BN(C) composite ceramic prepared from the mechanically alloyed SiBCN powder

Journal of Alloys and Compounds 537 (2012) 346–356

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Microstructural features and properties of the nano-crystalline SiC/BN(C) composite ceramic prepared from the mechanically alloyed SiBCN powder Pengfei Zhang, Dechang Jia ⇑, Zhihua Yang, Xiaoming Duan, Yu Zhou Institute for Advanced Ceramics, Harbin Institute of Technology, Harbin 150080, China

a r t i c l e

i n f o

Article history: Received 3 April 2012 Received in revised form 9 May 2012 Accepted 12 May 2012 Available online 26 May 2012 Keywords: Mechanical alloying Nano ceramic Microstructure SiBCN SiC/BN(C)

a b s t r a c t Nano SiC/BN(C) ceramic was prepared from the mechanically-alloyed amorphous SiBCN powder, hot pressed at 1900 °C under 80 MPa in the nitrogen atmosphere for 30 min. The microstructures and the properties of the prepared ceramic were carefully studied by XRD, TEM (SAED, HRTEM and EFTEM), EELS and property testing. Results show that the ceramic consists of b-SiC, a-SiC and BN(C). SiC has an average grain size of about 78.2 ± 32.4 nm, most of which have numerous stacking faults. BN(C) has small size, no fixed shape and turbostratic structure with heterogeneously distributed t-carbon layers, t-BN layers and B-doped t-carbon layers. It has uniform distribution in the ceramic, retarding the atomic diffusion and being responsible for the fine grains. The ceramic has room-temperature density, flexural strength, Young’s modulus and fracture toughness of 2.6 g/cm3, 331.1 MPa, 139.4 GPa and 2.8 MPa m1/2, respectively. The detailed investigation of the hot-pressed SiC/BN(C) ceramic is helpful for the further study of this material and other ceramics, targeted for the high-temperature applications. Ó 2012 Elsevier B.V. All rights reserved.

1. Introduction Organic polymer derived Si–B–C–N ceramics are well known for their outstanding features. For example, they have amorphous structures at temperatures up to 1800 °C [1,2]; they possess even better oxidation resistance than SiC or Si3N4 ceramics at 1700 °C in air [1–3]; and their creep rates show time-dependent decreasing during the testing at 1500 °C under 50 MPa [4,5]. Currently, Si–B– C–N ceramics attract increasing attention in both theoretical study and engineering applications [6,7]. It is accepted that the polymerderived amorphous Si–B–C–N ceramics generally consist of two phases, that is, amorphous SiCxN4 x (x = 1–4) and graphite-like BN(C), 1–2 nm in size [8,9]. BN(C) actually comprises of t-BN and t-carbon layers [8–10]. When annealed at temperatures higher than that for pyrolysis, SiCxN4 x may rapidly decompose into SiC and Si3N4 in less than 15 min, showing a thermally-activated crystallization process [11]. However, the decomposition and the following grain growth process may be greatly retarded by the turbostratic BN(C) [12,13]. The specific structure is exactly the reason why SiBCN ceramic is still X-ray amorphous at 1800 °C or higher temperatures, and why their grain sizes are still several hundred nanometers or so even the ceramic is annealed at 1800 °C for 10 h [12]. However, disadvantages also exist in the polymer pyrolysing method. For example, polymer pyrolysis generally produce about ⇑ Corresponding author. Address: Institute for Advanced Ceramics, Science Park, Harbin Institute of Technology, P.O. Box 3022, No. 2, Yikuang Street, Harbin 150080, PR China. Tel.: +86 451 86418792; fax: +86 451 86414291. E-mail address: [email protected] (D. Jia). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.05.073

30% volumetric shrinkage and 20–30 wt.% weight loss, as a result, pores or even cracks may appear in the products [12,13]. In addition, some organic solvents are very harmful and may cause severe pollution [12–14]. Thus, it is necessary to explore alternative ways for the Si–B–C–N ceramic preparation. Research has confirmed that mechanical alloying (MA) is an appropriate method for the preparation of various non-equilibrium materials, such as amorphous body, quasicrystal or supersaturated solid solution [15–17]. In our previous research [18–20], high-energy shaker ball milling and hot pressing were adopted to prepare SiBCN powder and corresponding nano ceramics. The as-milled powder was almost amorphous. Sintered at 1900 °C under 40 MPa in the nitrogen atmosphere for 30 min, the prepared ceramic has grain sizes of about 200–500 nm and a relative density of about 99.4%. It appears that the mechanically alloyed SiBCN powder is easy to be densified. However, due to the brittleness and the wearness of the zirconia vials and balls, a little ZrO2 impurity was often introduced to the powder during the milling process. Since zirconia has a low sintering temperature, it is inferred to act as sintering aid, which probably affects the atomic diffusion, the densification process, and hence the ceramic microstructures and properties. If the MA process is well controlled, it is most likely to prepare the completely amorphous SiBCN powder free from impurity contamination, as well as the nano ceramics with finer grains. Additionally, the microstructures study of the hot-pressed nano ceramic has not been carefully conducted yet. Hence, the present work strives to find better understandings on the nano SiC/BN(C) ceramic.

