cBN reinforced Y-α-SiAlON composites

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Journal of the European Ceramic Society 33 (2013) 2191–2198

cBN reinforced Y-␣-SiAlON composites J.C. Garrett a,∗ , I. Sigalas a , M. Herrmann b , E.J. Olivier c , J.H. O’Connell c a

DST/NRF Centre of Excellence in Strong Materials, School of Chemical and Metallurgical Engineering, University of the Witwatersrand, Private Bag 3, Wits, Johannesburg 2050, South Africa b Fraunhofer-Institute of Ceramic Technologies and Systems (IKTS), Dresden, Germany c Centre for High Resolution Transmission Electron Microscopy, Nelson Mandela Metropolitan University, Port Elizabeth, South Africa Received 14 December 2012; received in revised form 7 March 2013; accepted 9 March 2013 Available online 6 April 2013

Abstract Dense ␣-Sialon–cBN composites were produced by FAST/SPS–sintering at 1575–1625 ◦ C. The hardness of the materials increases only up to 21 GPa for materials with 10 vol.% cBN. On the other hand the fracture toughness increases up to nearly 8 MPa m0.5 with 30 vol.% cBN. The reason for the increase in fracture toughness is attributed to crack deflection at cBN grains due to the weak bonding of the grains in the matrix. The weak interfaces are also responsible for the moderate increase in hardness. Detailed investigation of the interface between cBN and the matrix was carried out by TEM. © 2013 Elsevier Ltd. All rights reserved. Keywords: Sialon; Boron nitride (cBN); Spark plasma sintering (SPS); Hardness; Toughness and toughening

1. Introduction Silicon nitride ceramics have been widely investigated due to their advantageous properties of high strength, fracture toughness, thermal shock, chemical resistance and hardness at both room and elevated operating temperatures.1–4 ␣-Sialon ceramics have improved hardness compared to ␤-silicon nitride ceramics, making them superior as cutting tool materials. The advantageous properties of Sialon ceramics have led to the advancement of Sialon as a cutting tool material selected for cast-iron and difficult to machine super alloys. Improvements in the wear resistance of ceramic cutting tools can be achieved through the addition of hard particulates with good adhesion to the ceramic matrix. One possible candidate for reinforcing is cubic boron nitride (cBN) which is the second hardest material known to and used commercially by man after diamond.5 It is known that cBN has a tendency to transform into

∗

Corresponding author at: 406A Richard Ward Building, East Campus, University of the Witwatersrand, Johannesburg 2050, South Africa. Tel.: +27 72 724 1639. E-mail addresses: [email protected], [email protected] (J.C. Garrett). 0955-2219/$ – see front matter © 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.jeurceramsoc.2013.03.014

its more stable hexagonal phase when sintered under relatively low pressures and temperatures above 1300 ◦ C.6 Therefore the use of spark plasma sintering (SPS) as a consolidation technique was selected due to the high heating and cooling rates achievable, which in turn reduces the exposure time of the cBN grains to high sintering temperatures, suppressing the transformation from the cubic to the hexagonal phase.7 Research into the fabrication of cBN containing Sialon composites has mainly been carried out with ␤-Sialon as the matrix material, where it was found that these ␤-Sialon/cBN composites exhibited maximum hardness, fracture toughness and flexural strength values of 15.4 GPa (Hv5 ), 6.8 MPa m0.5 and 432 MPa respectively with a 10 vol.% cBN addition.8 Additionally, the reinforcing cBN grains were observed to transform into the hexagonal structure at an approximate sintering temperature of 1650 ◦ C which consequently decreased the hardness of the composites. This effect became more pronounced with increased sintering temperature. Furthermore it was observed that increased heating rates could retard this transformation.9 These data also revealed that the period of time spent above the cBN to hBN transition temperature is the decisive factor to the degree of transformation experienced. However, the use of SiO2 -containing phases (SiO2 , SiAlON and mullite) as the parent matrix results in a significant reduction in the hBN

