- Email: [email protected]
Machinable ZrB2–SiC–BN composites fabricated by reactive spark plasma sintering
Materials Science & Engineering A 582 (2013) 41–46
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Machinable ZrB2–SiC–BN composites fabricated by reactive spark plasma sintering Wen-Wen Wu a,n, Mehdi Estili a, Toshiyuki Nishimura b, Guo-Jun Zhang c, Yoshio Sakka a a
Advanced Ceramics Group, National Institute for Materials Science (NIMS), 1-2-1 Tsukuba, Ibaraki 305-0047, Japan Sialon group, Sialon Unit, National Institute for Materials Science (NIMS), 1-1 Namiki, Tsukuba 305-0044, Japan c State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics, Shanghai 200050, China b
art ic l e i nf o
a b s t r a c t
Article history: Received 6 March 2013 Received in revised form 20 May 2013 Accepted 21 May 2013 Available online 25 June 2013
A solid state reaction of ZrH2, Si3N4 and B4C has been proposed to synthesize machinable ZrB2–SiC–BN composites. Dense ZrB2–SiC–BN composites with fine grain size and homogeneous microstructure were fabricated via spark plasma sintering at 1900 1C in vacuum. The kinetics of the reaction process was studied with XRD results. SEM and TEM results showed that the in-situ formed BN phase was composed of micro-sized intergranular and nano-sized intragranular h-BN particles. The influences of the BN content on the mechanical properties as well as the thermal conductivities of the composites were studied. The composites that contained 4 20 vol% BN exhibited excellent machinability and high strength. & 2013 Elsevier B.V. All rights reserved.
Keywords: ZrB2 BN Ceramics Reactive sintering Spark plasma sintering
1. Introduction Zirconium diboride (ZrB2) is one of a group of compounds known as ultra-high temperature ceramics (UHTCs), which exhibit melting temperatures in excess of 3000 1C [1]. In particular, the high flexural strength (500–1000 MPa) and good oxidation resistance ( 1600 1C) make the ZrB2-20 vol% SiC composite the baseline of UHTCs for a variety of applications such as high thermal protection systems for hypersonic aerospace vehicles, refractory linings or molten metal crucibles [2–4]. The strong covalent bonding and low self-diffusion, densification of the ZrB2-based materials by hot pressing (HP) with high temperatures ( 41800 1C) and external pressures (4 20 MPa) is the dominant technique. Spark plasma sintering (SPS) shows its advantage at rapid heating rate which enables densification without significant grain growth. However, the external pressure applied in both methods restricts the materials' applications in preparation of complex shape components. Besides, the hot pressing conditions would limit the depletion of some impurities especially B2O3, which enhances grain growth of ZrB2 and decreases the mechanical properties. The pressureless sintering (PL) of ZrB2, which has been realized by using additives such as
n
Corresponding author. Tel.: +81 2985 92461; fax: +81 2985 92401. E-mail addresses: [email protected] (W.-W. Wu), [email protected] (Y. Sakka). 0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.05.079
carbon, boron carbide and boron to reduce the surface oxides [5–7], offers the possibility to fabricate components to near net shape, however, the high sintering temperatures (1800–2000 1C) together with the long isothermal holds (1–2 h) lead to a significant grain growth during sintering, resulting in a degradation of the strength of the materials. One possible solution to overcome these shortcomings is the fabrication of the machinable h-BN containing ZrB2 based ceramics [8,9], which has been successfully applied in several other structural ceramics, such as SiC, Si3N4 and so on [10–12]. On the other hand, intrinsic characteristics of low fracture toughness and poor thermal shock resistance create an obstacle for the widespread implementation of ZrB2-based ceramics, especially in extreme environment applications. Boron nitride, which has a low Young's modulus and cleavage properties due to its graphite-like layered structure and weak van der Waals forces between the stacked layers, as well as its stability at high temperatures due to strong covalent bonding of boron and nitrogen in the planar hexagonal networks, has been widely used as the secondary phase to reduce the elastic modulus and improve the thermal shock resistance of monolithic ceramic materials [13–15]. It is expected that the incorporation of BN phase into ZrB2–SiC composites is expected to improve their machinability and thermal shock resistance without compromising the excellent refractoriness of ZrB2 and SiC. With the cleavage properties and low fracture strength of BN, it is difficult to obtain high strength materials by conventional
42
W.-W. Wu et al. / Materials Science & Engineering A 582 (2013) 41–46
processes. Large BN agglomerates or platelets act as fracture origins which consequently lowers material strength. In order to overcome these difficulties, several chemical routes have been developed to fabricate high-performance BN-containing composites with highly homogeneous dispersion of BN particles (normally nano-size). In general, BN phase was in-situ formed from the reactions between various boron sources (such as B, B4C, AlB2, SiB4, SiB6, B2O3 and H3BO3) and nitrogen source (such as N2, melamine, Si3N4 and AlN), or from the decomposition of NH4HB4O7 in ammonia [16–23]. The superior mechanical properties of the asreceived composites alleviated the shortcoming of low strength encountered in the conventional composites which were fabricated from the powder mixture with h-BN. In the current work, a new solid state reaction was proposed to fabricate the high performance ZrB2–SiC–BN composites, by using ZrH2, Si3N4 and B4C powders as the starting materials according to reaction (1). 4ZrH2 þ Si3 N4 þ 3B4 C ¼ 4ZrB2 þ 3SiC þ 4BN þ 4H2
ð1Þ
ΔG298 ¼ −988:0 kJ The reactive sintering of the ZrB2–SiC–BN composites with various concentration of the BN phase was conducted by SPS. By virtue of special heat effects such as Joule heat, electromagnetic field and electrical discharge, the reactive spark plasma sintering (RSPS) enables faster densification, shorter sintering time, and more homogeneous microstructure in comparison with the conventional sintering technique [24–26]. No necessary treatment of the raw powders was needed for the reactive sintering, which largely decreased the complexity of the processing. The synthesis, microstructure, mechanical properties, machinability as well as the thermal conductivity of the as-received materials were investigated. The influence of the BN content on the microstructure and properties was discussed.
2. Experimental procedure Four ZrB2–SiC–BN composites with various BN content were prepared. ZrH2 (5 μm, purity 98%, Kojundo Chemical Laboratory Co., Ltd., Japan), Si3N4 (UBE SNE10, Japan) and B4C (Grade HS, H.C. Starck, Berlin, Germany) were used to synthesize ZrB2–SiC–BN composite with 30 vol% BN phase (ZSBN30). Commercially available ZrB2 (5–10 μm, purity 97%, Kojundo Chemical Laboratory Co., Ltd., Japan) and β-SiC (0.03 μm, Sumimoto Osaka Cement Co., Ltd., Japan) powder were added to fabricate composites with 5 vol% (ZSBN5), 10 vol% (ZSBN10), and 20 vol% (ZSBN20) of BN phase, respectively. ZS samples with only ZrB2 and SiC phases were also prepared for comparison. Before mixing, ZrB2 powder was milled with SiC media in ethanol for 24 h by a planetary ball mill (Pulverisette 6, Fritsch GmbH, Germany) to decrease the particle size. The powder was then dried in air. The D50 of the received ZrB2 powder is 0.7 μm. The starting powder compositions and the final compositions were listed in Table 1. Powders were mixed and ball-milled for 24 h in a polyethylene bottle charged with ethanol using Si3N4 balls. Solvent was then removed in air for 24 h. The dried powder mixtures were sintered by SPS (100 kN SPS-1050, Syntex Inc., Japan) in vacuum at 1900 1C and 40 MPa pressure for 5 min in graphite dies, the heating rate was 50 1C/min. The densities of as-prepared samples were measured by the Archimedes method. Theoretical densities were assessed by the rule-of-mixture (see in Table 1). The phase compositions were identified by an X-ray diffraction (XRD) analyzer (Model RINT 2500, Rigaku Co., Tokyo, Japan, 40 kV–300 mA) with Cu Kα radiation. The morphologies of polished and fracture surface of samples were investigated by a scanning electron microscope (SEM) (JSM-6500F, JEOL Ltd., Japan). The flexural strength at room
temperature was conducted by a three-point bending test on a mechanical strength testing system (Model 4505, Instron Corp., MA), using 2 mm 1.5 mm 18 mm chamfered bars with a 16 mm span and a crosshead speed of 0.5 mm/min. Fracture toughness was evaluated by a single-edge notched beam test with a 16 mm span and a crosshead speed of 0.05 mm/min using 2 mm 4 mm 18 mm test bars, the notch was about 2 mm deep and 0.14 mm wide. Hardness was determined by Vickers indentation using a diamond indenter (MVK-E, Akashi Co., Japan) under a load of 98 N for 15 s. All the reported values were the average of 5 data measurements. The elastic property was evaluated by an ultrasonic equipment (TDS 3034B, Tektronix Inc., USA). The machinability of each composite was tested using a 1.7 mm diameter hard steel drill. Additionally, thermal conductivity measurements were performed by a Xenon flash apparatus (LFA447 Nanoflash, Netzsch, Germany) at room temperature. A disk specimen with a dimension of Φ10 mm 2 mm was used. A layer of colloidal graphite spray was coated on the surface of the sample for enhancing the absorption of Xenon light pulse energy and the emission of infrared (IR) radiation to the temperature detector.