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In the current work, to reduce the impurity as much as possible, the high-quality silicon nitride balls and vials, together with a planetary ball mill, are employed. The as-milled SiBCN powder is effectively amorphous and free from any crystalline impurity. Furthermore, the nano ceramic with finer grains is also successfully prepared. As revealed by the previous research [20,21], the variation of the chemical composition may affect the crystallinity and the content of each phase in the prepared ceramic, but the phase constitution is normally stable within certain composition limits. Namely, the nano ceramics prepared by the current method are generally composed of b-SiC, a-SiC and BN(C). To investigate the microstructures of each phase, it is advisable to carry out research on the ceramic with typical structures and favorable properties, thus the chemical composition in the current research is fixed as Si:C:BN = 2:3:1 in molar ratio. The microstructures and the roomtemperature properties of the prepared ceramic are carefully studied.

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transmission electron microscope (EFTEM) and electron energy loss spectrum (EELS) were also adopted to study the atomic arrangement, the phase distribution and the chemical bonds, respectively. The property measurement was carried out on omnipotence mechanics tester (Instron 5569, Instron Corp., USA), hardness tester (HVS-5, Laizhou Huayin Testing Instrument Co., Ltd., China) and thermal expansion device (DIL402C, Netzsch Company, Germany). The density of the ceramic was determined by mass/volume. The flexural strength and the Young’s modulus were obtained using three-point bending test on 3  4  26 mm bars with a span of 20 mm and a crosshead speed of 0.5 mm/min. The fracture toughness was evaluated by the single-edge notched beam test on 2  4  20 mm bars with a span of 16 mm and a crosshead speed of 0.05 mm/min. The Vickers hardness was measured on polished sample surface with a load of 10 kg and a holding time of 10 s. The thermal expansion coefficient was tested on 3  4  10 mm bar, which was heated to 1400 °C with a heating rate of 10 °C/min.

2. Experimental procedures 3. Results and discussion 2.1. Sample preparation 3.1. Structural features of the mechanically alloyed SiBCN powder The used raw materials were well crystalline cubic silicon (45.0 lm, 99.5% in purity, Beijing MounTain Technical Development Center, China), hexagonal boron nitride (0.6 lm, 98.0% in purity, Advanced Technology & Materials Co. Ltd., Beijing, China) and graphite powders (8.7 lm, 99.5% in purity, QingDao HuaTai Lubricant Sealing S&T Co., Ltd., China). For a more uniform distribution of different elements, finer grains in the prepared ceramics, and the comparison with the polymer-derived ceramics, silicon and graphite were used as starting powders in the current research instead of SiC. The composition was designed as Si:C:BN = 2:3:1 in molar ratio, a choice dependant on the literature, the current preparation technology and our previous research. The ball-to-powder mass ratio was set as 20:1. The powder mixture was loaded into silicon nitride vials along with silicon nitride balls under the argon atmosphere. Sealed with rubber rings, two vials were fixed on a planetary ball mill (Fritsch P4, Germany). The rotation speed of the main disk was set as 350 rpm and the vials 600 rpm in reverse. The machine rested for 30 min every 1 h, and the total milling time was 40 h. The milling time depends on the results of the optimization experiment on the milling parameters, which will be discussed in another research paper. A portion of the as-milled powder was removed for X-ray diffraction and TEM analysis, and the rest was hot pressed in HIGH MULTI 10000 furnace. The powder was heated to 1900 °C under a heating rate of 20 °C/min, and then kept at the target temperature and pressure in the nitrogen atmosphere (1 bar) for 30 min. The loading process of the axial pressure starts at 1200 °C and finishes at 1400 °C. In the cooling stage, the pressure was slowly unloaded in 5 min. 2.2. Structure analysis and performance evaluation The particle size distribution and the chemical composition of the as-milled powder was analyzed by the laser scattering particle-size analyzer (LA-920, Horiba Comp., Japan) and the chemical titration method, respectively. The prepared ceramic was cut, grinded and/ or ion-beam thinned for the property testing and the structure investigation. For the phase identification, X-ray diffractometer (40 KV/100 mA, D/max-cB CuKa, Rigaku Corporation, Japan) was used to obtain the X-ray diffraction (XRD) spectrum at 2h = 10o– 90o with a scanning speed of 4o/min. Transmission electron microscope (TEM, Tecnai F30, 300 KV, FEI Company, USA) was used to study the ceramic microstructures. In addition, high-resolution transmission electron microscope (HRTEM), energy filtering

