4C/h-BN composites are not as flat as that of pure B4C, which could be attributed to the effect of h-BN on the crack propagation path. Generally, microcracks propagate along the weak interface between h-BN and B4C grains, and deflect when meeting h-BN flakes and hard B4C. Thus, the h-BN flakes formed through the phase transformation of c-BN are essential for enhancing the mechanical properties of B4C-based ceramics. A TEM analysis was performed on the interface between B4C and the h-BN in the B4C/h-BN composites, as shown in Fig. 6. The EDS analysis (Fig. 6a, inset) indicates that particles B and A are h-BN and B4C, respectively. The HRTEM image of the boundary in Fig. 6b shows a clean and narrow grain boundary between the B4C and h-BN grains, suggesting strong bonding between the B4C and h-BN grains. Because the boundary strength between the matrix and the second phase significantly affects the mechanical properties of a composite, the B4C/hBN composites achieve enhanced mechanical properties. To reveal the grain boundaries and morphologies of the B4C-based ceramics, electrolytic etching was performed on the polished surface. Representative SEM images of the etched surfaces of the pure B4C ceramic and the B4C/h-BN composite with 2 vol.% c-BN are shown in Fig. 7. The B4C-based composite with added c-BN (Fig. 7b) showed a finer and less faceted grain structure than the pure B4C ceramic (Fig. 7a), with mean grain sizes of approximately 3.2 and 2.2 μm, respectively. When raw c-BN is added to the B4C matrix, the in situ synthesized h-BN grains are located at B4C boundaries and efficiently
(5)
The fit of the data in Fig. 4a to this formula gives a slope of -0.192, which is larger than -1/3, indicating that the in situ B4C/h-BN ceramics fabricated with 5 vol.% c-BN exhibit rising R-curve behaviour. The calculated families of KR(c) curves and R-curves for the in situ B4C/h-BN ceramics are shown in Fig. 4b. The envelope of the tangency points reveals rising R-curve behaviour, which suggests that the fracture toughness will increase with an increase in the crack length. In addition, Fig. 4c shows the load and displacement data recorded during five nano-indentation tests, and good repeatability is observed. Typical load-displacement curves for the B4C-based ceramics, obtained using a depth-sensing nano-indentation tester with a peak load of 2 N, are shown in Fig. 4d. Based on the Oliver–Pharr method [26], the nanoindentation tests were used to determine the elastic modulus of each B4C-based ceramic (Table 1). A high elastic modulus of approximately 400 GPa and degradation were observed at higher in situ h-BN contents. Furthermore, higher permanent deformation and higher elastoplasticity were also observed at higher in situ h-BN contents, which is consistent with the hardness results in Table 1 and reveals that the B4C/h-BN ceramics exhibit increased elastoplasticity. The R-curve behaviour and the load-displacement curves both confirm that in situ synthesized h-BN enhances the mechanical properties of the B4C ceramics. As shown by the SEM images of the fracture surfaces of the pure B4C and B4C/h-BN ceramics (Fig. 5), pores were hardly observed in the B4C/ h-BN composites, whereas some pores were observed in the pure B4C 5
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
Fig. 5. SEM images of the fracture surfaces of B4C/h-BN ceramics fabricated with (a) 0, (b) 2, (c) 5, (d) 7, and (e) 10 vol.% c-BN.
inhibit the growth of B4C grains. Generally, finer grains will significantly improve the mechanical properties of ceramics. Fig. 8 shows various microstructural evidence of the toughening by h-BN platelets. Homogeneously dispersed h-BN platelets absorb the crack propagation energy via different toughening mechanisms such as
pulling out of h-BN, crack bridging, crack deflection, and crack branching. These mechanisms are similar for all filler-reinforced ceramics, but differ in their frequency and effectiveness of load transfer between the matrix and the filler. The h-BN platelets are insert between the matrix grains because of the stress applied by neighbouring B4C
Fig. 6. TEM images of the morphology and grain boundary structure between B4C and h-BN. Inset: EDS analysis identifying particles A and B as B4C and h-BN, respectively. 6
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
Fig. 7. SEM images of the etched surfaces of (a) pure B4C and (b) B4C/h-BN with 2 vol.% c-BN.