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transformation, in turn making cBN stable at higher temperatures as compared to other cBN containing composites (Al2 O3 , Al2 O3 -Ni, TiN and WC-Co).10 Zhao and Wang11 specified that a 50 vol.% cBN containing ␤Sialon composite with a relative density of 98.6 ± 0.9% could be obtained through SPS (Dr Sinter) at 1450 ◦ C; additionally it was claimed that a HV0.5 of 48.0 ± 0.9 GPa and fracture toughness of 11.5 MPa m0.5 were obtained from a load of 4.9 N. The exceptionally high micro hardness values obtained by Zhao et al., indicate that the measurements were obtained solely on the cBN grains which had an average size of 7.5 ␮m. That value is close to the Vickers hardness of pure cBN. Currently no data exists concerning the interaction and behaviour of cBN in ␣-Sialon matrix compositions. This study investigates the reinforcement of Y-␣-Sialon ceramics through the addition of ≈10 ␮m cBN grains as the reinforcing agent. 2. Experimental procedure The starting powders used to fabricate the Y-␣-SiAlON matrices were ␣-Si3 N4 (SN-10, UBE), AlN (Grade H, Tokuyama), Al2 O3 (AKP 50, Sumitomo-Chemical) and Y2 O3 (Grade C, H.C. Starck). The varying compositions of the Sialon ceramics were calculated on the basis of the overall formula of the Sialon phase Ym/3 Si12−(m+n) Al(m+n) On N16−n . The two compositions investigated had significantly different values of m and n and were named M1025 (m = 1.0 and n = 2.5) and M2045 (m = 2.0 and n = 4.5), respectively. The oxygen content of the nitride starting powders corresponding to 2.5 wt.% SiO2 and Al2 O3 in the Si3 N4 and AlN powders respectively was accounted for in the calculation of the Sialon compositions. An additional 2 wt.% of Y2 O3 was added in order to aid densification through the formation of a permanent liquid phase. The composition of M2045 was designed for comparison to that presented by Ye et al.8 however the location of the M2045 Sialon composition is slightly above the Sialon plane in the Jänecke prism due to the addition of the 2 wt.% Y2 O3 . The starting powders were mixed using a planetary ball mill for 4 h at a speed of 200 rpm. Agate milling media (bowl and 1 cm balls) was used along with isopropanol as the medium, after which the mixture was dried through the use of a rotary evaporator. The wear of the agate media was found to be negligible. The reinforcing cBN grains (Grade 9, Element 6 Pty. Ltd.) were introduced into the Sialon matrix in 10, 20 and 30 vol.% additions through the use of a dry Turbula mixing route (60 rpm for 2 h). Sintering of discs 20 mm in diameter and approximately 5 mm thickness was carried out with the use of a SPS furnace (HP-D5 FCT, Germany) under vacuum with a constant dwell time of 5 min and uni-axial pressure of 50 MPa. The heating profile consisted of a two-step heating ramp rate, firstly 250 ◦ C/min up to 1350 ◦ C and secondly 50 ◦ C/min up to the desired sintering temperature, cooling however was constant at 200 ◦ C/min. The sintering temperatures investigated were 1550, 1575, 1600 and 1625 ◦ C. The densities of the sintered compacts were measured through the use of the Archimedes method.12 The crystalline