After being milled for 40 h, the composite powder is completely X-ray amorphous, as shown in Fig. 1(a). This is attributed to the repeated physical and chemical processes during the high-energy ball milling, including fierce collisions, rupture, local high temperature, fusion and short-range diffusion [15–17]. The amorphous structure is further corroborated by the selected-area electron diffraction (SAED) pattern, as displayed in the inset in Fig. 1(b), where a large diffuse scattering spot appears. As shown in the bright field image in Fig. 1(b), the powder has particle sizes of about 50–110 nm. This is in conformity with the statistical 116 ± 34 nm, revealed by the morphology in Fig. 1(c). However, it is known that the nano particles, formed during the high-energy ball milling, generally have large surface energy and tend to get together for lowering the total energy. As a result, numerous near-spherical agglomerates appear in the powder, which can be seen in Fig. 1(d). The particle size distribution of the as-milled SiBCN powder is shown in Fig. 2. It is found that about 90% of the particles are smaller than 14 lm, and the average particle size is calculated as 6.5 ± 5.4 lm. Actually, after the amorphization of the mixed powders, the as-milled powders always consist of near-spherical agglomerates, in spite of the varied milling time. This is the general case using mechanical alloying to prepare ceramic powders [15]. The influences of the milling parameters, including rotation speed, ball-to-powder mass ratio and milling time, on the microstructures and the morphologies of the asmilled SiBCN powder will be discussed in another research paper. The chemical composition of the as-milled powder is analyzed by the chemical titration method and the results are listed in Table 1. It is found that except for Si, B, C and N, a small proportion of oxygen is also included in the powder. The oxygen may be introduced by the raw materials, the impurities in glove box or the exposure to air before analyzing. The element molar ratio is calculated as Si:B:C:N:O  2.00:0.99:2.74:0.81:0.24, giving a stoichiometric composition of Si2.00B0.99C2.74N0.81O0.24. The source of oxygen and its influence on the preparation, the structures and the properties of the prepared ceramic will be carefully studied in the future work. 3.2. Structural features of the prepared SiC/BN(C) ceramic As mentioned above, the phase constitution of the ceramics, prepared by the current method, is generally stable within certain chemical composition limits. In the current research, it is found that the principal processing variables (temperature and pressure)

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Fig. 1. Structural features of the mechanically alloyed SiBCN powder. (a) XRD spectrum; (b) Bright field image and SAED pattern studied by TEM; (c), (d) Powder morphologies investigated by SEM (15 KV).

Fig. 2. Particle size distribution of the as-milled SiBCN powder.

Table 1 Chemical composition of the mechanically alloyed SiBCN powder prepared in the current research. Element

Si

B

C

N

O

Relative content, % atomic ratio

29.5

14.6

40.4

12.0

3.5

have the similar effect. As shown by the XRD spectra in Fig. 3(a), the ceramics sintered at different pressures (80 and 50 MPa) or different temperatures (1900 and 1800 °C) generally consist of the same phases, that is, b-SiC, a-SiC and BN(C). The only difference is the crystallinity and relative content of each phase. On the other hand, the mechanically alloyed SiBCN powder free from zirconia impurity is found hard to densify. Lowering the sintering temper-

ature or the applied pressure may lead to the sharp weakening of the mechanical properties, as discussed in the latter section. Thus, the ceramic prepared at the higher temperature and the higher pressure is expected to possess higher mechanical properties and typical microstructure. Therefore, the microstructural investigation in the present work focuses on one composition (Si:C:BN = 2:3:1) and one sintering condition (1900 °C/80 MPa). This section discusses the general microstructure of the prepared SiC/BN(C) ceramic. As indicated by the XRD spectrum (1) in Fig. 3(a), the amorphous powder crystallizes to form crystalline SiC and BN(C), when hot pressed at 1900 °C under 80 MPa for 30 min. The XRD pattern shows typical peaks of b-SiC, a-SiC and BN(C). In detail, peaks at 35.8°, 41.6°, 60.2°, 72.0° and 75.6° represent the diffraction of bSiC, with the plane indices of (1 1 1), (2 0 0), (2 2 0), (3 1 1) and (2 2 2), respectively. Peaks at 33.8°, 38.3°, 65.8° and 73.6° correspond to the diffraction of hexagonal a-SiC. The broad peak at 26.2° is assigned to BN(C), implying a graphite-like structure and a low crystallinity. It is well known that silicon carbide includes b- and a-SiC, which are stable at low and high temperatures, respectively. At temperatures higher than 2100 °C, b-SiC may transform into a-SiC spontaneously [22,23]. In addition, favorable kinetic conditions, such as high-energy ball milling, high pressure or long-time annealed at high temperature, may also facilitate the formation of a-SiC [24]. Hence, it is reasonable to believe that aSiC exists as transition phase in the prepared ceramic. If the chemical composition is designed as Si:C:BN = 1:1:1 in molar ratio, that is, after the formation of SiC, no ‘‘free’’ carbon is left. The XRD spectrum of the prepared ceramic may generate a sharp peak for h-BN at 26.7° [25], which matches that of (0 0 2) planes. However, in the recorded XRD pattern in Fig. 3(a), a broad peak appears at about 26.2°, showing the shift towards lower angle and implying the expansion of interplanar spacing. Because of their high similarities in crystal

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Fig. 3. Structural features of the prepared SiC/BN(C) ceramic hot pressed at 1900 °C/80 MPa. (a) XRD spectra of the ceramics hot pressed at different temperatures and pressures: (1) 1900 °C/80 MPa, (2) 1900 °C/50 MPa, (3) 1800 °C/80 MPa; (b), (c) Bright field images investigated by TEM; (d) SAED pattern collected from the area in (b).

structure and lattice parameters, the excess graphite may be mixed with h-BN to form a graphite-like structure, whose features require more experimental data to clarify. Fig. 3(b) and (c) show the TEM microstructures of the SiC/BN(C) ceramic prepared at 1900 °C/80 MPa. Nano-sized SiC and BN(C) have relatively uniform distribution in the ceramic, without abnormal growth. The SiC grains are counted and the size distribution is shown in Fig. 4, giving an average grain size of about 78.2 ± 32.4 nm. Additionally, lots of light-and-dark stripes can be observed in many SiC grains, implying the existence of numerous

Fig. 4. SiC grain size distribution in the SiC/BN(C) ceramic prepared at 1900 °C/ 80 MPa.