3. Results and discussion 3.1. Synthesis ZrB2–SiC–BN composites by reactive spark plasma sintering Fig. 1 shows the XRD results of ceramics containing various contents of BN reactively spark plasma sintered at 1900 1C. The reactions were completed during the sintering, as all samples contained phases of ZrB2, β-SiC and h-BN, except for ZS which is free of BN. From ZS to ZSBN30, the intensity of h-BN peaks Table 1 Compositions of the prepared ZrB2–SiC–BN composite samples. Starting composition (wt%)
ZSBN30 ZSBN20 ZSBN10 ZSBN5 ZS
Final Composition (vol%)
ZrH2
B4C
Si3N4 ZrB2
SiC
ZrB2 SiC BN
54.38 32.73 15.11 7.29 0
24.70 14.87 6.87 3.31 0
20.91 12.59 5.81 2.80 0
0 2.82 7.59 9.73 11.64
46 60 50 75 80
0 37.00 64.62 76.86 88.36
24 20 20 20 20
30 20 10 5 0
Theoretical density (g/cm3)
4.22 4.74 5.13 5.33 5.52
Fig. 1. XRD patterns of the ZrB2–SiC–BN composites fabricated by RSPS at 1900 1C: (a) ZSBN30, (b) ZSBN20, (c) ZSBN10, (d) ZSBN5 and (e) ZS.
W.-W. Wu et al. / Materials Science & Engineering A 582 (2013) 41–46
43
In addition, the appearance of ZrO2 could be attributed to the incomplete oxide impurities' evaporation, due to the uniaxial pressure applied during the RSPS process. 3.2. Microstructure
Fig. 2. XRD patterns of the intermediate products that reactive spark plasma sintered from ZrH2–Si3N4–B4C raw powder mixture at: (a) 900 1C; (b) 1000 1C; (c) 1200 1C; (d) 1400 1C and (e) 1600 1C in vacuum.
increased, and that of ZrB2 and SiC has no obvious change. Low peaks of α-SiC were detected in some samples, indicating that phase transformation of β-SiC into α-SiC occurred during sintering. The presence of ZrO2 in all the reactive sintered samples can be attributed to the oxide impurities on the surface of the starting powders. Fig. 2 illustrates the conversion of crystalline phases in the products synthesized by heating a mixture of ZrH2, Si3N4 and B4C at a rate of 50 1C/min using RSPS. At 900 1C, only small amount of ZrH was left, most of the ZrH2 has been decomposed into Zr and H2. The Si3N4 phase was retained without decomposition or reaction. Peaks of B4C were not found probably because the diffusion and dissolution of amorphous boron and carbon into the Zr metal [27,28] and the low sensitivity of the light elements to X-ray. During the heating from 900 1C to 1000 1C, a sharp decrease of vacuum level (from 6 Pa to 70 Pa, then increased again to o10 Pa in 1 min) in the SPS furnace chamber occurred. XRD patterns at 1000 1C indicated no decomposition of Si3N4 but the formation ZrC and ZrB2, with unchanged Si3N4 peaks and the much higher peaks of ZrC than those of ZrB2. The easier formation of ZrC than ZrB2 could be attributed to the activation energy for diffusion of boron in Zr (34.5 7 8 Kcal/mol [29]) is higher than that for carbon in Zr (17.975.6 Kcal/mol [30]). The temporary degradation of the vacuum may associate with evaporation of the B2O3 impurities on the surface of the B4C raw powders because the formation of ZrC and ZrB2 are highly exothermic. Further increase of the temperature to 1200 1C and 1400 1C, the intensity of peaks of ZrC and Si3N4 decreased, in contrast, the intensity of ZrB2 increased. Upon heat treatment at 1600 1C, the peaks of Si3N4 and ZrC disappeared, SiC and the broad peak of BN was observed, meanwhile, the peaks of ZrO2 appeared. According to the above results, the reactions between ZrH2, Si3N4 and B4C took place in the following steps, as described in Eqs. (2)–(7): ZrH2 ¼ Zr þ H2
ð2Þ
Zr þ B4 C ¼ ZrC þ 4B ðamorphousÞ
ð3Þ
Zr þ 2B ðamorphousÞ ¼ ZrB2
ð4Þ
ZrC þ 2B ðamorphousÞ ¼ ZrB2 þ C ðamorphousÞ
ð5Þ
Si3 N4 þ 4B ðamorphousÞ þ 3C ðamorphousÞ ¼ 3SiC þ 4BN
ð6Þ
3ZrC þ Si3 N4 þ 10B ðamorphousÞ ¼ 3ZrB2 þ 3SiC þ 4BN
ð7Þ
All the samples were fully densified at the sintering temperature of 1900 1C with relative densities 499%. Fig. 3(a–e) shows the SEM images of the polished surfaces of the reactive sintered ZrB2– SiC–BN composites. The composites were free of significant porosity, although voids associated with the peeling off of the weak BN were occasionally observed. All samples exhibit fine and homogeneous microstructures. The BN content dependence of the mean grain size of ZrB2 and SiC in the ZrB2–SiC–BN composites is shown in Fig. 3(f). The mean sizes of the SiC grains (about 1 μm) were not significantly influenced by the amount of BN, while the mean size of the ZrB2 grains first increased and then decreased. The maximum ZrB2 particle size is about 2.3 μm when the BN volume is 5%. The promoted grain growth of ZrB2 with small amount addition of BN is because the sinterbility of the composite was enhanced by the reactive sintering introduced into the ZrB2 and SiC powder mixtures. However, further increasing of the BN phase inhibited the grain growth, because of the poor sinterbility of h-BN which has strong covalent nature and plate-like structure. Although ZrB2 and SiC phases in the samples (except ZSN30) come from two ways: directly from the raw ceramic powders and by solid reaction formation, it is very difficult to distinguish them in microstructures, because no phase and morphology difference was observed. Both ZrB2 and SiC grains in all composites are equiaxed. The added ZrB2 and SiC powders has a much smaller grain size (0.7 μm for ZrB2 and 30 nm for SiC) than the final grain sizes ( 2 μm for ZrB2 and 1 μm for SiC), indicating the significant grain growth for both the added phases and the in-situ formed phases. The micro-sized h-BN grains that homogeneously located at the grain boundaries of ZrB2 and SiC grains had a flakelike morphology, with a length of 2–3 μm, and thickness o1 μm. The anisotropic grain growth of the h-BN plates were found in the composites that contain BN content lower than 20 vol%. As the amount of the BN phase increased, the size and the aspect ratio of the BN grains decreased, because of the distortion of the BN flakes to accommodate the grain boundary space in the composites. The buckled and delaminated h-BN flakes that located in the vicinity of the ZrB2 and SiC interfaces were observed on the fracture surfaces of ZSBN20 and ZSBN30 that fabricated by RSPS, as can be seen in Fig.4. It is thought that the residual thermal stress in the sample have played an important role in creating such phenomena. ZrB2 has a mean thermal expansion coefficient of 5.2 10−6 1C−1 while SiC is 3.3 10−6 1C−1. Watts et al. [31] estimated that a compressive stress of 880 MPa in the SiC phase, and 450 MPa tensile in the ZrB2 were generated in the ZrB2–SiC composites, during the cooling process, due to the thermal expansion mismatch of the two phases. The thermal expansion coefficient of h-BN is highly anisotropic, it shows a high value of ∼40 10−6 1C−1 and low value of ≤10−7 1C−1 in the directions parallel and perpendicular to the c-axis, respectively [23,32]. The thermal mismatch between h-BN, ZrB2, and SiC made the intergranular h-BN particle subject to a tensile stress and a compressive stress in the directions parallel and perpendicular to the c-axis when the sample was cooled down to room temperature after sintering. The tensile stress led to the delamination of h-BN, while the compressive stress caused the bucking of the h-BN layers. Fig. 5 shows a representative TEM micrograph for the reactive spark plasma sintered ZrB2–SiC–BN composite processed in this work. Some h-BN particles (as arrowed in Fig. 5) were entrapped into ZrB2 grains. The intragranular h-BN particles showed a dimension of less than 200 nm, which is much less than that of
44
W.-W. Wu et al. / Materials Science & Engineering A 582 (2013) 41–46
Fig. 3. SEM micrographs of polished surfaces of the ZrB2-SiC-BN composites that were fabricated by RSPS: (a) ZS, (b) ZSBN5, (c) ZSBN10, (d) ZSBN20, (e) ZSBN30 and (f) BN content dependence of the grain size of ZrB2 and SiC.
