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Spark plasma sintering and high-temperature strength of B6O–TaB2 ceramics
Journal of the European Ceramic Society 37 (2017) 3009–3014
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Short communication
Spark plasma sintering and high-temperature strength of B6 O–TaB2 ceramics Dmytro Demirskyi a,∗ , Oleg Vasylkiv b a b
Nanyang Technological University, 50 Nanyang Avenue, 639798 Singapore, Singapore National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan
a r t i c l e
i n f o
Article history: Received 19 May 2016 Received in revised form 14 February 2017 Accepted 25 February 2017 Available online 6 March 2017 Keywords: High-temperature strength Tantalum diboride Boron suboxide Elevated temperature fracture toughness
a b s t r a c t Tantalum diboride – boron suboxide ceramic composites were densified by spark plasma sintering at 1900 ◦ C. Strength and fracture toughness of these bulk composites at room temperature were 490 MPa and 4 MPa m1/2 , respectively. Flexural strength of B6 O–TaB2 ceramics increased up to 800 ◦ C and remained unchanged up to 1600 ◦ C. At 1800 ◦ C a rapid decrease in strength down to 300 MPa was observed and was accompanied by change in fracture mechanisms suggestive of decomposition of boron suboxide grains. Fracture toughness of B6 O–TaB2 composites showed a minimum at 800 ◦ C, suggestive a relaxation of thermal stresses generated from the mismatch in coefficients of thermal expansion. Flexural strength at elevated temperatures for bulk TaB2 reference sample was also investigated. Results suggest that formation of composite provides additional strengthening/toughening as in all cases flexural strength and fracture toughness of the B6 O–TaB2 ceramic composite was higher than that reported for B6 O monoliths. © 2017 Elsevier Ltd. All rights reserved.
1. Introduction Monolithic boron suboxide (B6 O) possesses high hardness (>28 GPa), high strength (400 MPa), and good high-temperature properties making it as attractive material for high-temperature applications and cutting tools [1–10]. Unlike another commonly used hard covalent boron-rich ceramic such as boron carbide (B4 C), stoichiometric boron suboxide powders are relatively hard to synthesize [6]. Therefore, majority of methods for synthesis of this compound included high-pressure technique, where high pressure (>5 GPa) and relatively high temperatures (∼1600 ◦ C) were used for synthesis and consolidation of B6 O. Recent advances in synthesis showed that highly stoichiometric B6 Ox (where x > 0.85) powders can be mass-synthesized at ambient pressure [11], opening a door to exploration of B6 O-based ceramic composites. The B6 O-based composites are solving the problem monolithic boron suboxide densification, as the oxygen deficiency in the rhombohedral cell at the 6c position, (x in B6 Ox ) affects the ‘sinterability’ of as-synthesized boron suboxide powders [4,5]. Anyhow, covalent B6 O-based ceramics still require application of external pressure (<90 MPa) and high temperatures (>1800 ◦ C) to
∗ Corresponding author. E-mail addresses: [email protected] (D. Demirskyi), [email protected] (O. Vasylkiv). http://dx.doi.org/10.1016/j.jeurceramsoc.2017.02.052 0955-2219/© 2017 Elsevier Ltd. All rights reserved.
activate the consolidation process. Hence techniques such as hotpressing (HP) and spark plasma sintering (SPS) are commonly used. Among B6 O-based ceramics composites the B6 O–B4 C [12,13], B6 O–TiB2 [14,15] and B6 O–ZrB2 [16] are among first to be consolidated using powder metallurgy approach using HP [14] or SPS [13,15,16]. For B6 O–B4 C and B6 O–TiB2 composites reactive and nonreactive approaches where investigated. In case of B6 O–B4 C composites application of reactive SPS consolidation resulted in increase of hardness and fracture toughness of composites [13]. While in the case of B6 O–TiB2 system reactive approach resulted in increase of final density in case of HP [14], while did not affect substantially hardness or fracture hardness of composites. In [15] SPS set-up where graphite dies with BN protection was used (hence current was not directly passing though the specimen). The reactive approach resulted in higher hardness and no improvement in fracture toughness. Keeping in mind the high strength of B4 C–TiB2 ceramics [17–19] and the fact that coefficients of thermal expansions (CTE) of B6 O and B4 C are quite close [15,20], good flexural strength of B6 O–TiB2 ceramics is anticipated. Other ceramics composites in B6 O–MeIV−V B2 system are also within reach. Therefore, this study will be focused on the consolidation of the large size specimens (i.e.∼60 mm) using non-reactive powder mixture of B6 O powder with a low oxygen deficiency level (i.e. high x
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Fig. 1. SEM structures of initial powders and microstructures (polished and fractured at room temperature) of B6 O–TaB2 and TaB2 ceramics consolidated at 1900 ◦ C.
Fig. 2. Configuration used in the present study in order to evaluate fracture toughness of the composites using the three-point flexure.
in B6 Ox ) and commercially available TaB2 powders, with SPS as the consolidation method. To decrease level of carbon diffusion into B6 O or TaB2 during high-temperature consolidation protective Ta foil was used. In addition, the reference monolithic TaB2 was also consolidated by SPS. The strength and toughness of bulk B6 O–TaB2 ceramics at high temperatures of up to 1800 ◦ C is reported for the first time.
2. Experimental Commercially available tantalum diboride powder (TaB2 , dav = 1.0–1.5 m, <0.5 wt.% C, <0.5 wt.% N, <0.7 wt.% O, Lot # T301160, Japan New Metals Co. Ltd., Japan) and mass-synthesized B6 O powder [11] were used as the starting material (Fig. 1). Analysis of previous work on composites of TaB2 with B4 C [21,22] or B6 O with TiB2 [15] indicate that addition of 18–2 vol.% of diboride will yield best ratio between final density and mechanical properties. Hence, in the present study powder mixtures of B6 O and 52 wt.% (18 vol.%) TaB2 were prepared by wet mixing in alcohol followed by drying at about 100 ◦ C. A reference sample of 100% TaB2 diboride was homogenized using wet mixing in alcohol using same procedure as B6 O–TaB2 mixture. The resultant powders were screened through a 60-mesh screen. The homogenized powder mixtures were loaded into a graphite die with an inner diameter of 60 mm and subjected to SPS. The outer surface of the die was wrapped in 5-mm-thick graphite felt to homogenize the temperature distribution and reduce heat loss by
Fig. 3. XRD pattern of B6 O–TaB2 ceramic after SPS at 1900 ◦ C for 1 min.