stacking faults. BN(C) has smaller size, no fixed shape and mainly distributes between SiC grains or in their junction regions. It follows that the nano SiC/BN(C) ceramic with fine grains is successfully prepared in the current work using the improved mechanical alloying and hot pressing techniques. The SAED pattern in Fig. 3(d) records the information within the area in Fig. 3(b). The bright diffraction rings, with the interplanar spacing d = 0.2524, 0.2197, 0.1551, and 0.1323 nm, denote a large quantity of b-SiC grains. Similarly, the bright diffraction ring, with d = 0.3504 nm, reveals a great amount of BN(C). For the existence of a-SiC, there are two pieces of evidence in the SAED pattern. One is the broadened diffraction ring with interplanar spacing d = 0.2524 nm. In fact, this broad ring represents a varied d values from 0.2364 to 0.2654 nm, covering d(0 0 2), d(1 0 0), d(1 0 1), d(1 0 2), d(1 0 3) and d(1 0 4) of hexagonal a-SiC (2H-, 4H-, 6H- and 8H-SiC). This is in accordance with the XRD peaks from 2h = 33.8° (d = 0.2647 nm) to 2h = 38.3° (d = 0.2346 nm). The other evidence is the darker diffraction ring with the interplanar spacing d = 0.1451 nm. This is in agreement with the XRD peak at 2h = 65.8° (d = 0.1419 nm), corresponding to d(1 0 3), d(1 0 6), d(1 0 9) and d(1 0 1 2) of hexagonal a-SiC. Although the microstructures of SiC, BN(C) and their interfacial regions require further clarification, it is now clear that the prepared ceramic is composed of b-SiC, graphite-like BN(C) and a small amount of hexagonal a-SiC, with the grain size of 100 nm or so. The small grain size is closely related to the mechanical alloying process, during which the high rotation speed and the fierce collisions can produce as many as possible metallurgical reactions and mixing powders as uniform as possible. The relatively uniform distribution of different elements may retard the

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atomic diffusion, limit the grain growth and greatly contribute to the fine grains. In addition, no Si3N4 appears in the prepared ceramic. This may be related to the designed chemical composition and the applied processing techniques, requiring more research work to clarify. 3.3. Microstructural features of b- and a-SiC in the prepared SiC/BN(C) ceramic In Fig. 3, the sharp XRD peaks and the bright SAED rings suggest a high crystallinity of SiC. However, there are also various indications for the existence of many plane faults in the SiC grains. Hints include the appearance of small steps at the left side of 2h = 35.8° and 38.3° in the XRD spectrum, the black-and-white stripes in some grains in the bright field images and the broadened SAED ring with the interplanar spacing d = 0.2364–0.2654 nm. The

microstructural features of b- and a-SiC grains are carefully discussed in this section, with the results in Figs. 5 and 6. In Fig. 5(a), the image shows two b-SiC grains with no stripes. Under HRTEM, it is found that the atoms join together in regular arrays, as displayed in Fig. 5(b). The interplanar spacing is measured as d = 0.2528 nm, which is consistent with that of d(1 1 1) in b-SiC. The inserted fast Fourier transform (FFT) pattern indicates that the incident electron beam runs parallel to the [0 1 1] zone axis. Besides the well crystallized grains, b-SiC containing volumes of stacking faults is also common in the ceramic, with an example displayed in Fig. 5(c). The black-and-white stripes show a periodic variation of phase contrast, implying the existence of plane defects in the lattices. Under HRTEM, the defects are found to be densely populated stacking faults with the fault plane of (1 1 1), as revealed in Fig. 5(d). The significantly elongated spots, as suggested by the FFT pattern in Fig. 5(d), also imply the existence of numerous

Fig. 5. Structural features of b-SiC in the SiC/BN(C) ceramic prepared at 1900 °C/80 MPa. (a) Two well crystallized grains; (b) HRTEM image collected from the grain in (a); (c) b-SiC grain with lots of stacking faults; (d) HRTEM image of the grain marked in (c); (e) Inverse FFT image of the region marked by the white box in (d); (f) EDX spectrum acquired from certain b-SiC grain.

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Fig. 6. Structural features of a-SiC in the SiC/BN(C) ceramic prepared at 1900 °C/80 MPa. (a) Bright field image of one a-SiC grain; (b) HRTEM image of the grain marked in (a); (c) FFT pattern of the HRTEM image in (b); (d) Inverse FFT image of the area marked by the white box in (b).

stacking faults. Fig. 5(e) is the inverse FFT image of the area marked by the white box in Fig. 5(d), showing the dislocation between adjacent atomic planes, as marked by the white arrows. The energy dispersive X-ray (EDX) spectrum of one SiC grain is displayed in Fig. 5(f), from which it is found that the grain contains Si, C and a small amount of O. The oxygen contamination may come from the raw powders, the impurity in glove box or the exposure to air before analyzing. While it is convenient to study the microstructure of b-SiC, the investigation on a-SiC is difficult, because of its low content in the prepared ceramic. Fortunately, one a-SiC grain is found after much effort, and its microstructure is recorded in Fig. 6. The bright field image in Fig. 6(a) shows that the black-and-white stripes also appear in the a-SiC grain, again implying the existence of stacking faults. The microstructures of the a-SiC grain and the plane defect are revealed by the HRTEM image in Fig. 6(b) and the FFT pattern in Fig. 6(c). The interplanar spacing is measured as d = 0.5125 nm, approaching to that of d(0 0 2) in hexagonal a-SiC. The incident electron beam runs parallel to the [2–10] zone axis, and the elongated spots are caused by stacking faults. Fig. 6(d) is the inverse FFT image of the area marked by the white box in Fig. 6(b), showing the atomic arrangement in hexagonal a-SiC. It is generally accepted that a-SiC (2H-, 4H-, 6H- and 8H-SiC) possesses hexagonal crystal lattice and has a symmetry of P63mc in space group. In this structure, positive ions are integrated as hexagonal close packed (HCP) system, and negative ions hold the interstitial sites in tetrahedrons. Observing along the c axis of HCP lattice, the atomic stacking sequence is ABAB. . ., which is identical to the results in Fig. 6(d). However, in some cases, plane defects may occur in lattice, such as the stacking