Fig. 4. Fracture surface of ZSBN20 that fabricated by reactive spark plasma sintering.
Fig. 5. TEM image of ZSBN30 that fabricated by reactive spark plasma sintering. The inserts are the electron diffraction of the particles of circle A and B.
3.3. Mechanical properties the intergranular h-BN flakes that were located at the grain boundaries. This result indicates that the originally in-situ formed BN phase was nano-sized, significant grain growth occurred during the sintering at a relatively high temperature of 1900 1C.
Fig. 6(a) is the hardness and Young's modulus of the ZrB2–SiC–BN ceramics as a function of the content of BN. As expected from the rule of mixtures for the mixture of ZrB2, SiC and h-BN, both of the values decreased with the increasing of BN content, because BN is a soft and
W.-W. Wu et al. / Materials Science & Engineering A 582 (2013) 41–46
45
Fig. 6. Young's modulus and hardness (a) and the flexural strength and fracture toughness and (b) of the reactive spark plasma sintered ZrB2–SiC–BN composites as a function of the content of BN.
Griffith equation 1 K IC pffiffiffi sf ¼ Y c
Fig. 7. SEM micrographs of the polished surface of ZSBN5. The microcracks formed by the residual stress are pointed by black arrows.
weak phase owning to its graphite-like crystal structure, and Young's modulus of BN (65 GPa [18]) is much lower than ZrB2 and SiC (489 GPa [33] and 427 GPa [17], respectively). Fig. 6(b) shows the variation in flexural strength and fracture toughness for the ZrB2–SiC–BN ceramics as a function of the content of BN. The fracture toughness increased from 5.8 70.3 MPa m1/2 for the ZS composite to a maximum value of 6.6 70.2 MPa m1/2 for the ZSBN5 composite, then decreased to 4.3 70.1 MPa m1/2 with the increasing of the BN content to 30 vol %. There are two reasons for the increasing of the fracture toughness. First, the existence of the softness of h-BN and the relatively large ZrB2 grain size at the small h-BN amount led to the crack of the composite propagating preferentially along the weak boundary particles or their interfaces, resulting crack bridging, deflection, or a tortuous crack path. Second, the thermal mismatches of the phases and the grain growth of ZrB2 produced high local residual stresses providing crack-particle interaction and the potential for microcracking (as pointed by black arrows in Fig. 7). The flexural strength decreased as the BN content increased, similar to the other BN containing microcomposites in which the BN grain size is in micron scale. The weak h-BN at the interface of ZrB2 and SiC changed the fracture mode from intragranular into intergranular, and decreased the strength of the composites. Nevertheless, the measured flexural strength from ZSN20 and ZSBN30 (518.8 723.4 MPa and 457.7 726.4 MPa, respectively) are comparable with that of the ZrB2–SiC and higher that the ZrB2–SiC–BN composites which were sintered by the conventional sintering methods [33–35]. The fine and homogeneous microstructure is believed to have an important role in the improvement of the mechanical properties. The flexural strength measured from ZSBN5 was relatively low compared to that of ZSBN10, which can be explained by the
where sf , K IC , c, and Y are flexural strength, fracture toughness, one-half of the width of the initial flaw, and the geometrical parameter of a flaw, respectively. Based on this fracture mechanics, the strength of the brittle ceramics is controlled much stronger by the critical flaw size than the fracture toughness. In a result, although the fracture toughness of ZSN5 was higher, the large flaw size associated with the large ZrB2 grains in ZSBN5 decreased the strength. The increasing of the BN content decreased the grain size of ZrB2, in addition, the 10 vol% of BN did not reach the threshold for the formation of BN a percolation network. As a result, the strength of ZSN10 increased slightly. Further increasing of the BN content to above 10 vol% degraded the material strength, because the BN percolation network can facilitate defect formation at the grain boundaries [36]. Several reports have demonstrated that the nanocomposites which have nano-sized h-BN coated on the matrix particles show much higher flexural strength in comparison with the conventionally fabricated microcomposites [10,18,19,21,23]. The reactive sintering has shown its advantage at refining the microstructure of the composites by low temperature sintering which inhibited grain growth of ZrB2-based ceramics [37,38]. It is expected that the flexural strength of the ZrB2–SiC–BN composites fabricated in this study can be further enhanced by careful processing control in the future. 3.4. Machinability Fig. 8 shows the indent induced by a load of 10 kg during hardness test on polished surfaces of ZSBN30. The absence of classical radial cracks around the indentation site indicates a remarkable contact damage resistance of the composite [39]. The formation of weak interfaces and the cleavage behavior of h-BN particles absorbed the fracture energy during the hardness test and prevented the crack propagation. Fig. 9 shows the photograph of a ZSBN20 plate, on which holes were drilled with a hard steel drill. The ZrB2–SiC–BN composites with BN content higher than 20 vol% were highly machinable. The intergranular h-BN is thought to have elevated the machinability and deformability due to their easy deformation and cleavage characteristics [11]. 3.5. Thermal conductivity The measured heat capacities, thermal diffusivity and the calculated thermal conductivities of the ZrB2–SiC–BN composites fabricated by RSPS are listed in Table 2. The thermal conductivity of the composites decreased with the increase of BN, because of the increased grain boundaries and the random orientation of the
46
W.-W. Wu et al. / Materials Science & Engineering A 582 (2013) 41–46
increased and then decreased with the increasing of the BN content. The flexural strength and the thermal conductivities of the composites decreased with the increasing of BN, while the fracture toughness first increased and then decreased. The composite that contained 420 wt% BN exhibited excellent machinability. Although most of the in-situ formed h-BN particle have a relative large grain size (2–3 μm high flexural strength of 518.8 7 23.4 MPa for the composite with 20 wt% BN was retained.