radiation. To prevent the diffusion of carbon into the powder mixture, tantalum foil (Sigma-Aldrich Chemie, 0.025 mm thick, 99.9+% metal basis) was inserted between the powder and the graphite
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Fig. 4. Flexural strengths in B6 O–TaB2 system at room and elevated temperatures. Left image shows typical loading diagram of the B6 O–TaB2 and TaB2 ceramics tested at different temperatures by the three-point flexural testing. Right image shows temperature dependence of flexural strength of B6 O–TaB2 ceramics. Data presented with upper open triangle are for B6 O–TaB2 composites with low oxygen occupancy starting boron suboxide powder (B6 O0.76 ). The dashed lines in the case of B4 C and B6 O ceramics indicate the general tendencies observed in previous studies [4,27–29]. Closed symbols indicate that strength was measured using 4P set-up, semi-closed and open symbols are for 3P flexural strength tests.
Fig. 5. Microstructures of fractured surfaces of boron suboxide – tantalum diboride ceramics after three-point flexural strength test at different testing temperatures.
Fig. 6. SEM structures of initial powders and microstructures of fractured surfaces of B6 O–TaB2 ceramic composite prepared using B6 O powder with different level of oxygen site occupancy (i.e., B6 O0.76 ), after three-point flexural strength test at different testing temperatures: (b) room temperature, (c) 800 ◦ C and (d,e) 1600 ◦ C. White arrow in (c) indicates the place of specific pore morphology; while the (e) shows magnified area dashed in (d) for the composite fractured at 1600 ◦ C.
foil. The mold system containing the powder mixture was placed in an SPS furnace (HP D125, FCT, Germany). Initially, a pressure of 20 MPa was applied to ensure sufficient electrical contact between the powder tablet and the graphite die, which was then increased to 80 MPa at 800 ◦ C and the temperature was increased to 1900 ◦ C.
A dwell time of 1 min at 800 ◦ C was used to increase the pressure. Then we increased the temperature at a rate of 110 ◦ C min−1 up to a sintering temperature of 1900 ◦ C with a dwell time of 1 min. Each specimen was gradually cooled to 600 ◦ C at a rate of 100 ◦ C min−1 and then naturally to room temperature in the furnace. In case of
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the monolithic TaB2 , 10 min dwell time at 1900 ◦ C was used. The sintering process was performed in argon gas with a flow rate of 2 L min−1 . The sintered specimens were ground with diamond disks with a particle size of up to 0.5 m. Then, the density of the samples was measured by the Archimedes method using ethanol as a medium in accordance with ASTM B 963–08. The three-point flexural strength was determined using rectangular blocks (2 × 2.5 × 20 mm) cut from specimens with a diameter of 60 mm using electric discharge machining. Their lateral surfaces were ground and polished using diamond pastes. Three-point flexural strength tests were conducted at room temperature and at high temperatures of up to 1800 ◦ C in argon using a Shimadzu AG-X plus system (Shimadzu, Japan). The loading speed was 0.5 mm min−1 . Ten to twelve samples were tested at each temperature, and the measurement accuracy was taken as the standard deviation. For the high-temperature flexural strength tests, the following heating schedule was used: from room temperature to 200 ◦ C in 10 min and from 200 ◦ C to the testing temperature at a rate of 18 ◦ C min−1 . A dwell time of 5 min was employed before the flexural test at the testing temperature. After testing, cooling from the testing temperature to room temperature was performed at a rate of 20 ◦ C min−1 . The fracture toughness of composites was evaluated using specimens testing in bending which contained a single edge through-thickness notch in accordance with ASTM C1421–10. A general view of the used configuration and a notch profile are presented in Fig. 2. Five tests were conducted at room temperature and four specimens were tested at each of the same elevated temperatures used for flexure testing. To verify the obtained fracture toughness values, we conducted reference flexural strength tests before the respective fracture toughness tests at elevated temperatures. Loading rate of 0.5 mm/min was used. Heating and cooling schedules were identical to those used for flexural tests. Tests at elevated temperatures were performed in argon gas. Microstructural observations and analyses were carried out on the polished samples using SU 8000 cold-emission FE-SEM (Hitachi, Japan) and TM3000 (Hitachi, Japan) in back-scattered mode. X-ray diffraction (XRD) analysis (Rigaku RINT 2500 HLR, Japan) was performed on polished samples to identify the crystalline phases.