faults arrowed in Fig. 6(d). In the current situation, the local stacking sequence is changed into ABAABABAAB. . ., which can be regarded as a layer missing of B atoms or an intercalation of A atoms. The generation of stacking faults is common in SiC grains, due to the especially low stacking fault energy [26]. 3.4. Microstructural features of BN(C) in the prepared SiC/BN(C) ceramic It is now definite that the currently prepared ceramic is made up of b-SiC, a-SiC and BN(C). For BN(C), various hints imply a graphite-like structure. Hints include the broad XRD peak at 2h = 26.2°, the irregular morphology of the phase located between SiC grains and the SAED ring with d = 0.3504 nm. This section discusses the microstructures of BN(C) based on further evidence. Fig. 7 shows the detailed features of BN(C), covering its size, morphology, distribution and atomic arrangement. The bright field image in Fig. 7(a) indicates that BN(C) has no fixed shape and mainly appears between SiC grains, with a thickness of about 15– 50 nm. Cluster region with larger size also exists, but its size is at most 100 nm or so. Under HRTEM, BN(C) is found to have a turbostratic structure. This structure is characterized by a layer structure similar to graphite or h-BN, but the layers may rotate by a random angle around the c-axis or the direction perpendicular to c-axis [27]. As a result, the interplanar spacings along the c axis are generally enlarged by varied degrees. This is in agreement with the results in Fig. 7(b). The inserted FFT pattern shows a greatly broadened SAED ring with the minimum interplanar spacing d  0.3457 nm, which is slightly larger than d(0 0 2) in graphite or h-BN.

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Fig. 7. Structural features of BN(C) in the SiC/BN(C) ceramic prepared at 1900 °C/80 MPa. (a) Bright field image showing the shape, the size and the distribution of BN(C); (b) HRTEM image and FFT pattern of certain BN(C) region.

To verify the turbostratic structure of BN(C), the XRD spectra at different angle range are precisely measured for the prepared ceramic, using a scanning speed of 1o/min. Fig. 8 shows the partial XRD spectra of BN(C) in the prepared ceramic, as a comparison with the spectra of the raw graphite and h-BN powders. In Fig. 8(a), the (0 0 2) peak of BN(C) has wider full width at half maximum (FWHM) and smaller diffraction angle, as compared with that of raw powders. The (0 0 2) peak shifts towards the lower angle by about 0.72o, resulting in the expansion of d(0 0 2) by 0.0091 nm.

Fig. 8. Partial XRD spectra of the prepared ceramic, together with the spectra of the raw graphite and h-BN powders in the same angle range. The noticeable (hkl) (l – 0) peaks show the three-dimensional ordered structure.

Since the interlayer-spacing fluctuation in graphite or h-BN is generally accompanied by rotation, translation or curvature of graphene or single BN layer [27], the peak shift towards lower angle implies the decrease of the three-dimensional order degree and the formation of turbostratic structure. Fig. 8(b) shows the threedimensional order (1 0 1) peak of h-BN, (1 0 1), (1 0 2), (1 0 3) and (1 0 4) peaks of graphite, and the XRD spectrum of the prepared ceramic at the same angle range. Since graphite and h-BN have three-dimensional ordered structures with regular atomic arrangement, their XRD spectra generally have noticeable (hkl) (l – 0) peaks, as shown in Fig. 8(b). However, in the angle range of 41.5o–47.0o, no identifiable peak appears in the XRD spectrum of the prepared ceramic, indicating that BN(C) in the prepared ceramic has rather low three-dimensional order degree. Most probably, various distortion factors, such as rotation, translation and curvature of single atomic layer occur in the crystal lattice, and this is in agreement with the peak shift and broadening revealed in Fig. 8(a). In other words, the BN and C layers in BN(C) are turbostratic BN (t-BN) and turbostratic carbon (t-carbon), respectively, not h-BN and graphite. To investigate the chemical bonds between B, C, and N atoms in BN(C), and to further verify the above conclusions, EELS is adopted and three typical spectra acquired from BN(C) are listed in Fig. 9. In region (1), only characteristic edges appear for carbon atoms, that is, the absorption edge for electron transition 1s ? p⁄ at 284 eV and 1s ? r⁄ at 289 eV. These are the features of carbon atoms bonded with other carbon atoms in the well sp2 hybridized graphite rings [28]. It is generally accepted that in sp2 hybridized graphite, every carbon atom holds r bond with three other ones in the ring, and possesses p bond with the counterpart in the adjacent atom layer. Based on the molecular orbital theory, antibonding orbits normally exist, which are empty and generally labeled as r⁄ and p⁄. When the incident electron collides with the electron in carbon 1s core level, the latter may absorb characteristic energy and jump into p⁄ or r⁄ orbit. The energy difference between p⁄, r⁄ orbits and 1s level approximately falls on 284 and 289 eV, respectively. This is the reason why the absorption edges appear in the successively falling curve of the energy-loss dependent electron counts. The EELS spectrum in region (1) reveals that only carbon atoms appear in this area, and the atoms form a graphite-like structure. Taking the results in the HRTEM image and the XRD spectra of BN(C) into consideration, it is confirmed that the carbon atoms in region (1) exist as t-carbon. In region (2), the EELS spectrum shows distinct K-edges for B and C atoms, while N K-edge is very weak. For B atoms, the absorption