Such an interaction decreases the stress intensity at the crack top and even at the temporary crack top. Thus, it is believed that such a crack deflection mechanism can help to increase the toughness of B4C-based ceramics by creating a curving path. Another toughness mechanism, crack branching, can disperse and decrease the stress at the crack top, as observed in Fig. 8b. This mechanism can also significantly contribute to improving fracture toughness of B4C-based ceramics.
grains. Moreover, in situ synthesized h-BN platelets have a high 2D aspect ratio and are homogeneously dispersed in the B4C matrix. Thus, the contact area between h-BN and the B4C matrix is large and is expected to be greater than that between fibres and a matrix. Thus, the energy required to pull out an h-BN platelet is expected to be greater than that required to pull out a fibre or h-BN nanotube. The pulling out of h-BN platelets is believed to be an effective toughening mechanism for the composites, as shown in Fig. 8a, owing to the friction resistance at the interface between the h-BN platelets and the B4C matrix grains, which significantly dissipates the crack propagation energy. Additionally, h-BN platelets could also behave as bridging ligaments in B4C-based ceramics. As h-BN is too rigid to be broken by internal cracks in B4C-based ceramics, h-BN can bridge the cracks and absorb the crack propagation energy, as shown in Fig. 8c. Crack deflection and a change in the direction of crack propagation are shown in Fig. 8d when the crack meets h-BN platelets. The deflection length and deflection angle range from 0.5 μm to several microns and from 20° to 90°, respectively.
4. Conclusion In this study, B4C/h-BN composites prepared using the phase transformation of c-BN to h-BN were characterized. The SEM results and the density analysis revealed that the densification behaviour of B4Cbased ceramics was improved by the energy released by the transformation of c-BN to h-BN, and almost completely densified B4C/h-BN composites were fabricated by SPS. The results showed that in situ synthesized h-BN can be homogeneously dispersed in the B4C matrix
Fig. 8. Toughening mechanisms for B4C/h-BN composites: (a) pulling out of h-BN platelets, (b) crack branding, (c) crack bridging, and (d) crack deflection. 7
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
J. Sun, et al.
microstructure. Moreover, finer B4C grains were obtained in the composite owing to the inhibition of B4C grain growth by in situ synthesized of h-BN at the grain boundaries. The flexural strength and fracture toughness of the B4C ceramics reinforced by in situ synthesized h-BN were significantly higher than those of pure B4C. The sample prepared with 5 vol.% c-BN as an additive exhibited excellent integrated mechanical properties with a hardness of 30.5 GPa, a bending strength of 470 MPa, and a fracture toughness of 3.84 MPa⋅ m1/2. Increasing the cBN content further decreased the hardness but did not significantly improve the bending strength and fracture toughness. Various toughening mechanisms, such as pulling out of h-BN, crack deflection, crack branching, and crack bridging, were responsible for the enhanced mechanical properties of the B4C/h-BN composites. This study showed that the in situ h-BN from c-BN was effective to reinforce B4C ceramics as a structural analogue of GPLs. The exothermal phase transformation could be a new method to improve the densification process of the ceramics and simultaneously offer reinforcement material for ceramic matrix.
Mater. 17 (1999) 27–32, https://doi.org/10.1016/S0263-4368(98)00052-3. [12] H. Zhao, K. Hiragushi, Y. Mizota, Densification and mechanical properties of mullite/TiO2-coated B4C composites, J. Eur. Ceram. Soc. 23 (2003) 1485–1490, https://doi.org/10.1016/S0955-2219(02)00307-2. [13] W.-M. Guo, L.-X. Wu, Y. You, H.-T. Lin, G.-J. Zhang, Three-step reactive hot pressing of B4C–ZrB2 ceramics, J. Eur. Ceram. Soc. 36 (2016) 951–957, https://doi. org/10.1016/j.jeurceramsoc.2015.11.022. [14] L. Levin, N. Frage, M.P. Dariel, A novel approach for the preparation of B4C-based cermets, Int. J. Refract. Met. Hard Mater. 18 (2000) 131–135, https://doi.org/10. 