3. Results and discussion Fig. 1 shows microstructures of B6 O–TaB2 and TaB2 ceramics consolidated at 1900 ◦ C. Image of B6 O–TaB2 ceramics polished surface indicates rather homogeneous mixing, although some zones of B6 O with size of 40–150 m were present regardless of mixing procedure. Grain size of both phases in dense B6 O–TaB2 composites was less than 5 m, only few 10–40 m aggregates of boron suboxide were noticed during the fractographic analysis. In the case of monolithic TaB2 specimens porosity of up to 6% was determined, most of pores had clear spherical shape and were located at the grain boundaries and inside the grains. The mean grain size of 15–20 m was estimated on polished specimens. The grain size and pore morphology suggest that rapid grain growth [23] occurred during high-temperature SPS at 1900 ◦ C. Furthermore, earlier study on TaC ceramics consolidated at 2000 ◦ C indicated similar density and grain size [24]. However, unlike TaC in case of TaB2 monoliths a minor amount of sharp-edged pores was observed. Mainly intergranular fracture was observed for B6 O–TaB2 and TaB2 ceramics. Some large grains were fractured in the transgranullar manner; surface marking and cleavage were also observed. XRD analysis (Fig. 3) showed that B6 O–TaB2 ceramics consisted only from boron suboxide (#50-1505) and tantalum diboride (#65-
0878) phases. The lattice parameters of B6 O phase were determined as a = 0.5388(2) nm, and c = 1.2348(3) nm (in hexagonal expression), this yielded oxygen occupancy x ∼ 0.85–0.88. Monolithic tantalum diboride ceramics (not shown) were indexed using (#650878) card, without any additional oxide or carbide phases. In the case of the B6 O–TaB2 ceramics no trace of B2 O3 or other impurities was observed. This situation is believed to be caused by (i) rapid heating schedule during SPS and (ii) the use of Ta foil which limits to a certain degree amount of graphite diffused into the powder mixture during the spark plasma sintering process [9] Fig. 4 [24–29] shows loading curves for ceramics in the B6 O–TaB2 system recorded during the three-point flexural strength tests at different temperatures. No change in fracture behavior between room temperature and 800 ◦ C test was observed. For the tests performed at 1600 ◦ C a change in testing curve slope was observed indicating change in the elastic modulus during testing. A slight decrease in strength was observed for monolithic TaB2 ceramics from 392 ± 13 to 385 ± 9 MPa, at 800 ◦ C and 1600 ◦ C, respectively. In case of the B6 O–TaB2 composites (i.e., B6 O0.85 –TaB2 ceramics) the strength was within 521 ± 13 MPa for both temperatures. Increase in the flexural test temperature up to 1800 ◦ C resulted in decrease in strength (Figure 4). However, in case of the B6 O–TaB2 composites more profound plastic behavior of the loading curve was observed. In the case of TaB2 reference specimens, and elastic behavior was observed at 1800 ◦ C, however significant drop in the slope of the curve also suggests decrease in elastic modulus. For TaB2 , even clearer strength decrease was observed in [25]: from 240 ± 94 MPa at room temperature to 39 ± 7 MPa at 1200 ◦ C. This was explained by the penetration of the residual oxygen in the furnace, although the flexural strength tests were performed in argon. The results of the present study also indicate minor oxidation of the TaB2 during 1800 ◦ C strength tests (Fig. 5). However, it is unclear how this process affected the observed strength (136 ± 13 MPa), however we should note that in the case of ZrB2 ceramics a strength of 220 ± 18 MPa was observed at 1800 ◦ C [30]. Lack of visual confirmation on fracture in [25,30] makes it impossible to make final conclusion on the effect of residual oxygen on the fracture behavior and thus strength of monolithic TaB2 . Hence, to verify this situation, further investigation of high-temperature strength and fracture behavior of other diborides would be necessary. In the case of B6 O–TaB2 composites a more complex situation was observed. Previous systematic study on strength of B6 O ceramics [4] suggested that flexural strength will increase with temperature up to 1400 ◦ C. Works on the strength of B4 C ceramics also confirm this trend [27–29]. Therefore, a slight increase in strength up to 1600 ◦ C may indicate that the flexural strength in B6 O–TaB2 composites is controlled by B6 O grains. In order to fully understand the effect of boron suboxide on strength behavior in B6 O–TaB2 composites, we synthesized boron suboxide powder with different value of oxygen occupancy. Synthesis procedure was performed according to [13], with synthesis time and temperature of 150 min and 1320 ◦ C. For thus synthesized B6 O powder, lattice parameters a = 0.5366(6) nm and c = 1.2329(8) nm (in hexagonal expression) were refined on the basis of the X-ray powder diffraction. Calculations using the refined lattice parameters from the powder XRD show that the estimated value of oxygen occupancy, x, was found to range between 0.76 and 0.81. This is consistent with analysis presented in [31]. Thus synthesized powder consisted of the irregular-shape plate-like particles (Fig. 6(a)), although stark-like and five-fold symetherical features were noticed at higher magnificantions. The strength of these ceramics was unchanged between room temperature and 1600 ◦ C, and the mean value of the strength was 443 ± 19 MPa. This is considerably lower compared with 474 ± 17 MPa obtained for composite consolidated by SPS using
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Fig. 7. Fracture toughness in B6 O–TaB2 system at elevated temperatures obtained using a loading rate of 0.5 mm/min (see data in Table 1). (a) shows typical loading curves recoded during fracture toughness tests of B6 O0.85 –TaB2 ceramics. (b) shows a temperature dependence of the KIC parameter for the B6 O–TaB2 composites, and includes data on fracture toughness at room temperature for TaB2 [26] and B6 O [3,35] monoliths. Table 1 Mechanical performance of B6 O − TaB2 composites at elevated temperatures. Material
Testing conditions
, MPa
ia , MPa
Fracture toughness, KIC , MPa m1/2 Mean
Actual values
B6 O0.85 –TaB2
25 ◦ C 800 ◦ C 1600 ◦ C 1800 ◦ C
474 ± 17 521 ± 14 523 ± 13 302 ± 22
451 512 530 290
3.81 3.33 3.83 3.50
(4.01 (3.12 (3.83 (3.47
3.97 3.52 4.08 3.36
3.88 3.41 3.69 3.50
3.52 3.27) 3.74) 3.69)
3.71)
B6 O0.76 –TaB2
25 ◦ C 800 ◦ C 1600 ◦ C
444 ± 12 440 ± 16 443 ± 19
446 452 439
3.66 3.22 3.79
(3.52 (3.21 (3.95
3.82 3.19 3.81
3.71 3.17 3.79
3.68 3.31) 3.63)
3.58)
a
i shows the value of the reference strength test performed before the fracture toughness tests at using identical loading conditions (0.5 mm/min).