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atoms are defined. Therefore, BN(C) in region (3) consists of t-BN, t-carbon and B-doped t-carbon layers. It is determined from the acquired results that BN(C) in the prepared ceramic has turbostratic structure, like that in turbostratic carbon or turbostratic BN. The BN(C) is composed of heterogeneously distributed t-carbon layers, t-BN layers and B-doped t-carbon layers. The formation of turbostratic BN(C) in the current research is greatly attributed to the mechanical alloying process and the easy-to-strip character of graphite and h-BN. The fierce collisions during mechanical alloying are able to break h-BN or graphite into tiny fragments, mix them as uniform as possible and produce reactions between different atoms. However, mechanical alloying is a mechanical process after all. For the large-scale formation of compound or the preparation of extremely uniform mixture, it cannot compare with the chemical methods, such as CVD, co-precipitation or organic polymer pyrolysis. 3.5. Features of the interfacial region between BN(C) and SiC grains

Fig. 9. Three typical EELS spectra of BN(C) in the SiC/BN(C) ceramic prepared at 1900 °C/80 MPa, marked as (1), (2) and (3).

edges at 189 eV and 192 eV coincide with the transition of 1s ? p⁄ and 1s ? r⁄ in h-BN [29], respectively. Furthermore, the N K-edge at 401 eV also implies the transition of 1s ? p⁄ in h-BN [30]. These results make us believe that in region (2), both B and N atoms emerge and they form turbostratic BN. For C atoms, the K-edges at 283 and 288 eV suggest the transition of 1s ? p⁄ and 1s ? r⁄ in the graphite-like structure. Hence, in region (2), BN(C) consists of t-BN and t-carbon. The EELS spectrum in region (3) shows evident absorption edges for B, C, and N atoms. The B K-edges at 192 and 199 eV shift towards higher energy by about 2 and 7 eV, respectively, as compared with that in h-BN. The current two K edges are consistent with the electron transition of 1s ? p⁄ and 1s ? r⁄ in boron carbide [31], respectively, indicating that B atoms are bonded with C atoms. The K edge shift towards higher energy may be related to the charge transfer from boron to carbon atoms, affecting the energy of core states and valence states. The K-edges at 284 eV (1s ? p⁄) and 290 eV (1s ? r⁄) imply sp2 hybridized carbon atoms in the graphite-like structure. At 401 eV (1s ? p⁄) and 406 eV (1s ? r⁄), well sp2 hybridized N atoms bonded with B

Since no low-melting-point sintering additives are included, the solid state diffusion may greatly contribute to the densification and the grain growth in the current ceramic. However, as Si–C, B–N and C–C are all strong covalent bonds, their self diffusion coefficients are rather small. In addition, the atomic diffusion may be further hindered by the relatively uniform distribution of turbostratic BN(C) and SiC in the ceramic. Evidence is provided by the results in the HRTEM and the EFTEM images in this section. The HRTEM image in Fig. 10(a) displays the microstructures of the interfacial region between SiC and BN(C). The atomic arrangement in the two phases is quite different. While the atoms are ordered in SiC, they have only short- or at best medium-range order in BN(C). Obviously, the SiC grain is enwrapped by the turbostratic BN(C). Consequently, the SiC grain growth may be suppressed in the encircled direction. Fig. 10(b) and (c) are the inverse FFT images of the marked areas B and C in Fig. 10(a), respectively. The images indicate that the interface region is limited to three to four atomic layers, with the thickness of about 0.78–1.12 nm. No impurity or any other phases exist in this region. The interfacial features, together with the phase composition, structure and distribution, may be responsible for the limited grain size. When the powder is hot pressed, prior crystallizing or ordering may occur among the neighboring atoms. When the adjacent atoms are all integrated into certain grain, further grain growth may mainly depend on the long range atomic diffusion. However, due to the restriction of turbostratic BN(C), the long range diffusion

Fig. 10. Structural features of the interfacial region between SiC and BN(C) in the ceramic prepared at 1900 °C/80 MPa. (a) HRTEM image of the interfacial region; (b), (c) Inverse FFT images of the areas marked by the white rectangles B and C in (a), respectively.