phases were identified through XRD analysis (D2, Bruker) with Cu K␣ radiation (30 kV and 10 mA). Measurements were taken between 10◦ and 60◦ with a step size of 0.02◦ . Microstructural investigations were carried out using a scanning electron microscope (FEI Quanta 400 FEG) accompanied with both secondary and back scattered electron detectors. Transmission electron microscopy investigations were also conducted to obtain a better understanding of the transformation/reaction zone between the Sialon matrix and cBN grains. TEM specimens were prepared by FIB (FEI Helios Nanolab 650) at 30 kV with final polishing at 500 V. The TEM investigation was carried out using a JEOL JEM 2100 microscope. Confirmation of boron nitride crystalline phases was completed through matching of experimental SAED patterns to simulated SAED patterns using JEMS simulation software package.13 The Vickers hardness (Hv5 ) was measured using a Leco V-100-A2, where at least five indentations were used per hardness value. The indentation fracture toughness was determined from crack measurements under the same load calculated through the description given by Anstis et al.,14 where the main requirements to be met are a homogenous microstructure with a parallel mirror polished surface free of pores and cracks.15 A Young’s modulus of 300 GPa was used for the Sialon matrix.16,17 The values determined by indentation fracture toughness have limitations; a detailed evaluation of this method was given by Quinn and Bradt.18 Nevertheless the method is mentioned in prCENT/TS 14425-1 as a possible method for evaluating the fracture toughness of ceramic materials where the only exclusions are very tough or porous materials. This is not the case for our materials. The absolute values determined should be used with caution however trends with changing composition and sintering the values do reflect and can be used for comparative purposes with other ceramic materials and composite measured through particular method. Furthermore the investigations of the crack path verify the indentation fracture toughness results. 3. Results and discussion 3.1. Densification The sintering conditions and achieved densities are given in Table 1. The densification curves of the pure matrices are given in Fig. 1. It is well recognised that the intensive densification of Sialon begins with the formation of the oxide liquid phase at about 1200–1250 ◦ C.2–4,19,20 Both composites start to densify in this temperature range predominantly through the rearrangement of the Si3 N4 particles. However, the onset temperature pertaining to the oxide liquid formation is lower for M2045 as well as the intensity of densification being higher than that of M1025, which is expected since it is more oxygen-rich than M1025 Sialon. However, the initial peak of densification rate occurs at approximately 1380 ◦ C for both materials (Fig. 1). After the first maximum of the densification rate (dL/dt; Fig. 1) a steady increase in densification is observed with increasing sintering temperature, as seen from the continuous piston travel with increasing sintering temperature (travel; Fig. 1). However, the densification rate of M2045 reduces in the temperature

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Table 1 Density, hardness and fracture toughness of the composite materials. For comparison the properties of the material obtained by Ye et al.8 are also given. Matrix type

Sintering temperature (◦ C)

cBN content (vol.%)

Absolute density (g/cm3 )

M1025

1550 1575 1575 1575 1575 1600 1600 1600 1600 1625 1625 1625 1625

0 0 10 20 30 0 10 20 30 0 10 20 30

2.85 3.33 3.32 3.31 3.31 3.33 3.32 3.29 3.29 3.32 3.30 3.28 3.28

± ± ± ± ± ± ± ± ± ± ± ± ±

M2045

1575 1600 1600 1625

0 0 30 0

3.23 3.39 3.34 3.42

± ± ± ±

Ye et al.8

1500a 1500a

0 30

– –

a

Density, % theoretical density

Hardness (HV5 ) (GPa)

Fracture toughness (MPa m0.5 )

0.03 0.01 0.01 0.01 0.01 0.01 0.01 0.01 0.01 0.01 0.01 0.01 0.01

85.6 100.0 99.3 98.5 98.1 100.0 99.1 97.9 97.5 99.7 98.7 97.6 97.2

– 18.2 ± 0.4 20.9 ± 0.5 20.2 ± 0.8 19.7 ± 0.2 19.1 ± 0.3 21.3 ± 0.8 19.5 ± 0.8 16.8 ± 0.9 19.7 ± 0.5 20.7 ± 0.3 16.2 ± 0.4 15.9 ± 0.3

– 3.6 ± 0.3 5.2 ± 0.2 5.2 ± 0.2 5.4 ± 0.3 3.7 ± 0.2 4.6 ± 0.4 5.9 ± 0.4 7.3 ± 0.3 3.7 ± 0.2 4.7 ± 0.2 7.0 ± 0.8 7.8 ± 0.7

0.01 0.01 0.01 0.01

94.4 99.2 97.1 100.0

– 15.5 ± 0.3 16.1 ± 0.4 16.3 ± 0.2

– 4.2 ± 0.1 6.9 ± 0.2 5.2 ± 0.2

14.0 14.5

4.0 6.3

99.6 94.3

Optical pyrometer is focused on the outer surface of the die which will in turn record a value lower to that experienced in the material.