B6 O0.85 powder with higher level of oxygen occupancy. Specimens of two B6 O–TaB2 composites with different starting boron suboxide powder had total porosity less than 2% and identical level of distribution of TaB2 in B6 O matrix; hence the difference in strength may be solely attributed to the quality of initial boron suboxide powder. Results on the B6 O bulks suggested that strength of B6 O0.85 was slightly lower than that of B6 O0.88 (i.e. 300 vs 314 MPa) [4]; other factors such as porosity or grain size may also contribute significantly. In terms of fractographic data both composites fractured in similar manner (compare Figs. 5 and 6), and the only distinctive feature of B6 O–TaB2 ceramics produced using B6 O0.76 powder was the presence of closed pores with a specific shape, which were previously reported for B4 C ceramics (see white arrow in Fig. 6(c)). Such pores were also reported in [32] for the pressureless sintering of boron carbide. These closed pores although rarely observed in boron-rich ceramics may influence strength. Some aspects related to pores shape influence on flexural strength are discussed for TiB2 –TaC composites in [24]. At 1600 ◦ C intergranular fracture was observed in the case of B6 O–TaB2 composites, surfaces were free of striations which were clearly observed at lower testing temperatures (Fig. 5). Increased number of grain pull-out sites was also observed. Only at high magnifications one can define grooving of grain surface by thermal etching and surface diffusion (Fig. 6(e)). At 1800 ◦ C a different surface structure of TaB2 and B6 O grains was observed. TaB2 grains were mostly round-shape and exhibited spiral marking that are result of thermal etching by gas phase or
were induced by surface diffusion. B6 O grains had an onion like shape, where the layers with same crystallographic orientation were stuck one on each other. This is visually similar to fine cleavage steps sometimes observed during the fracture of ceramics [24]. However, the prismatic shape of the observed B6 O layers is similar to that observed during the synthesis of boron-rich covalent solids in B-C-N-O system [33]. We suggest that in the present case, some dissociation of B6 O took place [34]. Hence, further decrease in strength of B6 O–TaB2 ceramics is anticipated if temperature of flexural strength test is increased. The results of the fracture toughness tests are shown in Fig. 7 and are summarized in Table 1. Fig. 7 shows that at ambient temperature a slight difference in toughness between the two B6 O–TaB2 composites is observed. Furthermore, for the monolithic boron suboxide ceramics toughness between 2 and 3 MPa m1/2 was previously reported [3,35]. Hence, one can consider that toughening by the particulate reinforcement allows increase in fracture toughness up to 4 MPa m1/2 , which is close to the toughness of respective diboride (4.5 MPa m1/2 ) [26]. Improvements in the toughness of two-phase particulate composites have been shown to be a result of the residual stresses generated from the CTE mismatch between phases [17]. Upon heating, these stresses relax and the toughness decreases. Hence, a decrease in toughness at 800 ◦ C is observed in both composites. At 1600 ◦ C toughness increases to the near value at room temperature, followed by a slight decrease to 3.5 MPa m1/2 for the specimens tested at 1800 ◦ C. The increase in toughness at 1600 ◦ C can be attributed to stress relief ahead of the crack tip by the plastic
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flow as suggested in [36]. Slight increase in toughness from 4.7 to 5.0 MPa m1/2 was observed for monolithic TiB2 , when temperature of the tests was increased to 960 ◦ C [37]. Similar dependence may be expected to be valid for other diborides. As summarized in [38], in general, for coarse grained monolithic diborides or carbides the change in strength or toughness due to activation of plastic deformation is comparatively small for temperatures below 1800 ◦ C. Furthermore, the increased contribution of plasticity to fracture process has been observed during flexural tests of ZrB2 , TiB2 , B4 C or their composites at temperatures below 2000 ◦ C by numerous researches [24,27–29,36–39]. In the framework of the present investigation, loading curves with plastic behavior during flexure were observed for B6 O–TaB2 composite at 1800 ◦ C. Similar situation was observed during fracture toughness tests at 1800 ◦ C (see Fig. 4(a)). One may except that further increase in temperature will lead to decrease in fracture toughness and flexural strength due to dissociation of the B6 O matrix. Furthermore, we suggest that data on fracture toughness of TaB2 , B6 O or B4 C monoliths may assist in deeper understanding of the toughening or strengthening mechanisms active in the B6 O–TaB2 system at elevated temperatures. 4. Conclusions Fully dense B6 O–TaB2 ceramics were produced by non-reactive consolidation using spark plasma sintering at 1900 ◦ C and 1 min dwell time. Use of tantalum foil during SPS ensured that composite was free from impurities or other minor phases. B6 O–TaB2 composites showed high room-temperature strength up to 490 MPa and toughness of 3.8 MPa m1/2 . Reference specimen of monolithic TaB2 with density of 94% exhibited strength of 400 ± 20 MPa from room temperature up to 1600 ◦ C. Temperature dependence of flexural strength and fracture toughness of B6 O–TaB2 composites suggested sufficient contribution of plasticity between 1600 and 1800 ◦ C. Finally, microstructure evolution during the hightemperature flexural strength tests indicated mostly intergranular fracture of composites. Acknowledgement We thank to Dr. Toshiyuki Nishimura (NIMS) for the use of the flexural strength measurement facility. References [1] M. Herrmann, H.-J. Kleebe, J. Raethel, K. Sempf, S. Lauterbach, M.M. Müller, I. Sigalas, Field-assisted densification of superhard B6 O materials with Y2 O3 /Al2 O3 addition, J. Am. Ceram. Soc. 92 (2009) 2368–2372. [2] M. Thiele, M. Herrmann, A. Michaelis, B6 O materials with Al2 O3 /Y2 O3 additives densified by FAST/SPS and HIP, J. Eur. Ceram. Soc. 33 (2013) 2375–2390. [3] M. Herrmann, A.K. Swarnakar, M. Thiele, O. Vander Biest, I. Sigalas, High temperature properties of B6 O-Materials, J. Eur. Ceram. Soc. 31 (2011) 2387–2392. [4] D. Demirskyi, I. Solodkyi, Y. Sakka, O. Vasylkiv, High-Temperature strength of boron suboxide ceramic consolidated by spark plasma sintering, J. Am. Ceram. Soc. 99 (2016) 2669–2777. [5] I. Solodkyi, D. Demirskyi, Y. Sakka, O. Vasylkiv, Hardness and toughness control of brittle boron suboxide ceramics by consolidation of star-shaped particles by spark plasma sintering, Ceram. Int. 42 (2016) 3525–3530. [6] O.O. Kurakevych, Superhard phases of simple substances and binary compounds of the B–C–N–O system: from diamond to the latest results (a review), J. Superhard Mater. 31 (2009) 139–157. [7] D.V. Turkevych, V. Bushlya, J.-E. Stahl, I.A. Petrusha, N.N. Belyavina, V.Z. Turkevich, HP-HT sintering, microstructure, and properties of B6O- and TiC-containing composites based on cBN, J. Superhard Mater. 37 (2015) 143–154. [8] H. Itoh, I. Maekawa, H. Iwahara, High pressure sintering of B6O powder and properties of the sintered compact, J. Soc. Mater. Sci. Jpn. 47 (1998) 1000–1005.