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Fig. 11. EFTEM images of the prepared ceramic. (a), (b), (c), (d) Z-contrast images of Si, B, C, N, respectively; (e) Color map combined from the above four Z-contrast images, showing the phase distribution. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

of Si and C atoms may be difficult. Similarly, it may be also a hard diffusion process for the long range diffusion of B and N atoms. The mutual restriction may greatly limit the grain growth. Besides the results in the bright field images and HRTEM images, the EFTEM pictures in Fig. 11 also provide reliable evidence. In the EFTEM images, the bright region is rich and the dark area is scarce in certain kind of atoms. After comparing the images in Fig. 11(a)–(d), it is found that the dwelling range of B atoms is approximately the same as that of N atoms. C atoms exist not only in silicon-rich regions, but also in areas abundant of B and N atoms. This is in agreement with the fact that the prepared ceramic is composed of SiC and BN(C). The relative distribution of each phase is shown by

the color map in Fig. 11(e), where the mauve areas represent SiC and the green regions stand for BN(C). It is found that the two phases indeed have rather small grain size and alternate distribution. The long range atomic diffusion may be rather difficult in such microstructure and the grain growth is most likely restricted. Besides the information on phase distribution, the EFTEM images also imply the inhomogeneous composition in BN(C). The brighter area (1) and darker area (2) in Fig. 11(c) suggest higher and lower content of C atoms, respectively, though the two areas all belong to the BN(C) region. The condition is opposite for B atoms in Fig. 11(b), where the atoms is rich in area (2) and scarce in area (1). These coincide well with the results from EELS spectra.

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Table 2 Room temperature properties of the SiC/BN(C) ceramics hot pressed at different temperatures and pressures. (Atmosphere: N2, 1 bar; Holding time: 30 min; Heating rate: 25 °C/ min).

a

Sintering conditions

Density (g/cm3)

Flexural strength (MPa)

Young’s modulus (GPa)

Fracture toughness (MPa m1/2)

Vickers’ hardness (GPa)

Coefficient of thermal expansiona (10 6/°C)

1900 °C/80 MPa 1900 °C/50 MPa 1800 °C/80 MPa

2.6 ± 0.1 2.2 ± 0.1 2.1 ± 0.1

331.1 ± 40.5 89.4 ± 9.3 66.3 ± 7.2

139.4 ± 16.0 33.6 ± 3.6 30.6 ± 2.4

2.8 ± 0.9 1.2 ± 0.1 0.8 ± 0.1

2.7 ± 0.4 – –

4.7 ± 0.2 – –

Average value from 50 –1400 °C.

3.6. Structure comparison between the hot pressed SiC/BN(C) and the polymer derived nano Si–B–C–N ceramic The above results and discussion reveal the detailed microstructure of the hot pressed SiC/BN(C) ceramic. To evaluate the value of the current method in the nano Si–B–C–N ceramic preparation, it is necessary to compare the microstructure of the ceramic prepared here with those produced by the polymer pyrolysing route. A noticeable resemblance of the ceramics prepared by the two methods is the structure and function of BN(C). In the ceramics, BN(C) has turbostratic structure and mainly consists of t-BN and t-carbon layers [32]. During the crystallization and the grain growth process, BN(C) has similar effects in hindering atomic diffusion, retarding grain growth and limiting grain size [33]. As a result, the grain sizes in the prepared ceramics have approximately the same order of magnitude, namely, 100 nm or so [34,35]. Another similarity is that no low-melting-point intercrystalline phases exist in the ceramics fabricated by the two methods. During PDCs’ preparation, the chemical composition can be precisely controlled and no additives are needed for pyrolysing [33]. In the current work, dense bulk ceramic can be prepared by high temperature and high pressure, without using any sintering additives. The phase composition of the current ceramic is similar to some polymer derived nano Si–B–C–N ceramics, especially when the chemical composition of the latter is located in the four-phase field of SiC + BN + C + B4C (1400 °C, 1 bar N2) [36]. Namely, depending on the composition, the polymer type and the way of synthesis, the polymer derived nano Si–B–C–N ceramics may also consist of bSiC, a-SiC and turbostratic BN(C) [35,36]. Except for the similarities, differences also exist in the ceramics prepared by the two methods. As a mechanical process, mechanical alloying cannot produce powder with extremely uniform elements, like that in polymer precursors. As a result, the amorphous SiBCN powder prepared by the current method may have lower crystallization temperature. The types and the contents of various chemical bonds in the as-milled amorphous SiBCN powder may be quite different from the polymer derived amorphous bodies. Perhaps, it is difficult using the current method to prepare nano ceramic containing a/b-Si3N4, like that by polymer pyrolysing route. These will be studied in the future work. Another difference is that by the current method, nitrogen cannot be introduced independently, and hence the composition design is less flexible. However, using mechanical alloying and hot pressing method, the dense bulk ceramic with large dimensions can be easily prepared, facilitating their property evaluation. Currently, it is still in its infancy using mechanical alloying and hot pressing to prepare Si–B–C–N ceramics, and in the future, lots of works will be carried out on the theoretical study and engineering applications.