Acknowledgments
Fig. 8. SEM micrographs of the indent that introduced on the polished surface of ZSBN30 with a diamond indenter at the load of 10 kg.
Financial supports from the National Natural Science Foundation of China (No. 51272266) and the bilateral project of NSFC-JSPS (No. 51111140017) are gratefully acknowledged. References
Fig. 9. Machinable ZrB2–SiC–BN composite containing 20 vol% BN fabricated by RSPS. The diameter of the disk is 25 mm.
Table 2 Heat capacity, thermal diffusivity and thermal conductivity of the ZrB2–SiC–BN composites fabricated by RSPS.
ZSBN30 ZSBN20 ZSBN10 ZSBN5 ZS
Heat capacity Cp (J (g K)−1)
Thermal diffusivity α (mm2/s)
Thermal conductivity kc (W (m K)−1)
0.52 0.53 0.53 0.53 0.56
31.09 32.11 31.18 27.34 27.58
69.3 76.8 79.2 76.9 86.6
low conductive h-BN (in a range of 30–60 W (m K)−1) [40–42]. Nevertheless, the values received in this work is comparable to that of the reported value of ZrB2–SiC composites (60–100 W (m K)−1) [43], which can be attributed to the clean grain boundary that was formed during the reactive sintering. 4. Conclusions Highly dense ZrB2–SiC–BN composites with fine and homogeneous microstructure were successfully fabricated from a mixture of ZrH2, Si3N4 and B4C powders, via reactive spark plasma sintering at 1900 1C in vacuum. The in-situ formed BN phase was composed of micro-sized intergranular and nano-sized intragranular h-BN particles. The particle size of the ZrB2 grains first
[1] K. Upadhya, J.M. Yang, W.P. Hoffman, Am. Ceram. Soc. Bull. 76 (1997) 51–56. [2] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, J. Mater. Sci. 39 (2004) 5887–5904. [3] E. Wuchina, E. Opila, M. Opeka, W. Fahrenholtz, I. Talmy, Electrochem. Soc. Interface 16 (2007) 30–36. [4] A. Paul, D.D. Jayaseelan, S. Venugopal, E. Zapata-Solvas, J. Binner, B. Vaidhyanathan, A. Heaton, P. Brown, W.E. Lee, Am. Ceram. Soc. Bull. 91 (2012) 22–28. [5] S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, J. Am. Ceram. Soc. 91 (2008) 26–32. [6] W.G. Fahrenholtz, G.E. Hilmas, S.C. Zhang, S. Zhu, J. Am. Ceram. Soc. 91 (2008) 1398–1404. [7] W.M. Guo, G.J. Zhang, Z.G. Yang, J. Am. Ceram. Soc. 95 (2012) 2470–2473. [8] G. Li, W.B. Han, B.L. Wang, Mater. Des. 32 (2011) 401–405. [9] C.C. Wei, X.H. Zhang, P. Hu, W.B. Han, G.S. Tian, Scr. Mater. 65 (2011) 791–794. [10] T. Kusunose, T. Sekino, Y.H. Choa, K. Niihara, J. Am. Ceram. Soc. 85 (2002) 2689–2695. [11] Y.L. Li, G.J. Qiao, Z.H. Jin, Mater. Res. Bull. 37 (2002) 1401–1409. [12] X.D. Wang, G.J. Qiao, Z.H. Jin, J. Am. Ceram. Soc. 87 (2004) 565–570. [13] K.S. Mazdiyasni, R. Ruh, J. Am. Ceram. Soc. 64 (1981) 415–419. [14] R. Ruh, A. Zangvil, R.R. Wills, Adv. Ceram. Mater. 3 (1988) 411–415. [15] E.H. Lutz, M.V. Swain, J. Am. Ceram. Soc. 75 (1992) 67–70. [16] G.J. Zhang, M. Ando, T. Ohji, S. Kanzaki, Int. J. Appl. Ceram. Technol. 2 (2005) 162–171. [17] G.J. Zhang, Y. Beppu, T. Ohji, S. Kanzaki, Acta Mater. 49 (2001) 77–82. [18] T. Kusunose, T. Sekino, Y.H. Choa, K. Niihara, J. Am. Ceram. Soc. 85 (2002) 2678–2688. [19] T. Kusunose, T. Sekino, Y. Ando, Nanotechnology 19 (2008) 275603. [20] T. Kusunose, N. Sakayanagi, T. Sekino, Y. Ando, J. Nanosci. Nanotechnol. 8 (2008) 5846–5853. [21] T. Kusunose, Y.H. Kim, T. Sekino, T. Matsumoto, N. Tanaka, T. Nakayama, K. Niihara, J. Mater. Res. 20 (2005) 183–190. [22] R. Shuba, I.W. Chen, J. Am. Ceram. Soc. 89 (2006) 2147–2153. [23] L. Gao, X.H. Jin, J.G. Li, Y.G. Li, J. Sun, Mater. Sci. Eng. A 415 (2006) 145–148. [24] S. Grasso, Y. Sakka, G. Maizza, Sci. Technol. Adv. Mater. 10 (2009) 053001. [25] R. Orru, R. Licheri, A.M. Locci, A. Cincotti, G.C. Cao, Mater. Sci. Eng. R 63 (2009) 127–287. [26] W.W. Wu, G.J. Zhang, Y.M. Kan, P.L. Wang, K. Vanmeensel, J. Vleugels, O. Van der Biest, Scr. Mater. 57 (2007) 317–320. [27] G.J. Zhang, Z.Y. Deng, N. Kondo, J.F. Yang, T. Ohji, J. Am. Ceram. Soc. 83 (2000) 2330–2332. [28] G.J. Zhang, M. Ando, J.F. Yang, T. Ohji, S. Kanzaki, J. Eur. Ceram. Soc. 24 (2004) 171–178. [29] L.V. Strashinskaya, A.N. Stepanchuk, Fiz.-Khim. Mekh. Mater. 6 (1970) 76–79. [30] N.L. Peterson, Wadd Technical Report, 1960 pp. 60–793. [31] J. Watts, G. Hilmas, W.G. Fahrenholtz, D. Brown, B. Clausen, J. Eur. Ceram. Soc. 31 (2011) 1811–1820. [32] W. Sinclair, H. Simmons, J. Mater. Sci. Lett. 6 (1987) 627–629. [33] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, J. Am. Ceram. Soc. 90 (2007) 1347–1364. [34] S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, J. Eur. Ceram. Soc. 31 (2011) 893–901. [35] H.T. Wu, W.G. Zhang, J. Eur. Ceram. Soc. 30 (2010) 1035–1042. [36] G.J. Zhang, T. Ohji, J. Mater. Res. 15 (2000) 1876–1880. [37] W.W. Wu, G.J. Zhang, Y.M. Kan, P.L. Wang, J. Am. Ceram. Soc. 91 (2008) 2501–2508. [38] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, J. Am. Ceram. Soc. 89 (2006) 3638–3645. [39] X.T. Wang, N.P. Padture, H. Tanaka, Nat. Mater. 3 (2004) 539–544. [40] M. Hubacek, M. Ueki, J. Am. Ceram. Soc. 82 (1999) 156–160. [41] D.W. Ni, G.J. Zhang, Y.M. Kan, Y. Sakka, J. Am. Ceram. Soc. 94 (2011) 1397–1404. [42] A. Lipp, K.A. Schwetz, K. Hunold, J. Eur. Ceram. Soc. 5 (1989) 3–9. [43] J.W. Zimmermann, G.E. Hilmas, W.G. Fahrenholtz, R.B. Dinwiddie, W.D. Porter, H. Wang, J. Am. Ceram. Soc. 91 (2008) 1405–1411.