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McMillan, High-Pressure high-Temperature synthesis and characterization of boron suboxide (B6O), Chem. Mater. 10 (1998) 1530–1537. [32] N. Cho, Z. Bao, R.F. Speyer, Density- and herdness-optimized pressureless sintered and post-hot pressed B4C, J. Mater. Res. 20 (2005) 2110–2116. [33] H. Hubert, L.A.J. Garvie, P.R. Buseck, W.T. Petuskey, P.F. McMillan, High-Pressure, high-temperature syntheses in the B-C-N-O system, J. Solid State Chem. 133 (1997) 356–364. [34] H.F. Rizzo, W.C. Simmons, H.O. Bielstein, The existence and formation of the Solid B6 O, J. Electrochem. Soc. 109 (1962) 1079–1082. [35] D. Demirskyi, O. Vasylkiv, Microstructure and mechanical properties of boron suboxide ceramics prepared by pressureless microwave sintering, Ceram. Int. 42 (2016) 14282–14286. [36] G.E. Neuman, Ultra-high temperature mechanical properties of a zirconium diboride–zirconium carbide ceramic, J. Am. Ceram. Soc. 99 (2016) 597–603. [37] H.R. Baumgartner, R.A. 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Short communication
Spark plasma sintering and high-temperature strength of B6 O–TaB2 ceramics Dmytro Demirskyi a,∗ , Oleg Vasylkiv b a b
Nanyang Technological University, 50 Nanyang Avenue, 639798 Singapore, Singapore National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan
a r t i c l e
i n f o
Article history: Received 19 May 2016 Received in revised form 14 February 2017 Accepted 25 February 2017 Available online 6 March 2017 Keywords: High-temperature strength Tantalum diboride Boron suboxide Elevated temperature fracture toughness
a b s t r a c t Tantalum diboride – boron suboxide ceramic composites were densified by spark plasma sintering at 1900 ◦ C. Strength and fracture toughness of these bulk composites at room temperature were 490 MPa and 4 MPa m1/2 , respectively. Flexural strength of B6 O–TaB2 ceramics increased up to 800 ◦ C and remained unchanged up to 1600 ◦ C. At 1800 ◦ C a rapid decrease in strength down to 300 MPa was observed and was accompanied by change in fracture mechanisms suggestive of decomposition of boron suboxide grains. Fracture toughness of B6 O–TaB2 composites showed a minimum at 800 ◦ C, suggestive a relaxation of thermal stresses generated from the mismatch in coefficients of thermal expansion. Flexural strength at elevated temperatures for bulk TaB2 reference sample was also investigated. Results suggest that formation of composite provides additional strengthening/toughening as in all cases flexural strength and fracture toughness of the B6 O–TaB2 ceramic composite was higher than that reported for B6 O monoliths. © 2017 Elsevier Ltd. All rights reserved.
1. Introduction Monolithic boron suboxide (B6 O) possesses high hardness (>28 GPa), high strength (400 MPa), and good high-temperature properties making it as attractive material for high-temperature applications and cutting tools [1–10]. Unlike another commonly used hard covalent boron-rich ceramic such as boron carbide (B4 C), stoichiometric boron suboxide powders are relatively hard to synthesize [6]. Therefore, majority of methods for synthesis of this compound included high-pressure technique, where high pressure (>5 GPa) and relatively high temperatures (∼1600 ◦ C) were used for synthesis and consolidation of B6 O. Recent advances in synthesis showed that highly stoichiometric B6 Ox (where x > 0.85) powders can be mass-synthesized at ambient pressure [11], opening a door to exploration of B6 O-based ceramic composites. The B6 O-based composites are solving the problem monolithic boron suboxide densification, as the oxygen deficiency in the rhombohedral cell at the 6c position, (x in B6 Ox ) affects the ‘sinterability’ of as-synthesized boron suboxide powders [4,5]. Anyhow, covalent B6 O-based ceramics still require application of external pressure (<90 MPa) and high temperatures (>1800 ◦ C) to
∗ Corresponding author. E-mail addresses: [email protected] (D. Demirskyi), [email protected] (O. Vasylkiv). http://dx.doi.org/10.1016/j.jeurceramsoc.2017.02.052 0955-2219/© 2017 Elsevier Ltd. All rights reserved.
activate the consolidation process. Hence techniques such as hotpressing (HP) and spark plasma sintering (SPS) are commonly used. Among B6 O-based ceramics composites the B6 O–B4 C [12,13], B6 O–TiB2 [14,15] and B6 O–ZrB2 [16] are among first to be consolidated using powder metallurgy approach using HP [14] or SPS [13,15,16]. For B6 O–B4 C and B6 O–TiB2 composites reactive and nonreactive approaches where investigated. In case of B6 O–B4 C composites application of reactive SPS consolidation resulted in increase of hardness and fracture toughness of composites [13]. While in the case of B6 O–TiB2 system reactive approach resulted in increase of final density in case of HP [14], while did not affect substantially hardness or fracture hardness of composites. In [15] SPS set-up where graphite dies with BN protection was used (hence current was not directly passing though the specimen). The reactive approach resulted in higher hardness and no improvement in fracture toughness. Keeping in mind the high strength of B4 C–TiB2 ceramics [17–19] and the fact that coefficients of thermal expansions (CTE) of B6 O and B4 C are quite close [15,20], good flexural strength of B6 O–TiB2 ceramics is anticipated. Other ceramics composites in B6 O–MeIV−V B2 system are also within reach. Therefore, this study will be focused on the consolidation of the large size specimens (i.e.∼60 mm) using non-reactive powder mixture of B6 O powder with a low oxygen deficiency level (i.e. high x
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Fig. 1. SEM structures of initial powders and microstructures (polished and fractured at room temperature) of B6 O–TaB2 and TaB2 ceramics consolidated at 1900 ◦ C.
Fig. 2. Configuration used in the present study in order to evaluate fracture toughness of the composites using the three-point flexure.
in B6 Ox ) and commercially available TaB2 powders, with SPS as the consolidation method. To decrease level of carbon diffusion into B6 O or TaB2 during high-temperature consolidation protective Ta foil was used. In addition, the reference monolithic TaB2 was also consolidated by SPS. The strength and toughness of bulk B6 O–TaB2 ceramics at high temperatures of up to 1800 ◦ C is reported for the first time.