range 1470–1560 ◦ C before it starts to increase again (Fig. 1). Such retardation is less pronounced in the M1025 material. The reduced densification in this temperature range for M2045 is most likely attributed to the intermediate crystallisation of phases which were not investigated in detail. The formation of a nitrogen rich liquid phase is caused by the dissolution of AlN and ␣-Si3 N4 within the oxide liquid which subsequently becomes richer in nitrogen resulting in the precipitation of Sialon polytypes and ␣/␤-Sialon (see also20 ); this is signified as the last peak of the shrinkage rate (Fig. 1). Despite the more intensive sintering of M2045 Sialon at low temperatures, full densification occurs at a lower temperature for M1025 (1575 ◦ C) as compared to M2045 (1625 ◦ C) (Fig. 2). An increase in the sintering temperature from 1500 ◦ C to 1575 ◦ C results in increased density, as expected; however for M1025

Fig. 1. Piston travel and piston travel rate (shrinkage) during densification in the SPS of M1025 and M2045 Sialons sintered at 1600 ◦ C under 50 MPa uni-axial pressure.

Sialon sintered above 1575 ◦ C an apparent decrease is observed. This was attributed to a density change caused by increasing formation of ␣-Sialon since no porosity was observed through SEM analysis (see next paragraph). Fig. 3 shows the dependence of the relative density with respect to cBN content in the Sialon composites. It is evident that the densification is retarded with an increase in cBN content due to the formation of a rigid skeleton structure. However, all composites sintered at and above 1575 ◦ C for M1025 Sialon had a density higher than 97% of the respective theoretical densities. The same is found for M2045, but the sintering temperature required to obtain full densification is higher (1600 ◦ C). The densification is also affected by increasing sintering temperatures due to the transformation from cBN to hBN. The

Fig. 2. Dependence of the relative density of pure M1025 and M2045 materials on the sintering temperature.

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Fig. 3. Dependence of the density of the composite materials on the cBN content for samples sintered at different temperatures.

transformation creates a volume increase which reduces the relative density of the material even if it is 100% dense. The sintering temperatures measured were lower than those used by Hotta and Goto9 who had found that ␤-Sialon/cBN composites could be fully densified at 1700 ◦ C with 10 vol.% cBN without holding. However, the temperature for full densification of M2045 observed is higher than that determined by Ye et al.8 The most probable reason for this difference is due to the assorted SPS machines using various temperature measuring systems. These differences can result in more than 100 ◦ C deviation in the nominal temperature.21 3.2. Microstructure and phase formation Fig. 4 shows an overview of the microstructure of the two Sialon matrices and Fig. 5 summarises the results of the XRD analysis for M1025 and M2045. M1025 consists primarily of ␣-Sialon whereas in M2045 there is, beside ␣-Sialon the 12H polytype as a main phase. The ␣-Sialon can be identified in the SEM micrographs by the more equiaxed grey grains, whereas the polytype and ␤-Sialon appear darker with the polytype phase being present in the form of clusters. The microstructures of the materials sintered at different temperatures are given in Figs. 6 and 7. They show a homogeneous

Fig. 5. Results of the XRD-analysis of the pure Sialon matrix as well as Sialon composites containing 30 vol.% cBN sintered at 1600 ◦ C, for M1025 (a) and M2045 (b).

distribution of the cBN and no porosity of the materials. A detailed analysis of the cBN/matrix interface reveals that an interfacial layer is formed. SEM micrographs of the cross section of M1025 sintered with 10 vol.% cBN at 1625 ◦ C (Fig. 6) reveal a thickness for the reaction zone in the region of 1–2 ␮m on the surface of the cBN grains. The rough surfaces of cBN grains preferentially reacted with the matrix, whereas some of the well facetted surfaces do not show any evidence of this reaction.

Fig. 4. SEM micrographs showing the differing grain morphologies of (a) M1025 Sialon and (b) M2045 Sialons sintered at 1600 ◦ C and analysed under back-scattered detection mode.

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Fig. 6. SEM micrograph showing the inhomogeneous hexagonal transformation of cBN grains within the M1025 Sialon matrix when sintered at 1625 ◦ C.