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Machinable ZrB2–SiC–BN composites fabricated by reactive spark plasma sintering Wen-Wen Wu a,n, Mehdi Estili a, Toshiyuki Nishimura b, Guo-Jun Zhang c, Yoshio Sakka a a
Advanced Ceramics Group, National Institute for Materials Science (NIMS), 1-2-1 Tsukuba, Ibaraki 305-0047, Japan Sialon group, Sialon Unit, National Institute for Materials Science (NIMS), 1-1 Namiki, Tsukuba 305-0044, Japan c State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics, Shanghai 200050, China b
art ic l e i nf o
a b s t r a c t
Article history: Received 6 March 2013 Received in revised form 20 May 2013 Accepted 21 May 2013 Available online 25 June 2013
A solid state reaction of ZrH2, Si3N4 and B4C has been proposed to synthesize machinable ZrB2–SiC–BN composites. Dense ZrB2–SiC–BN composites with fine grain size and homogeneous microstructure were fabricated via spark plasma sintering at 1900 1C in vacuum. The kinetics of the reaction process was studied with XRD results. SEM and TEM results showed that the in-situ formed BN phase was composed of micro-sized intergranular and nano-sized intragranular h-BN particles. The influences of the BN content on the mechanical properties as well as the thermal conductivities of the composites were studied. The composites that contained 4 20 vol% BN exhibited excellent machinability and high strength. & 2013 Elsevier B.V. All rights reserved.
Keywords: ZrB2 BN Ceramics Reactive sintering Spark plasma sintering
1. Introduction Zirconium diboride (ZrB2) is one of a group of compounds known as ultra-high temperature ceramics (UHTCs), which exhibit melting temperatures in excess of 3000 1C [1]. In particular, the high flexural strength (500–1000 MPa) and good oxidation resistance ( 1600 1C) make the ZrB2-20 vol% SiC composite the baseline of UHTCs for a variety of applications such as high thermal protection systems for hypersonic aerospace vehicles, refractory linings or molten metal crucibles [2–4]. The strong covalent bonding and low self-diffusion, densification of the ZrB2-based materials by hot pressing (HP) with high temperatures ( 41800 1C) and external pressures (4 20 MPa) is the dominant technique. Spark plasma sintering (SPS) shows its advantage at rapid heating rate which enables densification without significant grain growth. However, the external pressure applied in both methods restricts the materials' applications in preparation of complex shape components. Besides, the hot pressing conditions would limit the depletion of some impurities especially B2O3, which enhances grain growth of ZrB2 and decreases the mechanical properties. The pressureless sintering (PL) of ZrB2, which has been realized by using additives such as
n
Corresponding author. Tel.: +81 2985 92461; fax: +81 2985 92401. E-mail addresses: [email protected] (W.-W. Wu), [email protected] (Y. Sakka). 0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.05.079
carbon, boron carbide and boron to reduce the surface oxides [5–7], offers the possibility to fabricate components to near net shape, however, the high sintering temperatures (1800–2000 1C) together with the long isothermal holds (1–2 h) lead to a significant grain growth during sintering, resulting in a degradation of the strength of the materials. One possible solution to overcome these shortcomings is the fabrication of the machinable h-BN containing ZrB2 based ceramics [8,9], which has been successfully applied in several other structural ceramics, such as SiC, Si3N4 and so on [10–12]. On the other hand, intrinsic characteristics of low fracture toughness and poor thermal shock resistance create an obstacle for the widespread implementation of ZrB2-based ceramics, especially in extreme environment applications. Boron nitride, which has a low Young's modulus and cleavage properties due to its graphite-like layered structure and weak van der Waals forces between the stacked layers, as well as its stability at high temperatures due to strong covalent bonding of boron and nitrogen in the planar hexagonal networks, has been widely used as the secondary phase to reduce the elastic modulus and improve the thermal shock resistance of monolithic ceramic materials [13–15]. It is expected that the incorporation of BN phase into ZrB2–SiC composites is expected to improve their machinability and thermal shock resistance without compromising the excellent refractoriness of ZrB2 and SiC. With the cleavage properties and low fracture strength of BN, it is difficult to obtain high strength materials by conventional
42
W.-W. Wu et al. / Materials Science & Engineering A 582 (2013) 41–46
processes. Large BN agglomerates or platelets act as fracture origins which consequently lowers material strength. In order to overcome these difficulties, several chemical routes have been developed to fabricate high-performance BN-containing composites with highly homogeneous dispersion of BN particles (normally nano-size). In general, BN phase was in-situ formed from the reactions between various boron sources (such as B, B4C, AlB2, SiB4, SiB6, B2O3 and H3BO3) and nitrogen source (such as N2, melamine, Si3N4 and AlN), or from the decomposition of NH4HB4O7 in ammonia [16–23]. The superior mechanical properties of the asreceived composites alleviated the shortcoming of low strength encountered in the conventional composites which were fabricated from the powder mixture with h-BN. In the current work, a new solid state reaction was proposed to fabricate the high performance ZrB2–SiC–BN composites, by using ZrH2, Si3N4 and B4C powders as the starting materials according to reaction (1). 4ZrH2 þ Si3 N4 þ 3B4 C ¼ 4ZrB2 þ 3SiC þ 4BN þ 4H2
ð1Þ
ΔG298 ¼ −988:0 kJ The reactive sintering of the ZrB2–SiC–BN composites with various concentration of the BN phase was conducted by SPS. By virtue of special heat effects such as Joule heat, electromagnetic field and electrical discharge, the reactive spark plasma sintering (RSPS) enables faster densification, shorter sintering time, and more homogeneous microstructure in comparison with the conventional sintering technique [24–26]. No necessary treatment of the raw powders was needed for the reactive sintering, which largely decreased the complexity of the processing. The synthesis, microstructure, mechanical properties, machinability as well as the thermal conductivity of the as-received materials were investigated. The influence of the BN content on the microstructure and properties was discussed.
2. Experimental procedure Four ZrB2–SiC–BN composites with various BN content were prepared. ZrH2 (5 μm, purity 98%, Kojundo Chemical Laboratory Co., Ltd., Japan), Si3N4 (UBE SNE10, Japan) and B4C (Grade HS, H.C. Starck, Berlin, Germany) were used to synthesize ZrB2–SiC–BN composite with 30 vol% BN phase (ZSBN30). Commercially available ZrB2 (5–10 μm, purity 97%, Kojundo Chemical Laboratory Co., Ltd., Japan) and β-SiC (0.03 μm, Sumimoto Osaka Cement Co., Ltd., Japan) powder were added to fabricate composites with 5 vol% (ZSBN5), 10 vol% (ZSBN10), and 20 vol% (ZSBN20) of BN phase, respectively. ZS samples with only ZrB2 and SiC phases were also prepared for comparison. Before mixing, ZrB2 powder was milled with SiC media in ethanol for 24 h by a planetary ball mill (Pulverisette 6, Fritsch GmbH, Germany) to decrease the particle size. The powder was then dried in air. The D50 of the received ZrB2 powder is 0.7 μm. The starting powder compositions and the final compositions were listed in Table 1. Powders were mixed and ball-milled for 24 h in a polyethylene bottle charged with ethanol using Si3N4 balls. Solvent was then removed in air for 24 h. The dried powder mixtures were sintered by SPS (100 kN SPS-1050, Syntex Inc., Japan) in vacuum at 1900 1C and 40 MPa pressure for 5 min in graphite dies, the heating rate was 50 1C/min. The densities of as-prepared samples were measured by the Archimedes method. Theoretical densities were assessed by the rule-of-mixture (see in Table 1). The phase compositions were identified by an X-ray diffraction (XRD) analyzer (Model RINT 2500, Rigaku Co., Tokyo, Japan, 40 kV–300 mA) with Cu Kα radiation. The morphologies of polished and fracture surface of samples were investigated by a scanning electron microscope (SEM) (JSM-6500F, JEOL Ltd., Japan). The flexural strength at room
temperature was conducted by a three-point bending test on a mechanical strength testing system (Model 4505, Instron Corp., MA), using 2 mm 1.5 mm 18 mm chamfered bars with a 16 mm span and a crosshead speed of 0.5 mm/min. Fracture toughness was evaluated by a single-edge notched beam test with a 16 mm span and a crosshead speed of 0.05 mm/min using 2 mm 4 mm 18 mm test bars, the notch was about 2 mm deep and 0.14 mm wide. Hardness was determined by Vickers indentation using a diamond indenter (MVK-E, Akashi Co., Japan) under a load of 98 N for 15 s. All the reported values were the average of 5 data measurements. The elastic property was evaluated by an ultrasonic equipment (TDS 3034B, Tektronix Inc., USA). The machinability of each composite was tested using a 1.7 mm diameter hard steel drill. Additionally, thermal conductivity measurements were performed by a Xenon flash apparatus (LFA447 Nanoflash, Netzsch, Germany) at room temperature. A disk specimen with a dimension of Φ10 mm 2 mm was used. A layer of colloidal graphite spray was coated on the surface of the sample for enhancing the absorption of Xenon light pulse energy and the emission of infrared (IR) radiation to the temperature detector.