2. Experimental Commercially available tantalum diboride powder (TaB2 , dav = 1.0–1.5 m, <0.5 wt.% C, <0.5 wt.% N, <0.7 wt.% O, Lot # T301160, Japan New Metals Co. Ltd., Japan) and mass-synthesized B6 O powder [11] were used as the starting material (Fig. 1). Analysis of previous work on composites of TaB2 with B4 C [21,22] or B6 O with TiB2 [15] indicate that addition of 18–2 vol.% of diboride will yield best ratio between final density and mechanical properties. Hence, in the present study powder mixtures of B6 O and 52 wt.% (18 vol.%) TaB2 were prepared by wet mixing in alcohol followed by drying at about 100 ◦ C. A reference sample of 100% TaB2 diboride was homogenized using wet mixing in alcohol using same procedure as B6 O–TaB2 mixture. The resultant powders were screened through a 60-mesh screen. The homogenized powder mixtures were loaded into a graphite die with an inner diameter of 60 mm and subjected to SPS. The outer surface of the die was wrapped in 5-mm-thick graphite felt to homogenize the temperature distribution and reduce heat loss by
Fig. 3. XRD pattern of B6 O–TaB2 ceramic after SPS at 1900 ◦ C for 1 min.
radiation. To prevent the diffusion of carbon into the powder mixture, tantalum foil (Sigma-Aldrich Chemie, 0.025 mm thick, 99.9+% metal basis) was inserted between the powder and the graphite
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Fig. 4. Flexural strengths in B6 O–TaB2 system at room and elevated temperatures. Left image shows typical loading diagram of the B6 O–TaB2 and TaB2 ceramics tested at different temperatures by the three-point flexural testing. Right image shows temperature dependence of flexural strength of B6 O–TaB2 ceramics. Data presented with upper open triangle are for B6 O–TaB2 composites with low oxygen occupancy starting boron suboxide powder (B6 O0.76 ). The dashed lines in the case of B4 C and B6 O ceramics indicate the general tendencies observed in previous studies [4,27–29]. Closed symbols indicate that strength was measured using 4P set-up, semi-closed and open symbols are for 3P flexural strength tests.
Fig. 5. Microstructures of fractured surfaces of boron suboxide – tantalum diboride ceramics after three-point flexural strength test at different testing temperatures.
Fig. 6. SEM structures of initial powders and microstructures of fractured surfaces of B6 O–TaB2 ceramic composite prepared using B6 O powder with different level of oxygen site occupancy (i.e., B6 O0.76 ), after three-point flexural strength test at different testing temperatures: (b) room temperature, (c) 800 ◦ C and (d,e) 1600 ◦ C. White arrow in (c) indicates the place of specific pore morphology; while the (e) shows magnified area dashed in (d) for the composite fractured at 1600 ◦ C.
foil. The mold system containing the powder mixture was placed in an SPS furnace (HP D125, FCT, Germany). Initially, a pressure of 20 MPa was applied to ensure sufficient electrical contact between the powder tablet and the graphite die, which was then increased to 80 MPa at 800 ◦ C and the temperature was increased to 1900 ◦ C.
A dwell time of 1 min at 800 ◦ C was used to increase the pressure. Then we increased the temperature at a rate of 110 ◦ C min−1 up to a sintering temperature of 1900 ◦ C with a dwell time of 1 min. Each specimen was gradually cooled to 600 ◦ C at a rate of 100 ◦ C min−1 and then naturally to room temperature in the furnace. In case of
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the monolithic TaB2 , 10 min dwell time at 1900 ◦ C was used. The sintering process was performed in argon gas with a flow rate of 2 L min−1 . The sintered specimens were ground with diamond disks with a particle size of up to 0.5 m. Then, the density of the samples was measured by the Archimedes method using ethanol as a medium in accordance with ASTM B 963–08. The three-point flexural strength was determined using rectangular blocks (2 × 2.5 × 20 mm) cut from specimens with a diameter of 60 mm using electric discharge machining. Their lateral surfaces were ground and polished using diamond pastes. Three-point flexural strength tests were conducted at room temperature and at high temperatures of up to 1800 ◦ C in argon using a Shimadzu AG-X plus system (Shimadzu, Japan). The loading speed was 0.5 mm min−1 . Ten to twelve samples were tested at each temperature, and the measurement accuracy was taken as the standard deviation. For the high-temperature flexural strength tests, the following heating schedule was used: from room temperature to 200 ◦ C in 10 min and from 200 ◦ C to the testing temperature at a rate of 18 ◦ C min−1 . A dwell time of 5 min was employed before the flexural test at the testing temperature. After testing, cooling from the testing temperature to room temperature was performed at a rate of 20 ◦ C min−1 . The fracture toughness of composites was evaluated using specimens testing in bending which contained a single edge through-thickness notch in accordance with ASTM C1421–10. A general view of the used configuration and a notch profile are presented in Fig. 2. Five tests were conducted at room temperature and four specimens were tested at each of the same elevated temperatures used for flexure testing. To verify the obtained fracture toughness values, we conducted reference flexural strength tests before the respective fracture toughness tests at elevated temperatures. Loading rate of 0.5 mm/min was used. Heating and cooling schedules were identical to those used for flexural tests. Tests at elevated temperatures were performed in argon gas. Microstructural observations and analyses were carried out on the polished samples using SU 8000 cold-emission FE-SEM (Hitachi, Japan) and TM3000 (Hitachi, Japan) in back-scattered mode. X-ray diffraction (XRD) analysis (Rigaku RINT 2500 HLR, Japan) was performed on polished samples to identify the crystalline phases.