Furthermore the yttrium-rich glassy phase was also observed to pool between neighbouring cBN grains particularly at higher cBN contents. This is caused by the applied external uniaxial pressure forcing the liquid phase to penetrate between the three dimensional cBN skeleton (Fig. 7). Very little to no porosity was observed from the micrographs (Figs. 6 and 7), revealing that the calculated relative densities in Table 1 seem to be lower than the actual composite densities. This is caused by the incomplete reaction of the starting powders into Sialon and the mentioned phase transformation of the cBN into hexagonal BN. Fig. 8 highlights the effect of hBN content on the relative density of cBN containing composites. The relative density of the M1025/cBN composites sintered at 1625 ◦ C was calculated under the assumption of different degree of conversion of cBN into hBN. The assumption of a 100 nm thick hBN layer result only in a slight change whereas a hBN layer thickness of 1 ␮m results in relative densities of more than 100%. Since the thickness of the hBN layer in the M1025/cBN composites sintered at 1625 ◦ C is nearly 1 ␮m thick (Fig. 6) the material must be nearly dense. This is in agreement with the results of the polished cross sections showing nearly no

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Fig. 8. Dependence of the calculated relative density of the composites with the matrix M1025 sintered at 1625 ◦ C as a function of the assumed thickness of the hBN layer on the surface of the cBN grains.

porosity. This was attributed to the formation of lower density hBN phase. Since a quantitative phase composition study could not be completed, changes in the respective phases with increasing cBN content could contribute additionally to changes in the theoretical densities. The interfacial layer between the cBN grains and the matrix was investigated in more detail using TEM (Figs. 9–11). The transformation of cBN grains to its hexagonal modification was observed, in this work through HRTEM, to occur from temperatures as low as 1575 ◦ C (Fig. 10). This transformation resulted in reaction zones exhibiting different structures: dense epitaxially grown hBN layers on the cBN surface and layers consisting of nano-sized hBN grains surrounded by the amorphous yttria-rich phase between hBN grains (Fig. 11). The growth of the reaction zone between the Sialon matrix and the cBN grains was promoted with an increase in sintering temperature e.g. compare Fig. 9a and b. The formation of hBN takes place by different mechanisms. On the one side direct conversion of cBN into hBN is very likely for the epitaxially grown layers (Fig. 10). On the other side a complex structure of the interfacial layer consisting of nano hBN grains and glassy oxinitride phases suggest that a solution precipitation mechanism could take place. Different transport mechanisms seem to be possible: Formation of the hBN at the interface can take place via the following reaction: Si3 N4 + 2B2 O3 (l) ↔ 4BN + 3SiO2 (l) (G(1600 ◦ C) = −291 kJ/mol)

(1)

Fig. 7. SEM micrograph showing the existence of the yttria rich pooling behaviour between cBN grains with the addition of 30 vol.% cBN to M2045 sintered at 1600 ◦ C.

The reaction is strongly driven into the direction of BN and SiO2 . But nevertheless a low amount of B2 O3 dissolved in the oxynitride liquid can be expected. A calculation using Factsage (Factsage 6.3) results in a solubility of about 2 Mol% B2 O3 in the SiO2 liquid at 1600 ◦ C. This solubility would be altered due to the other components in the melt, but no reliable data exists about the full system. Nevertheless, this small solubility would be enough for the slow dissolution of the metastable cBN in

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Fig. 11. HRTEM micrograph illustrating the contents of the reaction zone between the cBN grains and the Sialon matrix consisting of nano-sized hBN grains surrounded by an amorphous yttria-rich phase.

the liquid and precipitation as the stable hBN. However, this process is slow due to the low solubility. The experimental data suggest that this process is accelerated in materials with higher BN content. The absence of hBN layer on facetted surfaces of the cBN grains (Fig. 6) could be a result of a very low solubility rate or the nucleation is hindered and the dissolved BN precipitate in other areas. Additional transport can take place via B2 O3 and (BO)2 through the gas phase in the pores which exist in the early stages of densification. 3.3. Hardness and fracture toughness

Fig. 9. STEM image of M1025 sintered with 10 vol.% cBN at 1575 ◦ C (a) and 1625 ◦ C (b). The pull-out of the hBN grains is attributed to the poor wetting of hBN with that of the oxide rich glassy phase.