3. Results and discussion 3.1. Synthesis ZrB2–SiC–BN composites by reactive spark plasma sintering Fig. 1 shows the XRD results of ceramics containing various contents of BN reactively spark plasma sintered at 1900 1C. The reactions were completed during the sintering, as all samples contained phases of ZrB2, β-SiC and h-BN, except for ZS which is free of BN. From ZS to ZSBN30, the intensity of h-BN peaks Table 1 Compositions of the prepared ZrB2–SiC–BN composite samples. Starting composition (wt%)
ZSBN30 ZSBN20 ZSBN10 ZSBN5 ZS
Final Composition (vol%)
ZrH2
B4C
Si3N4 ZrB2
SiC
ZrB2 SiC BN
54.38 32.73 15.11 7.29 0
24.70 14.87 6.87 3.31 0
20.91 12.59 5.81 2.80 0
0 2.82 7.59 9.73 11.64
46 60 50 75 80
0 37.00 64.62 76.86 88.36
24 20 20 20 20
30 20 10 5 0
Theoretical density (g/cm3)
4.22 4.74 5.13 5.33 5.52
Fig. 1. XRD patterns of the ZrB2–SiC–BN composites fabricated by RSPS at 1900 1C: (a) ZSBN30, (b) ZSBN20, (c) ZSBN10, (d) ZSBN5 and (e) ZS.
W.-W. Wu et al. / Materials Science & Engineering A 582 (2013) 41–46
43
In addition, the appearance of ZrO2 could be attributed to the incomplete oxide impurities' evaporation, due to the uniaxial pressure applied during the RSPS process. 3.2. Microstructure
Fig. 2. XRD patterns of the intermediate products that reactive spark plasma sintered from ZrH2–Si3N4–B4C raw powder mixture at: (a) 900 1C; (b) 1000 1C; (c) 1200 1C; (d) 1400 1C and (e) 1600 1C in vacuum.
increased, and that of ZrB2 and SiC has no obvious change. Low peaks of α-SiC were detected in some samples, indicating that phase transformation of β-SiC into α-SiC occurred during sintering. The presence of ZrO2 in all the reactive sintered samples can be attributed to the oxide impurities on the surface of the starting powders. Fig. 2 illustrates the conversion of crystalline phases in the products synthesized by heating a mixture of ZrH2, Si3N4 and B4C at a rate of 50 1C/min using RSPS. At 900 1C, only small amount of ZrH was left, most of the ZrH2 has been decomposed into Zr and H2. The Si3N4 phase was retained without decomposition or reaction. Peaks of B4C were not found probably because the diffusion and dissolution of amorphous boron and carbon into the Zr metal [27,28] and the low sensitivity of the light elements to X-ray. During the heating from 900 1C to 1000 1C, a sharp decrease of vacuum level (from 6 Pa to 70 Pa, then increased again to o10 Pa in 1 min) in the SPS furnace chamber occurred. XRD patterns at 1000 1C indicated no decomposition of Si3N4 but the formation ZrC and ZrB2, with unchanged Si3N4 peaks and the much higher peaks of ZrC than those of ZrB2. The easier formation of ZrC than ZrB2 could be attributed to the activation energy for diffusion of boron in Zr (34.5 7 8 Kcal/mol [29]) is higher than that for carbon in Zr (17.975.6 Kcal/mol [30]). The temporary degradation of the vacuum may associate with evaporation of the B2O3 impurities on the surface of the B4C raw powders because the formation of ZrC and ZrB2 are highly exothermic. Further increase of the temperature to 1200 1C and 1400 1C, the intensity of peaks of ZrC and Si3N4 decreased, in contrast, the intensity of ZrB2 increased. Upon heat treatment at 1600 1C, the peaks of Si3N4 and ZrC disappeared, SiC and the broad peak of BN was observed, meanwhile, the peaks of ZrO2 appeared. According to the above results, the reactions between ZrH2, Si3N4 and B4C took place in the following steps, as described in Eqs. (2)–(7): ZrH2 ¼ Zr þ H2
ð2Þ
Zr þ B4 C ¼ ZrC þ 4B ðamorphousÞ
ð3Þ
Zr þ 2B ðamorphousÞ ¼ ZrB2
ð4Þ
ZrC þ 2B ðamorphousÞ ¼ ZrB2 þ C ðamorphousÞ
ð5Þ
Si3 N4 þ 4B ðamorphousÞ þ 3C ðamorphousÞ ¼ 3SiC þ 4BN
ð6Þ
3ZrC þ Si3 N4 þ 10B ðamorphousÞ ¼ 3ZrB2 þ 3SiC þ 4BN
ð7Þ
All the samples were fully densified at the sintering temperature of 1900 1C with relative densities 499%. Fig. 3(a–e) shows the SEM images of the polished surfaces of the reactive sintered ZrB2– SiC–BN composites. The composites were free of significant porosity, although voids associated with the peeling off of the weak BN were occasionally observed. All samples exhibit fine and homogeneous microstructures. The BN content dependence of the mean grain size of ZrB2 and SiC in the ZrB2–SiC–BN composites is shown in Fig. 3(f). The mean sizes of the SiC grains (about 1 μm) were not significantly influenced by the amount of BN, while the mean size of the ZrB2 grains first increased and then decreased. The maximum ZrB2 particle size is about 2.3 μm when the BN volume is 5%. The promoted grain growth of ZrB2 with small amount addition of BN is because the sinterbility of the composite was enhanced by the reactive sintering introduced into the ZrB2 and SiC powder mixtures. However, further increasing of the BN phase inhibited the grain growth, because of the poor sinterbility of h-BN which has strong covalent nature and plate-like structure. Although ZrB2 and SiC phases in the samples (except ZSN30) come from two ways: directly from the raw ceramic powders and by solid reaction formation, it is very difficult to distinguish them in microstructures, because no phase and morphology difference was observed. Both ZrB2 and SiC grains in all composites are equiaxed. The added ZrB2 and SiC powders has a much smaller grain size (0.7 μm for ZrB2 and 30 nm for SiC) than the final grain sizes ( 2 μm for ZrB2 and 1 μm for SiC), indicating the significant grain growth for both the added phases and the in-situ formed phases. The micro-sized h-BN grains that homogeneously located at the grain boundaries of ZrB2 and SiC grains had a flakelike morphology, with a length of 2–3 μm, and thickness o1 μm. The anisotropic grain growth of the h-BN plates were found in the composites that contain BN content lower than 20 vol%. As the amount of the BN phase increased, the size and the aspect ratio of the BN grains decreased, because of the distortion of the BN flakes to accommodate the grain boundary space in the composites. The buckled and delaminated h-BN flakes that located in the vicinity of the ZrB2 and SiC interfaces were observed on the fracture surfaces of ZSBN20 and ZSBN30 that fabricated by RSPS, as can be seen in Fig.4. It is thought that the residual thermal stress in the sample have played an important role in creating such phenomena. ZrB2 has a mean thermal expansion coefficient of 5.2 10−6 1C−1 while SiC is 3.3 10−6 1C−1. Watts et al. [31] estimated that a compressive stress of 880 MPa in the SiC phase, and 450 MPa tensile in the ZrB2 were generated in the ZrB2–SiC composites, during the cooling process, due to the thermal expansion mismatch of the two phases. The thermal expansion coefficient of h-BN is highly anisotropic, it shows a high value of ∼40 10−6 1C−1 and low value of ≤10−7 1C−1 in the directions parallel and perpendicular to the c-axis, respectively [23,32]. The thermal mismatch between h-BN, ZrB2, and SiC made the intergranular h-BN particle subject to a tensile stress and a compressive stress in the directions parallel and perpendicular to the c-axis when the sample was cooled down to room temperature after sintering. The tensile stress led to the delamination of h-BN, while the compressive stress caused the bucking of the h-BN layers. Fig. 5 shows a representative TEM micrograph for the reactive spark plasma sintered ZrB2–SiC–BN composite processed in this work. Some h-BN particles (as arrowed in Fig. 5) were entrapped into ZrB2 grains. The intragranular h-BN particles showed a dimension of less than 200 nm, which is much less than that of
44
W.-W. Wu et al. / Materials Science & Engineering A 582 (2013) 41–46
Fig. 3. SEM micrographs of polished surfaces of the ZrB2-SiC-BN composites that were fabricated by RSPS: (a) ZS, (b) ZSBN5, (c) ZSBN10, (d) ZSBN20, (e) ZSBN30 and (f) BN content dependence of the grain size of ZrB2 and SiC.