3. Results and discussion Fig. 1 shows microstructures of B6 O–TaB2 and TaB2 ceramics consolidated at 1900 ◦ C. Image of B6 O–TaB2 ceramics polished surface indicates rather homogeneous mixing, although some zones of B6 O with size of 40–150 m were present regardless of mixing procedure. Grain size of both phases in dense B6 O–TaB2 composites was less than 5 m, only few 10–40 m aggregates of boron suboxide were noticed during the fractographic analysis. In the case of monolithic TaB2 specimens porosity of up to 6% was determined, most of pores had clear spherical shape and were located at the grain boundaries and inside the grains. The mean grain size of 15–20 m was estimated on polished specimens. The grain size and pore morphology suggest that rapid grain growth [23] occurred during high-temperature SPS at 1900 ◦ C. Furthermore, earlier study on TaC ceramics consolidated at 2000 ◦ C indicated similar density and grain size [24]. However, unlike TaC in case of TaB2 monoliths a minor amount of sharp-edged pores was observed. Mainly intergranular fracture was observed for B6 O–TaB2 and TaB2 ceramics. Some large grains were fractured in the transgranullar manner; surface marking and cleavage were also observed. XRD analysis (Fig. 3) showed that B6 O–TaB2 ceramics consisted only from boron suboxide (#50-1505) and tantalum diboride (#65-
0878) phases. The lattice parameters of B6 O phase were determined as a = 0.5388(2) nm, and c = 1.2348(3) nm (in hexagonal expression), this yielded oxygen occupancy x ∼ 0.85–0.88. Monolithic tantalum diboride ceramics (not shown) were indexed using (#650878) card, without any additional oxide or carbide phases. In the case of the B6 O–TaB2 ceramics no trace of B2 O3 or other impurities was observed. This situation is believed to be caused by (i) rapid heating schedule during SPS and (ii) the use of Ta foil which limits to a certain degree amount of graphite diffused into the powder mixture during the spark plasma sintering process [9] Fig. 4 [24–29] shows loading curves for ceramics in the B6 O–TaB2 system recorded during the three-point flexural strength tests at different temperatures. No change in fracture behavior between room temperature and 800 ◦ C test was observed. For the tests performed at 1600 ◦ C a change in testing curve slope was observed indicating change in the elastic modulus during testing. A slight decrease in strength was observed for monolithic TaB2 ceramics from 392 ± 13 to 385 ± 9 MPa, at 800 ◦ C and 1600 ◦ C, respectively. In case of the B6 O–TaB2 composites (i.e., B6 O0.85 –TaB2 ceramics) the strength was within 521 ± 13 MPa for both temperatures. Increase in the flexural test temperature up to 1800 ◦ C resulted in decrease in strength (Figure 4). However, in case of the B6 O–TaB2 composites more profound plastic behavior of the loading curve was observed. In the case of TaB2 reference specimens, and elastic behavior was observed at 1800 ◦ C, however significant drop in the slope of the curve also suggests decrease in elastic modulus. For TaB2 , even clearer strength decrease was observed in [25]: from 240 ± 94 MPa at room temperature to 39 ± 7 MPa at 1200 ◦ C. This was explained by the penetration of the residual oxygen in the furnace, although the flexural strength tests were performed in argon. The results of the present study also indicate minor oxidation of the TaB2 during 1800 ◦ C strength tests (Fig. 5). However, it is unclear how this process affected the observed strength (136 ± 13 MPa), however we should note that in the case of ZrB2 ceramics a strength of 220 ± 18 MPa was observed at 1800 ◦ C [30]. Lack of visual confirmation on fracture in [25,30] makes it impossible to make final conclusion on the effect of residual oxygen on the fracture behavior and thus strength of monolithic TaB2 . Hence, to verify this situation, further investigation of high-temperature strength and fracture behavior of other diborides would be necessary. In the case of B6 O–TaB2 composites a more complex situation was observed. Previous systematic study on strength of B6 O ceramics [4] suggested that flexural strength will increase with temperature up to 1400 ◦ C. Works on the strength of B4 C ceramics also confirm this trend [27–29]. Therefore, a slight increase in strength up to 1600 ◦ C may indicate that the flexural strength in B6 O–TaB2 composites is controlled by B6 O grains. In order to fully understand the effect of boron suboxide on strength behavior in B6 O–TaB2 composites, we synthesized boron suboxide powder with different value of oxygen occupancy. Synthesis procedure was performed according to [13], with synthesis time and temperature of 150 min and 1320 ◦ C. For thus synthesized B6 O powder, lattice parameters a = 0.5366(6) nm and c = 1.2329(8) nm (in hexagonal expression) were refined on the basis of the X-ray powder diffraction. Calculations using the refined lattice parameters from the powder XRD show that the estimated value of oxygen occupancy, x, was found to range between 0.76 and 0.81. This is consistent with analysis presented in [31]. Thus synthesized powder consisted of the irregular-shape plate-like particles (Fig. 6(a)), although stark-like and five-fold symetherical features were noticed at higher magnificantions. The strength of these ceramics was unchanged between room temperature and 1600 ◦ C, and the mean value of the strength was 443 ± 19 MPa. This is considerably lower compared with 474 ± 17 MPa obtained for composite consolidated by SPS using
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Fig. 7. Fracture toughness in B6 O–TaB2 system at elevated temperatures obtained using a loading rate of 0.5 mm/min (see data in Table 1). (a) shows typical loading curves recoded during fracture toughness tests of B6 O0.85 –TaB2 ceramics. (b) shows a temperature dependence of the KIC parameter for the B6 O–TaB2 composites, and includes data on fracture toughness at room temperature for TaB2 [26] and B6 O [3,35] monoliths. Table 1 Mechanical performance of B6 O − TaB2 composites at elevated temperatures. Material
Testing conditions
, MPa
ia , MPa
Fracture toughness, KIC , MPa m1/2 Mean
Actual values
B6 O0.85 –TaB2
25 ◦ C 800 ◦ C 1600 ◦ C 1800 ◦ C
474 ± 17 521 ± 14 523 ± 13 302 ± 22
451 512 530 290
3.81 3.33 3.83 3.50
(4.01 (3.12 (3.83 (3.47
3.97 3.52 4.08 3.36
3.88 3.41 3.69 3.50
3.52 3.27) 3.74) 3.69)
3.71)
B6 O0.76 –TaB2
25 ◦ C 800 ◦ C 1600 ◦ C
444 ± 12 440 ± 16 443 ± 19
446 452 439
3.66 3.22 3.79
(3.52 (3.21 (3.95
3.82 3.19 3.81
3.71 3.17 3.79
3.68 3.31) 3.63)
3.58)
a
i shows the value of the reference strength test performed before the fracture toughness tests at using identical loading conditions (0.5 mm/min).