The hardness and fracture toughness values are given in Table 1. Additionally the dependencies of the values on the sintering temperatures and cBN content are given in Figs. 12 and 13. The hardness of the matrix materials increases with increasing sintering temperature due to the increased formation of the ␣Sialon phase (Fig. 14). It is evident from Fig. 14 that a 10 vol.% cBN addition to M1025 Sialon matrix results in considerable increase in the hardness which is nearly in agreement with the values obtained by rule of mixtures of Sialon and cBN (Fig. 13)

Fig. 10. HRTEM micrograph showing the transformation interface. The stacking of the cBN basal planes is much tighter than that of the hBN and more ordered. Inserts show the SAED patterns for hBN (left) and cBN (right), respectively. Fig. 12. Dependence of the hardness and fracture toughness of pure M1025 and M2045 Sialon ceramics on the sintering temperature.

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Fig. 13. Effect of the cBN content on the hardness and fracture toughness of M1025 Sialon ceramics sintered at 1575, 1600 and 1625 ◦ C and theoretical composite hardness calculated using the rule of mixtures (RoM).

Fig. 14. XRD analysis showing phase formation in M1025 as a function of sintering temperature.

for sintering temperature 1575 ◦ C but is less than expected for the materials sintered at higher temperature. With increasing sintering temperature the thickness of the hBN layer around the hBN grains increases. Therefore this is the reason for the

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Fig. 15. TEM bright field image showing the crack path through the reaction zone followed by intergranular crack propagation through the Sialon matrix.

lower increase in hardness. However, it has to be mentioned that thin hBN layers in the sintered at 1575 ◦ C material have only a minor influence for the material with 10 vol.% cBN. Further addition of cBN results in the decrease in the hardness as compared to M1025-10 cBN composites. The deterioration is more pronounced at the higher sintering temperatures. The cross sections of the materials showed that there were nearly no differences in porosity therefore this cannot be the main reason for the reduction of hardness. This reduction of the hardness is caused by the weak bonding of the cBN grains in the matrix and is more pronounced with increasing hexagonal transformation of the cBN grains. With increasing cBN content the weakly bonded cBN grains start to form a three dimensional network which soften the material. The fracture toughness values increase with increasing sintering temperatures and increasing cBN content (Fig. 13). Again due to the weak bonding of cBN with the matrix, as a result of the hBN formation, an interlayer crack deflection mechanism is strongly activated as illustrated in the TEM bright field image (Fig. 15). The increasing thickness of the hBN interlayers with increasing temperature weaken the bonding of the cBN grains and result in strong intergranular fracture through crack deflection and crack bridging (Fig. 16a) which were observed by tracking the movement of propagating cracks resultant from

Fig. 16. Intergranular (a) crack propagation around cBN grains and transgranular (b) crack propagation through cBN grains for M1025-30cBN-1625.

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the indentation measurements. Some transgranular crack propagation through cBN grains was also evident, especially at higher sintering temperatures as it is mainly attributed to the damage of the cBN grains during sintering (Fig. 16b). The microstructural investigations are therefore in agreement with the fracture toughness measurements. The investigations showed that Y-␣-Sialon/cBN composites had a considerably higher hardness (by as much as 5 GPa) than that of cBN composites prepared with ␤-Sialon as the matrix material, reported by Ye et al.8 and Hotta and Goto.22 Although the hardness for cBN containing composites has been improved by selection of ␣-Sialon as the matrix over that of ␤-Sialon, the composite hardness is not as high as that obtained for Al2 O3 /cBN (HV = 26 GPa)6 and Al2 O3 -Ni/cBN (HV = 27 GPa).23 4. Conclusion Although the mechanical properties of Sialon materials were improved with the addition of cBN as a reinforcing agent, significant improvement in hardness was limited to a 10 vol.% cBN addition as well as a sintering temperature of 1575 ◦ C, where the hexagonal transformation of the cBN grains into hBN could be minimised. At these sintering temperatures dense composite materials have been obtained. An increase in sintering temperature promotes the formation of ␣-Sialon which in turn improves the hardness of the Sialon matrix. However, increasing sintering temperature also enhances the phase transformation of boron nitride from the metastable superhard cubic structure to the soft hexagonal one. TEM investigations have shown that both direct cBN–hBN transformation and a solution precipitation mechanism via the liquid phase takes place. This transformation results in increased fracture toughness and reduced hardness of the composites due to the formation of weak Sialon/cBN interfaces. Further improvement of the materials requires better control of the formation of the Sialon matrix, the phase transformation of the cBN as well as the precise control of the interface between the cBN and Sialon matrix. References 1. Hampshire S. Silicon nitride ceramics – review of structure, processing and properties. J Achieve Mater Manufact Eng 2007;24:43–50. 2. Petzow G, Herrmann M. Silicon nitride ceramics. Structure and bonding, vol. 102. Berlin, Heidelberg: Springer-Verlag; 200247–167.