Fig. 4. Fracture surface of ZSBN20 that fabricated by reactive spark plasma sintering.
Fig. 5. TEM image of ZSBN30 that fabricated by reactive spark plasma sintering. The inserts are the electron diffraction of the particles of circle A and B.
3.3. Mechanical properties the intergranular h-BN flakes that were located at the grain boundaries. This result indicates that the originally in-situ formed BN phase was nano-sized, significant grain growth occurred during the sintering at a relatively high temperature of 1900 1C.
Fig. 6(a) is the hardness and Young's modulus of the ZrB2–SiC–BN ceramics as a function of the content of BN. As expected from the rule of mixtures for the mixture of ZrB2, SiC and h-BN, both of the values decreased with the increasing of BN content, because BN is a soft and
W.-W. Wu et al. / Materials Science & Engineering A 582 (2013) 41–46
45
Fig. 6. Young's modulus and hardness (a) and the flexural strength and fracture toughness and (b) of the reactive spark plasma sintered ZrB2–SiC–BN composites as a function of the content of BN.
Griffith equation 1 K IC pffiffiffi sf ¼ Y c
Fig. 7. SEM micrographs of the polished surface of ZSBN5. The microcracks formed by the residual stress are pointed by black arrows.
weak phase owning to its graphite-like crystal structure, and Young's modulus of BN (65 GPa [18]) is much lower than ZrB2 and SiC (489 GPa [33] and 427 GPa [17], respectively). Fig. 6(b) shows the variation in flexural strength and fracture toughness for the ZrB2–SiC–BN ceramics as a function of the content of BN. The fracture toughness increased from 5.8 70.3 MPa m1/2 for the ZS composite to a maximum value of 6.6 70.2 MPa m1/2 for the ZSBN5 composite, then decreased to 4.3 70.1 MPa m1/2 with the increasing of the BN content to 30 vol %. There are two reasons for the increasing of the fracture toughness. First, the existence of the softness of h-BN and the relatively large ZrB2 grain size at the small h-BN amount led to the crack of the composite propagating preferentially along the weak boundary particles or their interfaces, resulting crack bridging, deflection, or a tortuous crack path. Second, the thermal mismatches of the phases and the grain growth of ZrB2 produced high local residual stresses providing crack-particle interaction and the potential for microcracking (as pointed by black arrows in Fig. 7). The flexural strength decreased as the BN content increased, similar to the other BN containing microcomposites in which the BN grain size is in micron scale. The weak h-BN at the interface of ZrB2 and SiC changed the fracture mode from intragranular into intergranular, and decreased the strength of the composites. Nevertheless, the measured flexural strength from ZSN20 and ZSBN30 (518.8 723.4 MPa and 457.7 726.4 MPa, respectively) are comparable with that of the ZrB2–SiC and higher that the ZrB2–SiC–BN composites which were sintered by the conventional sintering methods [33–35]. The fine and homogeneous microstructure is believed to have an important role in the improvement of the mechanical properties. The flexural strength measured from ZSBN5 was relatively low compared to that of ZSBN10, which can be explained by the
where sf , K IC , c, and Y are flexural strength, fracture toughness, one-half of the width of the initial flaw, and the geometrical parameter of a flaw, respectively. Based on this fracture mechanics, the strength of the brittle ceramics is controlled much stronger by the critical flaw size than the fracture toughness. In a result, although the fracture toughness of ZSN5 was higher, the large flaw size associated with the large ZrB2 grains in ZSBN5 decreased the strength. The increasing of the BN content decreased the grain size of ZrB2, in addition, the 10 vol% of BN did not reach the threshold for the formation of BN a percolation network. As a result, the strength of ZSN10 increased slightly. Further increasing of the BN content to above 10 vol% degraded the material strength, because the BN percolation network can facilitate defect formation at the grain boundaries [36]. Several reports have demonstrated that the nanocomposites which have nano-sized h-BN coated on the matrix particles show much higher flexural strength in comparison with the conventionally fabricated microcomposites [10,18,19,21,23]. The reactive sintering has shown its advantage at refining the microstructure of the composites by low temperature sintering which inhibited grain growth of ZrB2-based ceramics [37,38]. It is expected that the flexural strength of the ZrB2–SiC–BN composites fabricated in this study can be further enhanced by careful processing control in the future. 3.4. Machinability Fig. 8 shows the indent induced by a load of 10 kg during hardness test on polished surfaces of ZSBN30. The absence of classical radial cracks around the indentation site indicates a remarkable contact damage resistance of the composite [39]. The formation of weak interfaces and the cleavage behavior of h-BN particles absorbed the fracture energy during the hardness test and prevented the crack propagation. Fig. 9 shows the photograph of a ZSBN20 plate, on which holes were drilled with a hard steel drill. The ZrB2–SiC–BN composites with BN content higher than 20 vol% were highly machinable. The intergranular h-BN is thought to have elevated the machinability and deformability due to their easy deformation and cleavage characteristics [11]. 3.5. Thermal conductivity The measured heat capacities, thermal diffusivity and the calculated thermal conductivities of the ZrB2–SiC–BN composites fabricated by RSPS are listed in Table 2. The thermal conductivity of the composites decreased with the increase of BN, because of the increased grain boundaries and the random orientation of the
46
W.-W. Wu et al. / Materials Science & Engineering A 582 (2013) 41–46
increased and then decreased with the increasing of the BN content. The flexural strength and the thermal conductivities of the composites decreased with the increasing of BN, while the fracture toughness first increased and then decreased. The composite that contained 420 wt% BN exhibited excellent machinability. Although most of the in-situ formed h-BN particle have a relative large grain size (2–3 μm high flexural strength of 518.8 7 23.4 MPa for the composite with 20 wt% BN was retained.