B6 O0.85 powder with higher level of oxygen occupancy. Specimens of two B6 O–TaB2 composites with different starting boron suboxide powder had total porosity less than 2% and identical level of distribution of TaB2 in B6 O matrix; hence the difference in strength may be solely attributed to the quality of initial boron suboxide powder. Results on the B6 O bulks suggested that strength of B6 O0.85 was slightly lower than that of B6 O0.88 (i.e. 300 vs 314 MPa) [4]; other factors such as porosity or grain size may also contribute significantly. In terms of fractographic data both composites fractured in similar manner (compare Figs. 5 and 6), and the only distinctive feature of B6 O–TaB2 ceramics produced using B6 O0.76 powder was the presence of closed pores with a specific shape, which were previously reported for B4 C ceramics (see white arrow in Fig. 6(c)). Such pores were also reported in [32] for the pressureless sintering of boron carbide. These closed pores although rarely observed in boron-rich ceramics may influence strength. Some aspects related to pores shape influence on flexural strength are discussed for TiB2 –TaC composites in [24]. At 1600 ◦ C intergranular fracture was observed in the case of B6 O–TaB2 composites, surfaces were free of striations which were clearly observed at lower testing temperatures (Fig. 5). Increased number of grain pull-out sites was also observed. Only at high magnifications one can define grooving of grain surface by thermal etching and surface diffusion (Fig. 6(e)). At 1800 ◦ C a different surface structure of TaB2 and B6 O grains was observed. TaB2 grains were mostly round-shape and exhibited spiral marking that are result of thermal etching by gas phase or
were induced by surface diffusion. B6 O grains had an onion like shape, where the layers with same crystallographic orientation were stuck one on each other. This is visually similar to fine cleavage steps sometimes observed during the fracture of ceramics [24]. However, the prismatic shape of the observed B6 O layers is similar to that observed during the synthesis of boron-rich covalent solids in B-C-N-O system [33]. We suggest that in the present case, some dissociation of B6 O took place [34]. Hence, further decrease in strength of B6 O–TaB2 ceramics is anticipated if temperature of flexural strength test is increased. The results of the fracture toughness tests are shown in Fig. 7 and are summarized in Table 1. Fig. 7 shows that at ambient temperature a slight difference in toughness between the two B6 O–TaB2 composites is observed. Furthermore, for the monolithic boron suboxide ceramics toughness between 2 and 3 MPa m1/2 was previously reported [3,35]. Hence, one can consider that toughening by the particulate reinforcement allows increase in fracture toughness up to 4 MPa m1/2 , which is close to the toughness of respective diboride (4.5 MPa m1/2 ) [26]. Improvements in the toughness of two-phase particulate composites have been shown to be a result of the residual stresses generated from the CTE mismatch between phases [17]. Upon heating, these stresses relax and the toughness decreases. Hence, a decrease in toughness at 800 ◦ C is observed in both composites. At 1600 ◦ C toughness increases to the near value at room temperature, followed by a slight decrease to 3.5 MPa m1/2 for the specimens tested at 1800 ◦ C. The increase in toughness at 1600 ◦ C can be attributed to stress relief ahead of the crack tip by the plastic
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flow as suggested in [36]. Slight increase in toughness from 4.7 to 5.0 MPa m1/2 was observed for monolithic TiB2 , when temperature of the tests was increased to 960 ◦ C [37]. Similar dependence may be expected to be valid for other diborides. As summarized in [38], in general, for coarse grained monolithic diborides or carbides the change in strength or toughness due to activation of plastic deformation is comparatively small for temperatures below 1800 ◦ C. Furthermore, the increased contribution of plasticity to fracture process has been observed during flexural tests of ZrB2 , TiB2 , B4 C or their composites at temperatures below 2000 ◦ C by numerous researches [24,27–29,36–39]. In the framework of the present investigation, loading curves with plastic behavior during flexure were observed for B6 O–TaB2 composite at 1800 ◦ C. Similar situation was observed during fracture toughness tests at 1800 ◦ C (see Fig. 4(a)). One may except that further increase in temperature will lead to decrease in fracture toughness and flexural strength due to dissociation of the B6 O matrix. Furthermore, we suggest that data on fracture toughness of TaB2 , B6 O or B4 C monoliths may assist in deeper understanding of the toughening or strengthening mechanisms active in the B6 O–TaB2 system at elevated temperatures. 4. Conclusions Fully dense B6 O–TaB2 ceramics were produced by non-reactive consolidation using spark plasma sintering at 1900 ◦ C and 1 min dwell time. Use of tantalum foil during SPS ensured that composite was free from impurities or other minor phases. B6 O–TaB2 composites showed high room-temperature strength up to 490 MPa and toughness of 3.8 MPa m1/2 . Reference specimen of monolithic TaB2 with density of 94% exhibited strength of 400 ± 20 MPa from room temperature up to 1600 ◦ C. Temperature dependence of flexural strength and fracture toughness of B6 O–TaB2 composites suggested sufficient contribution of plasticity between 1600 and 1800 ◦ C. Finally, microstructure evolution during the hightemperature flexural strength tests indicated mostly intergranular fracture of composites. Acknowledgement We thank to Dr. Toshiyuki Nishimura (NIMS) for the use of the flexural strength measurement facility. References [1] M. Herrmann, H.-J. Kleebe, J. Raethel, K. Sempf, S. Lauterbach, M.M. Müller, I. Sigalas, Field-assisted densification of superhard B6 O materials with Y2 O3 /Al2 O3 addition, J. Am. Ceram. Soc. 92 (2009) 2368–2372. [2] M. Thiele, M. Herrmann, A. Michaelis, B6 O materials with Al2 O3 /Y2 O3 additives densified by FAST/SPS and HIP, J. Eur. Ceram. Soc. 33 (2013) 2375–2390. [3] M. Herrmann, A.K. Swarnakar, M. Thiele, O. Vander Biest, I. Sigalas, High temperature properties of B6 O-Materials, J. Eur. Ceram. Soc. 31 (2011) 2387–2392. [4] D. Demirskyi, I. Solodkyi, Y. Sakka, O. Vasylkiv, High-Temperature strength of boron suboxide ceramic consolidated by spark plasma sintering, J. Am. Ceram. Soc. 99 (2016) 2669–2777. [5] I. Solodkyi, D. Demirskyi, Y. Sakka, O. Vasylkiv, Hardness and toughness control of brittle boron suboxide ceramics by consolidation of star-shaped particles by spark plasma sintering, Ceram. Int. 42 (2016) 3525–3530. [6] O.O. Kurakevych, Superhard phases of simple substances and binary compounds of the B–C–N–O system: from diamond to the latest results (a review), J. Superhard Mater. 31 (2009) 139–157. [7] D.V. Turkevych, V. Bushlya, J.-E. Stahl, I.A. Petrusha, N.N. Belyavina, V.Z. Turkevich, HP-HT sintering, microstructure, and properties of B6O- and TiC-containing composites based on cBN, J. Superhard Mater. 37 (2015) 143–154. [8] H. Itoh, I. Maekawa, H. Iwahara, High pressure sintering of B6O powder and properties of the sintered compact, J. Soc. Mater. Sci. Jpn. 47 (1998) 1000–1005.
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