3.7. Properties of the hot pressed nano SiC/BN(C) ceramic Table 2 shows the room temperature mechanical properties and the average thermal expansion coefficient of the prepared nano

SiC/BN(C) ceramic. The ceramic hot prepared at 1900 °C/80 MPa has the room-temperature density, flexural strength, Young’s modulus, fracture toughness and Vickers’ hardness of 2.6 ± 0.1 g/cm3, 331.1 ± 40.5 MPa, 139.4 ± 16.0 GPa, 2.8 ± 0.9 MPa m1/2 and 2.7 ± 0.4 GPa, respectively. As zirconia contamination is well controlled during the milling process, the properties of the prepared SiC/BN(C) ceramic may be less affected by impurities. It is found in the current research that the mechanically alloyed SiBCN powder is very hard to densify, and the sintering temperature and pressure greatly affect the ceramic density and mechanical properties. As shown in Table 2, when the pressure is decreased to 50 MPa (under the same temperature of 1900 °C), the density, flexural strength, Young’s modulus and fracture toughness are rapidly reduced to 2.2 g/cm3, 89.4 MPa, 33.6 GPa and 1.2 MPa m1/2, respectively. On the other hand, when the temperature is lowered to 1800 °C (the pressure is maintained at 80 MPa), these properties are decreased to 2.1 g/cm3, 66.3 MPa, 30.6 GPa and 0.8 MPa m1/2, respectively. Using the current method, bulk ceramic with large dimensions can be easily prepared, and various mechanical properties can be evaluated conveniently. This is superior to the organic polymer pyrolysing method. Actually, except for fibers, the mechanical property study of the polymer derived ceramics (PDCs) has been hindered mainly due to the limitation in the fabrication of suitable bulk test specimens. In the past research, two systems of PDCs have been investigated in some detail: Si–C–O and Si–C–N, whose mechanical properties have been measured at different pyrolysis temperatures [37–39], while a systematic investigation for the influence of chemical composition has been studied only for the Si–C–O system [38,40]. For polymer derived bulk Si–B–C–N ceramic, the mechanical property study is limited to the investigation of high temperature creep resistance [41]. As the current work has successfully prepared compact bulk SiC/BN(C) ceramic with similar microstructure and has obtained its mechanical properties, the results can be used for reference to some degree.

4. Conclusions In the current research, nano SiC/BN(C) ceramic was prepared by the mechanical alloying and hot pressing method, and its microstructures and properties were carefully studied. Results show that the mechanically alloyed SiBCN powder is completely amorphous, and consists of near-spherical agglomerates (deriving from the nano particles) with an average particle size of 6.5 ± 5.4 lm. The prepared ceramic is composed of b-SiC, BN(C) and a small amount of a-SiC. SiC has an average grain size of 78.2 ± 32.4 nm and contains considerable stacking faults in most grains. BN(C) has small size, no fixed shape, relatively uniform distribution and turbostratic structure, comprising of heterogeneously distributed t-carbon layers, t-BN layers and B doped t-carbon layers. The interfacial region between SiC and BN(C) is free from any impurity and is limited to three to four atomic layers. Furthermore, BN(C) enwraps SiC in the grain boundary, retarding the atomic diffusion and being responsible for the fine grains. The ceramic is difficult to densify,

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and the sintering parameters (temperature and pressure) greatly affect the ceramic density and mechanical properties. The ceramic hot pressed at 1900 °C/80 MPa has the room temperature density, flexural strength, Young’s modulus, fracture toughness and Vickers’ hardness of 2.6 ± 0.1 g/cm3, 331.1 ± 40.5 MPa, 139.4 ± 16.0 GPa, 2.8 ± 0.9 MPa m1/2 and 2.7 ± 0.4 GPa, respectively. Acknowledgements The authors are grateful to the financial support from the National Natural Science Foundation of China under Grant Nos. 51072041, 50902031 and 5102100. References [1] M. Weinmann, J. Schuhmacher, H. Kummer, S. Prinz, J.Q. Peng, H.J. Seifert, M. Christ, K. Muller, J. Bill, F. Aldinger, Chem. Mater. 12 (2000) 623–632. [2] R. Riedel, A. Kienzle, W. Dressler, L. Ruwisch, J. Bill, F. Aldinger, Nature 382 (1996) 796–798. [3] M. Weinmann, T.W. Kamphowe, J. Schuhmacher, K. Muller, F. Aldinger, Chem. Mater. 12 (2000) 2112–2122. [4] M. Christ, G. Thurn, M. Weinmann, J. Bill, F. Aldinger, J. Am. Ceram. Soc. 83 (2000) 3025–3032. [5] N.V.R. Kumar, R. Mager, Y. Cai, A. Zimmermann, F. Aldinger, Scripta Mater. 51 (2004) 65–69. [6] P. Greil, Adv. Eng. Mater. 2 (2000) 339–348. [7] E. Kroke, Y.L. Li, C. Konetschny, E. Lecomte, C. Fasel, R. Riedel, Mater. Sci. Eng., R 26 (2000) 97–199. [8] J. Haug, P. Lamparter, M. Weinmann, F. Aldinger, Chem. Mater. 16 (2004) 83– 92. [9] J. Schuhmacher, F. Berger, M. Weinmann, J. Bill, F. Aldinger, K. Muller, Appl. Organomet. Chem. 15 (2001) 809–819. [10] H. Schmidt, G. Borchardt, O. Kaitasov, B. Lesage, J. Non-Cryst. Solids 353 (2007) 4801–4805. [11] H. Schmidt, W. Gruber, G. Borchardt, P. Gerstel, A. Muller, N. Bunjes, J. Eur. Ceram. Soc. 25 (2005) 227–231. [12] R. Kumar, Y. Cai, P. Gerstel, G. Rixecker, F. Aldinger, J. Mater. Sci. 41 (2006) 7088–7095. [13] P. Colombo, G. Mera, R. Riedel, G.D. Soraru, J. Am. Ceram. Soc. 93 (2010) 1805– 1837.

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