3. Herrmann M, Hohn S, Bales A. Kinetics of rare earth incorporation and its role in densification and microstructure formation of alpha-Sialon. J Eur Ceram Soc 2012;32:1313–9. 4. Kumar A, Acikbas N, Kara F, Mandal H, Basu D. A comparative study of SiAlON ceramics. Ceram Int 2012;38:5757–67. 5. Wentorf R, De Vries R, Bundy F. Sintered superhard materials. Science 1980;208:873–80. 6. Hotta M, Goto T. Densification and microstructure of Al2 O3 –cBN composites prepared by spark plasma sintering. J Ceram Soc Jpn 2008;116:744–8. 7. Yaman B, Mandal H. Spark plasma sintering of Co-WC cubic boron nitride composites. Mater Lett 2009;63:1041–3. 8. Ye F, Hou Z, Zhang H, Liu L, Zhou Y. Spark plasma sintering of cBN/␤SiAlON composites. Mater Sci Eng A 2010;527:4723–6. 9. Hotta M, Goto T. Effect of time on microstructure and hardness of ␤SiAlON–cubic boron nitride composites during spark plasma sintering. Ceram Int 2011;37:521–4. 10. Zhang J, Tu R, Goto T. Densification, microstructure and mechanical properties of SiO2 –cBN composites by spark plasma sintering. Ceram Int 2012;38:351–6. 11. Zhao Y, Wang M. Effect of sintering temperature on the structure and properties of polycrystalline cubic boron nitride prepared by SPS. J Mater Process Technol 2009;209:355–9. 12. American Society for Testing and Materials (ASTM C373-88 2006) obtained from http://www.astm.org/Standards/C373.htm 13. Stadelmann P. EMS – a software package for electron diffraction analysis and HREM image simulation in materials science. Ultramicroscopy 1987;21:131–45. 14. Anstis G, Chantikul P, Lawn B, Marshall D. A critical evaluation of indentation techniques for measuring fracture toughness: I. Direct crack measurements. J Am Ceram Soc 1981;64:533–8. 15. Rocha-Rangel E. Fracture toughness determination by means of indentation fracture. In: Nanocomposites with unique properties and applications in medicine and industry. Intech; 201121–37. 16. Jiang X, Baek Y, Lee S, Kang S. Formation of alpha-Sialon layer on beta-Sialon and its effect on mechanical properties. J Am Ceram Soc 1998;81:1907–12. 17. Liu L, Ye F, Zhou Y, Zhang Z, Hou Q. Fast bonding alpha-Sialon ceramics by spark plasma sintering. J Eur Ceram Soc 2010;30:2683–9. 18. Quinn G, Bradt R. On the vickers indentation fracture toughness test. J Am Ceram Soc 2007;90:673–80. 19. Peng H. Spark plasma sintering of Si3 N4 -based ceramics. PhD dissertation. Stockholm University; 2004. 20. Rosenflanz A, Chen I. Kinetics of phase transformations in SiAlON ceramics. I: Effects of cation size, composition and temperature. J Eur Ceram Soc 1999;19:2325–35. 21. Herrmann M, Räthel J, Schulz I. Spark plasma sintering/field assisted sintering. Int Ceram Rev 2009;58:109–14. 22. Hotta M, Goto T. Densification and phase transformation of betaSialon–cubic born nitride composites prepared by spark plasma sintering. J Am Ceram Soc 2009;92:1684–90. 23. Zhang J, Tu R, Goto T. Spark plasma sintering of Al2 O3 –cBN composites facilitated by Ni nanoparticle precipitation on cBN powder by rotary chemical vapour deposition. J Eur Ceram Soc 2011;31:2083–7.