1016/S0263-4368(00)00012-3. [15] M.W. Chen, J.W. McCauley, J.C. LaSalvia, K.J. Hemker, Microstructural characterization of commercial hot-pressed boron carbide ceramics, J. Am. Ceram. Soc. 88 (2005) 1935–1942, https://doi.org/10.1111/j.1551-2916.2005.00346.x. [16] A. Kovalčíková, R. Sedlák, P. Rutkowski, J. Dusza, Mechanical properties of boron carbide+graphene platelet composites, Ceram. Int. 42 (2016) 2094–2098, https:// doi.org/10.1016/j.ceramint.2015.09.139. [17] R. Sedlák, A. Kovalčíková, E. Múdra, P. Rutkowski, A. Dubiel, V. Girman, R. Bystrický, J. Dusza, Boron carbide/graphene platelet ceramics with improved fracture toughness and electrical conductivity, J. Eur. Ceram. Soc. 37 (2017) 3773–3780, https://doi.org/10.1016/j.jeurceramsoc.2017.04.061. [18] M. Wojtoniszak, X. Chen, R.J. Kalenczuk, A. Wajda, J. Łapczuk, M. Kurzewski, M. Drozdzik, P.K. Chu, E. Borowiak-Palen, Synthesis, dispersion, and cytocompatibility of graphene oxide and reduced graphene oxide, Colloids Surf., B 89 (2012) 79–85, https://doi.org/10.1016/j.colsurfb.2011.08.026. [19] M. Kitiwan, A. Iton, T. Goto, Spark plasma sintering of TiN–TiB2–hBN composites and their properties, Ceram. Int. 41 (2015) 4498–4503, https://doi.org/10.1016/j. ceramint.2014.11.144. [20] B. Lee, D. Lee, J.H. Lee, H.J. Ryu, S.H. Hong, Enhancement of toughness and wear resistance in boron nitride nanoplatelet (BNNP) reinforced Si3N4 nanocomposites, Sci. Rep. 6 (2016) 27609, https://doi.org/10.1038/srep27609. [21] D. Lee, S.H. Song, J. Hwang, S.H. Jin, K.H. Park, B.H. Kim, S.H. Hong, S. Jeon, Enhanced mechanical properties of epoxy nanocomposites by mixing noncovalently functionalized boron nitride nanoflakes, Small 9 (2013) 2602–2610, https://doi. org/10.1002/smll.201203214. [22] W.S. Cho, Z.H. Piao, K.J. Lee, Y.C. Yoo, J.H. Lee, M.W. Cho, Y.C. Hong, K. Park, W.S. Hwang, Microstructure and mechanical properties of AlN–hBN based machinable ceramics prepared by pressureless sintering, J. Eur. Ceram. Soc. 27 (2007) 1425–1430, https://doi.org/10.1016/j.jeurceramsoc.2006.05.044. [23] X. Li, Y. Gao, W. Pan, X. Wang, L. Song, Z. Zhong, S. Wu, Fabrication and characterization of B4C-based ceramic composites with different mass fractions of hexagonal boron nitride, Ceram. Int. 41 (2015) 27–36, https://doi.org/10.1016/j. ceramint.2014.07.042. [24] B. Zhong, G.L. Zhao, X.X. Huang, L. Xia, X.H. Tang, S.C. Zhang, G.W. Wen, Microstructure and mechanical properties of ZTA/BN machinable ceramics fabricated by reactive hot pressing, J. Eur. Ceram. Soc. 35 (2015) 641–649, https://doi. org/10.1016/j.jeurceramsoc.2014.09.002. [25] J.C. Garrett, I. Sigalas, A.K. Wolfrum, M. Herrmann, Effect of cubic boron nitride grain size in the reinforcing of α-Sialon ceramics sintered via SPS, J. Eur. Ceram. Soc. 35 (2015) 451–462, https://doi.org/10.1016/j.jeurceramsoc.2014.09.036. [26] F. Khodabakhshi, A.P. Gerlich, D. Verma, M. Haghshenas, Nano-indentation behavior of layered ultra-fine grained AA8006 aluminum alloy and AA8006-B4C nanostructured nanocomposite produced by accumulative fold forging process, Mater. Sci. Eng., A 744 (2019) 120–136, https://doi.org/10.1016/j.msea.2018.12.013. [27] S. Li, C. Wei, W. Wang, P. Wang, Y. Liu, S. Li, G. Wen, Fracture toughness and Rcurve behavior of laminated ZrB2–SiC/SiCw ceramic, J. Alloys. Compd. 784 (2019) 96–101, https://doi.org/10.1016/j.jallcom.2018.12.304. [28] T. Jiang, Z. Jin, J. Yang, G. Qiao, Mechanical property and R-curve behavior of the B4C/BN ceramics composites, Mater. Sci. Eng., A 494 (2008) 203–216, https://doi. org/10.1016/j.msea.2008.04.047. [29] I.A. Petrusha, Features of a cBN-to-graphite-like BN phase transformation under pressure, Diamond Relat. Mater. 9 (2000) 1487–1493, https://doi.org/10.1016/ S0925-9635(00)00273-9. [30] Y. Tan, H. Zhang, S. Peng, Electrically conductive graphene nanoplatelet/boron carbide composites with high hardness and toughness, Scr. Mater. 114 (2016) 98–102, https://doi.org/10.1016/j.scriptamat.2015.12.008. [31] A. Kovalčíková, J. Balko, C. Balázsi, P. Hvizdoš, J. Dusza, Influence of hBN content on mechanical and tribological properties of Si3N4/BN ceramic composites, J. Eur. Ceram. Soc. 34 (2014) 3319–3328, https://doi.org/10.1016/j.jeurceramsoc.2014. 02.021. [32] Y.K. Seo, Y.W. Kim, K.J. Kim, W.S. Seo, Electrically conductive SiC-BN composites, J. Eur. Ceram. Soc. 36 (2016) 3879–3887, https://doi.org/10.1016/j.jeurceramsoc. 2016.06.040. [33] T. Jiang, Z. Jin, J. Yang, G. Qiao, Investigation on the preparation and machinability of the B4C/BN nanocomposites by hot-pressing process, J. Mater. Process. Tech. 209 (2009) 561–571, https://doi.org/10.1016/j.jmatprotec.2008.02.026. [34] L.M. Braun, S.J. Bennison, B.R. Lawn, Objective evaluation of short-crack toughness curves using indentation flaws: case study on alumina-based ceramics, J. Am. Ceram. Soc. 75 (1992) 3049–3057, https://doi.org/10.1111/j.1151-2916.1992. tb04385.x.
Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements This work was financially supported by the National Natural Science Foundation of China (51502220, 51521001, 51672197), the Ministry of Science and Technology of the People’s Republic of China (2015DFR50650), the Self-determined and Innovative Research Funds of WUT (2017II17XZ), and the Open Project Program of the Key Laboratory of Inorganic Functional Materials and Devices, Chinese Academy of Sciences (Grant No. KLIFMD201606). References [1] B. Niu, F. Zhang, J. Zhang, W. Ji, W. Wang, Z. Fu, Ultra-fast densification of boron carbide by flash spark plasma sintering, Scr. Mater. 116 (2016) 127–130, https:// doi.org/10.1016/j.scriptamat.2016.02.012. [2] K.H. Kim, J.H. Chae, J.S. Park, J.P. Ahn, K.B. Shim, Sintering behavior and mechanical properties of B4C ceramics fabricated by spark plasma sintering, J. Ceram. Process. Res. 10 (2009) 716–720. [3] 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, https://doi.org/ 10.1179/095066009X12506721665211. [4] V. Domnich, S. Reynaud, R.A. Haber, M. Chhowalla, Boron carbide: structure, properties, and stability under stress, J. Am. Ceram. Soc. 94 (2011) 3605–3628, https://doi.org/10.1111/j.1551-2916.2011.04865.x. [5] T.K. Roy, C. Subramanian, A.K. Suri, Pressureless sintering of boron carbide, Ceram. Int. 32 (2006) 227–233, https://doi.org/10.1016/j.ceramint.2005.02.008. [6] Y. Tan, H. Luo, H. Zhang, S. Peng, Graphene nanoplatelet reinforced boron carbide composites with high electrical and thermal conductivity, J. Eur. Ceram. Soc. 36 (2016) 2679–2687, https://doi.org/10.1016/j.jeurceramsoc.2016.04.036. [7] W. Ji, S.S. Rehman, W. Wang, H. Wang, Y. Wang, J. Zhang, F. Zhang, Z. Fu, Sintering boron carbide ceramics without grain growth by plastic deformation as the dominant densification mechanism, Sci. Rep. 5 (2015) 15827, https://doi.org/ 10.1038/srep15827. [8] M. Mashhadi, E. Taheri-Nassaj, V.M. Sglavo, H. Sarpoolaky, N. Ehsani, Effect of Al addition on pressureless sintering of B4C, Ceram. Int. 35 (2009) 831–837, https:// doi.org/10.1016/j.ceramint.2008.03.003. [9] F. Ye, Z. Hou, H. Zhang, L. Liu, Densification and mechanical properties of spark plasma sintered B4C with Si as a sintering aid, J. Am. Ceram. Soc. 93 (2010) 2956–2959, https://doi.org/10.1111/j.1551-2916.2010.03931.x. [10] V.B. Fedorus, G.N. Makarenko, S.P. Gordienko, É.V. Marek, I.I. Timofeeva, Interaction between boron carbide and chromium oxide, Powder Metall. Met. Ceram. 34 (1996) 637–639, https://doi.org/10.1007/BF00559492. [11] P. Rogl, H. Bittermann, Ternary metal boron carbides, Int. J. Refract. Met. Hard
8