Acknowledgments
Fig. 8. SEM micrographs of the indent that introduced on the polished surface of ZSBN30 with a diamond indenter at the load of 10 kg.
Financial supports from the National Natural Science Foundation of China (No. 51272266) and the bilateral project of NSFC-JSPS (No. 51111140017) are gratefully acknowledged. References
Fig. 9. Machinable ZrB2–SiC–BN composite containing 20 vol% BN fabricated by RSPS. The diameter of the disk is 25 mm.
Table 2 Heat capacity, thermal diffusivity and thermal conductivity of the ZrB2–SiC–BN composites fabricated by RSPS.
ZSBN30 ZSBN20 ZSBN10 ZSBN5 ZS
Heat capacity Cp (J (g K)−1)
Thermal diffusivity α (mm2/s)
Thermal conductivity kc (W (m K)−1)
0.52 0.53 0.53 0.53 0.56
31.09 32.11 31.18 27.34 27.58
69.3 76.8 79.2 76.9 86.6
low conductive h-BN (in a range of 30–60 W (m K)−1) [40–42]. Nevertheless, the values received in this work is comparable to that of the reported value of ZrB2–SiC composites (60–100 W (m K)−1) [43], which can be attributed to the clean grain boundary that was formed during the reactive sintering. 4. Conclusions Highly dense ZrB2–SiC–BN composites with fine and homogeneous microstructure were successfully fabricated from a mixture of ZrH2, Si3N4 and B4C powders, via reactive spark plasma sintering at 1900 1C in vacuum. The in-situ formed BN phase was composed of micro-sized intergranular and nano-sized intragranular h-BN particles. The particle size of the ZrB2 grains first
[1] K. Upadhya, J.M. Yang, W.P. Hoffman, Am. Ceram. Soc. Bull. 76 (1997) 51–56. [2] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, J. Mater. Sci. 39 (2004) 5887–5904. [3] E. Wuchina, E. Opila, M. Opeka, W. Fahrenholtz, I. Talmy, Electrochem. Soc. Interface 16 (2007) 30–36. [4] A. Paul, D.D. Jayaseelan, S. Venugopal, E. Zapata-Solvas, J. Binner, B. Vaidhyanathan, A. Heaton, P. Brown, W.E. Lee, Am. Ceram. Soc. Bull. 91 (2012) 22–28. [5] S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, J. Am. Ceram. Soc. 91 (2008) 26–32. [6] W.G. Fahrenholtz, G.E. Hilmas, S.C. Zhang, S. Zhu, J. Am. Ceram. Soc. 91 (2008) 1398–1404. [7] W.M. Guo, G.J. Zhang, Z.G. Yang, J. Am. Ceram. Soc. 95 (2012) 2470–2473. [8] G. Li, W.B. Han, B.L. Wang, Mater. Des. 32 (2011) 401–405. [9] C.C. Wei, X.H. Zhang, P. Hu, W.B. Han, G.S. Tian, Scr. Mater. 65 (2011) 791–794. [10] T. Kusunose, T. Sekino, Y.H. Choa, K. Niihara, J. Am. Ceram. Soc. 85 (2002) 2689–2695. [11] Y.L. Li, G.J. Qiao, Z.H. Jin, Mater. Res. Bull. 37 (2002) 1401–1409. [12] X.D. Wang, G.J. Qiao, Z.H. Jin, J. Am. Ceram. Soc. 87 (2004) 565–570. [13] K.S. Mazdiyasni, R. Ruh, J. Am. Ceram. Soc. 64 (1981) 415–419. [14] R. Ruh, A. Zangvil, R.R. Wills, Adv. Ceram. Mater. 3 (1988) 411–415. [15] E.H. Lutz, M.V. Swain, J. Am. Ceram. Soc. 75 (1992) 67–70. [16] G.J. Zhang, M. Ando, T. Ohji, S. Kanzaki, Int. J. Appl. Ceram. Technol. 2 (2005) 162–171. [17] G.J. Zhang, Y. Beppu, T. Ohji, S. Kanzaki, Acta Mater. 49 (2001) 77–82. [18] T. Kusunose, T. Sekino, Y.H. Choa, K. Niihara, J. Am. Ceram. Soc. 85 (2002) 2678–2688. [19] T. Kusunose, T. Sekino, Y. Ando, Nanotechnology 19 (2008) 275603. [20] T. Kusunose, N. Sakayanagi, T. Sekino, Y. Ando, J. Nanosci. Nanotechnol. 8 (2008) 5846–5853. [21] T. Kusunose, Y.H. Kim, T. Sekino, T. Matsumoto, N. Tanaka, T. Nakayama, K. Niihara, J. Mater. Res. 20 (2005) 183–190. [22] R. Shuba, I.W. Chen, J. Am. Ceram. Soc. 89 (2006) 2147–2153. [23] L. Gao, X.H. Jin, J.G. Li, Y.G. Li, J. Sun, Mater. Sci. Eng. A 415 (2006) 145–148. [24] S. Grasso, Y. Sakka, G. Maizza, Sci. Technol. Adv. Mater. 10 (2009) 053001. [25] R. Orru, R. Licheri, A.M. Locci, A. Cincotti, G.C. Cao, Mater. Sci. Eng. R 63 (2009) 127–287. [26] W.W. Wu, G.J. Zhang, Y.M. Kan, P.L. Wang, K. Vanmeensel, J. Vleugels, O. Van der Biest, Scr. Mater. 57 (2007) 317–320. [27] G.J. Zhang, Z.Y. Deng, N. Kondo, J.F. Yang, T. Ohji, J. Am. Ceram. Soc. 83 (2000) 2330–2332. [28] G.J. Zhang, M. Ando, J.F. Yang, T. Ohji, S. Kanzaki, J. Eur. Ceram. Soc. 24 (2004) 171–178. [29] L.V. Strashinskaya, A.N. Stepanchuk, Fiz.-Khim. Mekh. Mater. 6 (1970) 76–79. [30] N.L. Peterson, Wadd Technical Report, 1960 pp. 60–793. [31] J. Watts, G. Hilmas, W.G. Fahrenholtz, D. Brown, B. Clausen, J. Eur. Ceram. Soc. 31 (2011) 1811–1820. [32] W. Sinclair, H. Simmons, J. Mater. Sci. Lett. 6 (1987) 627–629. [33] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, J. Am. Ceram. Soc. 90 (2007) 1347–1364. [34] S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, J. Eur. Ceram. Soc. 31 (2011) 893–901. [35] H.T. Wu, W.G. Zhang, J. Eur. Ceram. Soc. 30 (2010) 1035–1042. [36] G.J. Zhang, T. Ohji, J. Mater. Res. 15 (2000) 1876–1880. [37] W.W. Wu, G.J. Zhang, Y.M. Kan, P.L. Wang, J. Am. Ceram. Soc. 91 (2008) 2501–2508. [38] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, J. Am. Ceram. Soc. 89 (2006) 3638–3645. [39] X.T. Wang, N.P. Padture, H. Tanaka, Nat. Mater. 3 (2004) 539–544. [40] M. Hubacek, M. Ueki, J. Am. Ceram. Soc. 82 (1999) 156–160. [41] D.W. Ni, G.J. Zhang, Y.M. Kan, Y. Sakka, J. Am. Ceram. Soc. 94 (2011) 1397–1404. [42] A. Lipp, K.A. Schwetz, K. Hunold, J. Eur. Ceram. Soc. 5 (1989) 3–9. [43] J.W. Zimmermann, G.E. Hilmas, W.G. Fahrenholtz, R.B. Dinwiddie, W.D. Porter, H. Wang, J. Am. Ceram. Soc. 91 (2008) 1405–1411.